second israel materials engineering conference

INIS-mf—10066

SECOND ISRAELMATERIALS ENGINEERING

CONFERENCE

February 21-23,1984

BEN-GURION UNIVERSITY OF THE NEGEV

BEER-SHEVA, ISRAEL

CONFERENCE PROCEEDINGS

Edited by

A.GRILL & S.I.ROKKLIN

SECOND ISRAELMATERIALS ENGINEERING

CONFERENCE

February 21-23.1984

BEN-GURION UNIVERSITY OF THE NEGEV

BEER SHEVA, ISRAEL

CONFERENCE PROCEEDINGS

E d i t e d b y

A . G R I L L & S . I . R O K H L I N

ORGANIZING COMMITTEE

Chairman - Prof. S. Rokhlin, Ben-Gurion University of the NegevDr. U. Arnon, Israel Aircraft IndustryProf. D. Brandon, Technion - I s r a e l In s t i t u t e of Technology and

Is rae l Ins t i tu te of MetalsDr. B. Cina, Israel Metallurgical Society and Is rae l Aircraft IndustryProf. M. Dariel , Ben-Gurion University of the Negev and Nuclear Research

Centre - NegevDr. S. Kenig, Rafael - HaifaDr. G. Metzger, National Council for Research DevelopmentDr. H. Paruz, Ministry of DefenseProf. J . Pel leg, Ben-Gurion University of the NegevDr. M. Polak, Ben-Gurion University of the NegevProf. M. Ron, Technion - I s rae l Ins t i tu te of TechnologyDr. A. Stern, Nuclear Research Centre - NegevProf. B.Z. Weiss, Technion - Is rae l Ins t i tu te of Techology

EDITORIAL COMMITTEE

Chairman - Prof. A. G r i l l , Ben-Gurion Universi ty of the NegevProf. L. Kornbli t , Ben-Gurion University of the NegevProf. I . Minkoff, Technion - I s r a e l I n s t i t u t e of TechnologyMr. B. Rnbin, I s rae l Aircraf t IndustryProf. M. Rosen, Ben-Gurion University of the NegevProf. M. Schieber, The Hebrew University of JerusalemDr. J . Zahavi, I s r ae l I n s t i t u t e of MetalsMr. Z. Wagner, I s r ae l Mi l i t a ry Indust r ies

THE CONFERENCE WAS SPONSORED BY:

jpar tment of Materials Engineering, Ben-Gurion University of the NegevDepartment of Materials Engineering, Technion - I s r a e l I n s t i t u t e of

TechnologyFaculty of Engineering, Tel Aviv UniversitySchool of Applied Science and Technology, The Hebrew University of

JerusalemMinistry of DefenseNational Council for Research and DevelopmentIs rae l Atomic Energy Commission, Nuclear Research Centre - NegevIsrael Military IndustriesRAFAEL - Is rae l Armament Development AuthorityINTEL Electronics Ltd.ASHOT - ASHKELON Ltd.ISCAR BLADES Ltd.ISKOOR Ltd.URDAN, Associated Steel Foundries Ltd.

CONTENTS

I . TRANSFORMATIONS

Kinetics and Metastable Structure Formation in Rapid Solidif icat ionProcessing. (Invited).

D. Turnbull 1

Some Techniques for the Study of Atomic Motions with Applications toCeramic Materials. ( Invi ted) .

A.S. Nowick 11

An Atomic Resolution Study of Radiation-Induced Precipitat ion in FastNeutron Irradiated Tungsten (Rhenium) Alloys. (Invited) .

R. Herschitz, D.N. Seidman 21

Some Aspects of the Crystal l izat ion of Rare Earth-Noble MetalAmorphous Thin Film.

L. Shikhmanter, M. Talianker, M.P. Dariel 30

Dendritic Growth and Dendrite Arm Spacing in the Solidificationof S tee l .

M. Bamberger, I . Minkoff 35

Transition from Fibrous to Lamellar Morphology in UndierctionalSolidified Ni-W Eutectic.

S.F. Dirnfeld, Y. Zuta 39

The Microstructure of Rapidly Solidified Cu-Fe Alloys.A. Munitz, Z. Livne 46

Phase S tab i l i ty 5 Massive Transformation in Y203 - CompletelyStabilized Zirconia (Y-CSZ).

A.H. Heuer, R. Chaim, M. Ruhle 51

Microstructure Evolution S Ordering in Commercial MGO-PartiallyStabilized Zirconia (MG-PSZj.

R. Chaim, D.G. Brandon 55

Deformation Induced Decomposition of Uranium-Titanium Martensite.

G. Kimrnel, J. Sa r i e i , A. Landau, M. Talianker 59

I I . ALLOY PHASES AND STRUCTURE

Applications of Analytical Electron Microscopy to Materials.J . I . Goldstein, O.B. Williams, M.R. Notis 63

Field Electron and Ion Emission from Zirconiated and Zr FreeW Cathode.

J . Pelleg, J.L. Fink 67

The Transmission Electron Microscopy of NBi-5 Brazed Joint ofInconel 718.

B. Grushko, 0. Botstein, B.Z. Weiss 71

I I

Size Effect in Radiation Induced Segregation (RIS).L. Kornblit, A. Ignatiev 74

Applications of Advanced Color Computer Graphic DisplayTechniques to Materials Problems .Phase Equi l ibr ia § DiffusionalGrowth.

M.R. Notis, S.K. Tarby, J . I . Goldstein 77

Coercivity and Squareness Ratio in Co-W Thin Films.U. Admon, G. Kimmel, M.P. Dariel, E. Grunbaum, J.C. Lodder 82

Structural Analysis of High-Vacuum, High-Temperature BNi-5Brazed Joint of Inconel 718 Superalloy.

B. Grushko, B.Z. Weiss 88

Phase Relations in the Cu-Nd System on the Cu Rich Side.C. Laks, J . Pelleg, L. Zevin 92

Effect of Small Additions on Grain Refinement of a 14 CaratAu-Ag-Eu-Zn Alloy.

M. Fishman, L. Gal-Or, A. Iram . 95

Reinvestigation of the Pr-Ga System in the 66-100 at% Range.J . Pelleg, D. Dayan, G. Kimmel 100

Texture in Low Alloyed Uranium Alloys.J . Sa r i e l , G. Kimrael, J . Pelleg 104

I I I . MECHANICAL BEHAVIOR

Microstructure and Properties of Tungsten-Based Heavy Alloys.(Invited).

D.V. Edmonds 108

Investigation of the Creep Failure Mechanism in the Mo-5%W Alloy.A. Freund, D. Agronov, A. Rosen 119

Isochronous Creep of Copper-Beryllium-Nickel Alloy (Cu-0.4 Be - 2.0Ni) Solution Treated and Aged .

N. Nir 122

Hot Tearing of Lead Alloys.P. Arigur, F. Weinberg 128

The Embrittlement of Steels by Low Melting Point Metals.N. Breyer 133

The Effect of Substructure on Creep Properties of the TZM Alloy.D. Agronov, E. Freund , A. Rosen 137

Fracture Toughness Evaluation of Brittle Materials UsingIndentation Method.

Z. Nissenholz 142

I l l

IV. ENVIRONMENTAL BEHAVIOR

Failure of Welded Inconel-600 Pipe in the Cooling Systems of aNuclear Reactor.

G. Kohn, B. Herrman, E. Rabinovitz, A. S t e r n , S. Addess 146

Effects of Metal lurgical Variables on Hydrogen Embrittlement i nTypes 316, 321 and 347 S t a i n l e s s S t e e l s .

P. Rozenak, D. E l i eze r 152

Mar tens i t ic Transformations in 304L and 316L Types S ta in lessSteels Cathodically Hydrogen Charged.

E. Minkovitz, D. E l i e z e r 158

Mechanical Proper t ies Degradation of Hydrogenated Austeni t ic SS.I . Gi lad , Y. Katz, H. Mathias 163

Tensile Flow and Fracture Behaviour of Aus ten i t i c S ta in less S t e e lAfter Thermal Aging in Hydrogen Atmosphere.

Y. Rosenthal, M. Mark-Markowitch , A. S t e r n , D. E l iezer 167

Corrosion Behaviour of Al-Cu Alloy Thin Films in Microelect ronics .J . Zahavi, M. Rotel, H.C.W. Huang. P.A. To t t a . 175

V. PROCESSING AND TESTING

Quant i ta t ive Nondestructive Evaluation Using Ultrasonic Waves.(Invited)

L. Adler 185

On the Homogenization Problem in Sintered Alloys.L. Levin, A. Stern 201

Heat Transfer to Water and i t s Importance for Metal Casting andHeat Treatment.

M. Bamberger, B. Prinz 209

Fast, Non-Destructive Elechtrochemical Detection of SurfaceInclusions in Metallic Substrates.

I . Rubinstein 214

Telephone Tokens Produced by Powder Metallurgy.A. Sharon 216

Automatic Detection of Fluorescent Indicat ions.K.M. Jacobsen 221

Cryoforming of 301 and 302 Stainless Steel .T. Livni, S. Bar-Ziv, A. Rotem, A. Rosen 225

Ausforming of H-ll Mod SteelG. Elkabir, A. Rosen 232

IV

VI. SURFACE PHENOMENA

Thermodynamic and Kinetic Phenomena in Adsorbed Layers. ( Inv i t ed ) .M. Grunze

Supported Si lver Catalysts Have Some Important Propert ies in Commonwith Rough S i lver Films: EHipsometry § Raman Data.

P.H. McBreen, D. Hal l , J . Lalman, M. Moskovits

XPS Studies of Si Films Deposited from SiCl4 by an RF Cold PlasmaTechnique.

E. Grossman, M. Polak, A. Gri l l

Quantitative XPS of NaS-Alumina: Evidence for Sodium AnisotropicSegregation.

Y. Grinbaum, M. Polak

VII. SURFACE TREATMENT AND COATINGS

Structure of Protective Diffusion Coating for Niobium Alloys.M. Kazinets, 0. Gafri, L. Zevin, B. Rabin

The Effect of B and P Doping on the CI Concentration and theDeposition Rate of Si from SiCl 4 in RF Plasma.

R. Manory, E. Grossman, R. Avni, A. G r i l l

Silicon Ni t r ide Coatings by the Low Pressure RF Plasma Technique.U. Carmi, A. Raveh, A. Inspektor, R. Avni

Boridation of Steels in Low Pressure R.F. Plasma.A. Raveh, A. Inspektor, U. Carmi, E. Rabinovitz, R. Avni

Laser Induced Copper Elec t ro less Pla t ing.S. Tamir, J . Zahavi

Surface Hardening of Steel by Boriding in a Cold RF Plasma.I . Finberg, R. Avni, A. G r i l l , T. Spalvins , D. Buckley

VIII. MATERIALS AND PROCESSES FOR ELECTRONICS

Current Metal l izat ion Issues in Microelectronic Devices.( Invi ted) .

K.N. Tu

Schottky Barr ier Height for Ti-W on Si l icon.M.O. Aboelfotoh, K.N. Tu

In ter fac ia l Reactions in Laser Annealed Ni/GaAs Contacts.A. Lahav, T. Brat, C. Cyterman, M. Eizenberg

Enhanced Opto-Electronic Act ivi ty of Semiconductors by Means of(Photo) Electrochemical Etching.

R. Tenne, V. Marcu

235

245

249

253

258

262

267

272

277

283

288

291

299

304

Laser Induced Metal Deposi t ion on GaAs s u b s t r a t e s .J . Zahavi 308

Vapor Phase So lde r ing of Surface Mounted E l e c t r o n i c Assemblies.E. Falkenstein, I . Fainaro 315

Latt ice, Grain Boundary and Short Circuits Diffusion ofPhosphorus in TaSi Thin Films.

J . Pelleg 320

The LIMM Technique for Determination of the Spa t ia l Distributionof Polar izat ion of Space Charge in Polymer E l e c t r e t s .

S.B. Lang 324

Chalcogenide Infrared Glass FibersA. Bornstein, N. Croitoru 328

Programming of Crystal Diameter in Czochralski Growth by a CoolingPlot : Application to InSb.

M. Azoulay, Z. Burstein 332

IX. MATERIALS FOR ENERGY CONVERSION

The Dependence of the High Temperature, High Flux S tab i l i tyMaterials on Surface Structure and Composition. (Invited) .

A. Ignat iev 336

Transparent Conductor Films as a Material for PhotovoltaicJunctions with Polycrystal l ine Si l icon.

Z. Harzion, M. Zafir, J . Rishpon, S. Gottesfeld, N. Croitoru 343

Photoelectrochemical Characterization of ^ ,G. Dagan, G. Hodes, S. Endo, D. Cahen 347

MN. 15Pe~ oc - Hydride Compacts for Hydrogen Heat Pumps.

Y. Josephy, Y. Eisenberg, M. Ron 350

X. POLYMERS AND COMPOSITES

Crosslink Density of Polymers - Can it be Determined by SolventSwelling?

M. Gottlieb 356

Phase Separation in Rubber Modified Flame Retardant Epoxy Systems.Hemi Nae 361

Recent Advance i n the S t reng th and Time Dependent F a i l u r e Processof Kevlar Monofilaments and Composites.

H.D. Wagner, S.L. Phoneix, P . Schwartz 365

Deformation Processes in Impact Modified PVC.A. Hadas, A. Siegmann 369

Fat igue Crack Propagation Mechanisms in Polymers.A. Buss iba , Y. Katz, H. Mathias 373

AUTHORS' INDEX 378

KINETICS AND METASTABLE STRUCTURE FORMATION IN RAPID SOLIDIFICATIONPROCESSING

D. T u r n b u l l

Division of Applied Sciences, Harvard University, Cambridge, MA 02138 USA

ABSTRACT

The general procedures in metastable structure synthesis are surveyedand the reasons for kinetic preferences for metastable structures arediscussed. In melt solidification the motivation for forming metastablestates arises from the interfacial undercooling which results when highthermal gradients are imposed at the solidification front. The conditionsfor formation of glasses, supersaturated solutions and microcrystallinesolids in rapid solidification processing are discussed.

INTRODUCTION

The genera l procedure in low pressure me tas t ab le s t r u c t u r e syn the s i sis to energize a material--by melting, dissolution, irradiation or coldworking--and then deenergize i t by quenching or some condensation process[1,2]. To trap a metastable structure, the end stage of deenergizationmust be one in which the kinetic processes within the structure are toosluggish to alter its configurational state during the period of i t sintended study or use. Thus, the structure would be in a "configura-tionally frozen" rather than metastable equilibrium state. Nevertheless,it is common practice to apply to such states the label of the one fromwhich they were frozen.

Three types of metastability might be distinguished:

(11 Compositional--e.g. supersaturated solutions.(2) Topological--e.g. glasses and other amorphous solids and

certain crystalline phases.(3) Morphological--e.g. solids with a considerable density of

such extended imperfections as: dislocations, interphaseand intergrain boundaries.

Numerous examples of all three types of metastability have long beenrecognized. This paper surveys the formation of novel alloy structuresin the above three categories by the melt quenching technique. First,however, a digression on the problem of kinetic preference in metastablestructure synthesis is presented.

KINETIC PREFERENCE IN METASTABLE STRUCTURE FORMATION

Deenergization may open up p o s s i b i l i t i e s foT evolution of the systeminto some metastable, or oven unstable , s t a t e . I t s outcome wi l l be de ter -mined by: (a) what thermodynamic options develop, (b) which of theseoptions i s k i n e t i c a l l y p re fe r red , and (c'J whether the las t s tage ofdeenergization i s rapid enough t o k i n e t i c a l l y t r ap a meta- or unstables t a t e , i f formed.

Actual ly, much experience indicates t ha t metastable s t a t e s often arek i n e t i c a l l y preferred over the most s t ab l e ones. The ear ly experience oft h i s na ture so impiessed Ostwald that he formulated h i s well-known "stepr u l e " , t h a t , in deenergization a system evolves through the succession ofavai lab le metastable s t a t e s of decreasing f ree energy [ 3 ] . Such a ru lehas not been j u s t i f i e d t h e o r e t i c a l l y and i t must break down when thedriving f ree energy, AG, of forming the metastable s t ruc tu re i s su f f i -c i en t ly small. If force coupling cont r ibut ions are neglected the net fluxof matter, J j , in a pa r t i cu la r t ransi t ion is the product of a kineticfactor, p^, and a thermodynamic factor, f(AG^"), which reduces to AG:. inthe l inear kinetic regime, RT » | A G J | , thus :

p. = J. /f(AG.). (1)

Then the interesting question is how do the p- for the various possiblet rans i t ions correlate with the corresponding AG or other thermodynamicconstants.

While there are, indeed, many examples where kinetic preferences, asindicated by the p^, in evolving systems are different from those mostfavored thermodynamically, there are clear exceptions to t h i s behavior. Amore general pattern appears to be kinetic preference for the s ta te withentropy nearest to that of the i n i t i a l one. For example, in the down-quenching of vapor or l iquid, amorphous over crystal s ta te formation seemsalways to be preferred kinet ica l ly , while in crystal upquenching liquidalways forms faster than d i lu te vapor. Also, in phase separation, themost favored kinetic path is generally through a succession of s tates ofdecreasing dispersion. This pattern toward step entropy evolution inst ructural change may re f lec t that the processes most favored kinet ical lyare likely to be those requiring the least changes in the correlations ofthe positions and motions of atoms in ';he evolving group. For example, incondensation the atomic posit ional correlat ions which must develop areless extended for amorphous than for c rys ta l l ine s t ructures . However, thefavored sequence in the s tructural evolution of superheated quartz (quartz-*• liquid -*• Cristobali te) apparently departs from a step entropy pat tern.This path probably i s determined by some poorly understood features of theinterface structures (quartz - Cris tobal i te and quartz - l iqu id) . Itseems evident, also, that the step sequence in structural evolution may bealtered by heterogenous nucleation, e .g . , i t is likely that the f i r s tappearing phase in undercooled Fe alloy melts [4] should revert fromb . c . c . t o f . c . c . by addition of appropriately chosen heterogeneous nuclei .In these instances a step entropy preference, while not apparent in thebulk behavior, might s t i l l be exhibited by the interface processes.

From these considerations we expect that in rapid sol idi f icat ionprocessing formation of supersaturated solutions and microcrystall ine and

amorphous structures would be kinet ical ly favored. Whether or not thesestructures form will be determined by the magnitude of the thermodynamicdriving free energy which i s developed and t h e i r entrapment would bedetermined by the rate of quenching, T, through the configurationalfreezing temperature.

KINETICS OF SOLIDIFICATION

The optimum melt disposition for rapid quenching is as a thin layer

in good thermal contact with a massive heat sink of high thermal conduc-

tivity. The sink is usually composed of a crystalline metal or alloy with

a composition often very different from, but in some experiments identical

with, that of the melt. It is the undercooling, AT^, developed in the

melt at the interface during rapid quenching which may drive the formation

of non-equilibrium structures. AT- = Tβ - Tj where T^ is the liquidus and

Tj the interface temperature.

To sustain substantial undercooling a melt must be highly resistant

to homogeneous crystal nucleation. Simple ["classical") nucleation

theory relates the steady state nucleation frequency, I, to the scaled

undercooling, - T)/T£, as follov.s [5]:

~ nk. exp3(AT ) 2 T

(2)

where n = number density of melt atoms or moleculeskj = frequency of crystal-melt inter fac ia l rearrangement

3 = ASm/RAS,,, = molar entropy of melting, here assumed temperature independent

Tr =

T = absolute temperature of undercooled melt

assumed T independent and isotropic1 1/7

AHm

a = crystal-me±t inter fac ia l tensionAH,,, = molar enthalpy of meltingN = Avogadro's numberV = average molar volume

At a given value of the scaled crystal-melt interfacial tension, ex, andwith k^ in the range M/picosec., the calculated I (ATr) r ises sharply withATr from immeasurably low t o measurable levels at some onset undercoolingAT^(T] which depends on the quench rate , f, but only weakly in the highatomic mobility ("labile") range. Reflecting t h i s weak dependence, at thelowest real izable cooling r a t e AT'(T') approaches some minimum Thresholdvalue AT^Q which increases sharply with increase in a. At a > 1, 1 doesnot reach a measurable level at any undercooling.

When the interfacial rearrangement is thermally activated, i t s fre-quency decreases with fa l l ing T so that I(ATr) would be depressed to levelswell below those calculated with k; - 1/pic.osec. This effect has beenroughly evaluated for the case where k; scales as the shear relaxationfrequency. It is found [5] that with increasing scaled glass temperature.

T r g = Tg/T^ where T is the actual glass temperature, I(AT ) i s depressed,i t s maximum sharpens, and AT .Q i s increased. Also, the value of a atwhich the maximum in I(AT ) f a l l s below the measurable level decreases.

Experience indicates that AT' for pure metal and alloy melts i squite high, typical ly ^ 0 . 2 to 0.3 corresponding t o a ^ 0.6 whenk^ ^ 1/picosec. This AT^0 would be the upper limiting in te r fac ia l under-cooling in the rapid quenching of metal melts. I t s magnitude i s suffi-cient t o permit formation of structures which depart widely from that atequilibrium.

The movement of a planar crystal-melt front occurs in a sequence ofat least two steps: (1) in te r fac ia l rearrangement and (2) t ransport ofthe heat of crys ta l l iza t ion away from the interface . In alloys the move-ment may, in addition, be limited by transport of impurity from the in ter -face. In pure melts the speed, u, of the front is related to the in te r -facial rearrangement frequency, k j , as follows [6,7]:

AG/RT.u = fk.XCl-e *) (31

where f = fraction of surface sites at which growth can occurA = interatomic spacingAG = molar free energy of crystallization, equal to -AS AT.

at small departures from equilibrium

In the linear kinetic regime (RT » |AGJ1 equation (3) becomes

AT.u = fkjXB ~ (4)

i

At steady state u must also be proportional to the rate of transport ofthe crystallization heat from the front and neglecting temperature non-uniformity in the liquid, i t is related to (grad T)., the thermal gradientin the solid at the interface, as follows: ~l

K(grad T).Vu ~- sp-i- (51

m

where < i s the thermal conduct iv i ty in the s o l i d . Then from eqs . (4) and

K(grad T ) . VT.

fk.A BAH"3 m

SOLIDIFICATION OF PURE METAL MELTS

(61

Arguments have been presented elsewhere tha t in the c r y s t a l l i z a t i o nof pure metal melts (al f i s near un i ty , (bl t h e i n t e r f a c i a l rear range-ment process i s l imited only by the c o l l i s i o n frequency of atoms from themelt on the interface and i t needs no thermal ac t iva t ion f 8 ] . Thus, fk-A

should approach as an upper limiting value the sound speed, u s , in theliquid metal [1] . Also, i t follows that a melt in contact with a c rys t a l -l izat ion front would not form a glass having an appreciable lifetime atany rate of quenching.

Assigning typical values to the constants which appear in equation (6)(e.g^, fkjX = u5 = 4000 meters/sec, K - 0.2 cal/cm2sec.deg, 3 = 1 , andAHm/V = 300 cals/cm3) the maximum scaled interfacial undercooling at thecrysta l l izat ion front in a pure metal melt is estimated to be

A T i -q-=-=- = AT . ^ 1 x 10 " (grad T) . (7)

ij, ri l

Q

With the highest thermal gradients now realizable, M0 deg/cm inquenches (-T ^ 10*~deg/sec) following high energy pulses of a few pico-seconds duration [9], ATrj would only reach levels of order 0.10 whichare well below ATÔ, the scaled undercooling at measurable nucleationonset. Supporting this estimate, experiments of Lin and Spaepen [10]indicated that the crystallization front in pure liquid I-"e reachedvelocities as high as 500 meters/sec, in quenches following SO picosecondlaser pulses, corresponding to ATp 'v- 0.10 [eq. (4) with fk A *\J 4000meters/sec]. Thus, even at such quench rates there should be no measurablehomogeneous crystal nucleation in advance of the crystallization front inpure metal melts. Regrowth into pure molten layers produced by suchprocessing should be epitaxial as was observed by Lin and Spaep^n [101.However, the interfacial undercooling should be sufficient for the genera-tion of considerable densities of stacking raults and other extendedimperfections.

When liquid layers are quenched on substrates with compositions verydifferent from their own the crystallization front is formed by growth ofn"clei originating in the near vicinity of the melt-substrate interface,m-s. If the substrate structure is very different from that, of thecrystallized mets 1 the m-s temperature might fall to a level where-homogeneous nucleation occurs rapidly. Then a fine grained crystallizationfront would develop and evolve by further growth into a fine grainedcolumnar structure. A similar course of structural evolution would befollowed if nuclei formed heterogeneous ly a*: isolated sites on thesubstrate surface. Ther nucleation would begin at lesser undercooling andthe ultimate grain density/area would be proportional to the numberdensity of the sites activated.

In the foregoing it was assumed that the transient period for dei'el-opment of the steady r.ucleation frequency is negligible. A recent analysisof the transient by Kelton et a 1. [11] indicates that this assumption,while quite unsatisfactory in the glass transition regime, should beapproximately valid for the highly labile regime which appears to extendto Tr well below 0.5 for pure liquid metals.

SOLIDIFICATION OF ALLOY MELTS

Crystal Growth

In near equilibrium c r y s t a l l i z a t i o n homophase impurit ies in the melt

must, in general, be redistributed, but at high interfacial undercooling''partitionless" crystallization, i.e. that unattended by any long rangeimpurity transport, may become possible [12]. The thermodynamic condi-tions for non-equilibrium inclusion of impurities in growing crystals aswell as partitionless crystallization have been defined by Baker and Cahn[13,14] and recently reviewed by Boettinger et at. [15]. Partitionlesscrystallization may occur when the interfacial undercooling exceeds thatat temperature T0(x) where the free energies of crystal and melt alloyphases of identical composition, x, are the same. In this process, im-purity may actually be "trapped" in the crystal at a higher chemicalpotential than it had in the melt provided the total free energy of crys-tallization is negative. The course of T0(xJ for a binary alloy with afalling liquidus, T^(x) , is shown schematically in Figure 1. Often T0(x")is, for one or another alloy structure, well above 0°K at every compo-sition, meaning that at any x the alloy will have a lower energy in some"homogeneous crystallized chan in an amorphous solid form. However, thereare some alloy systems having solution thermodynamic parameters such thatthe To(x) for both terminal structures plunge to zero, say at x0 and xo,on non-intersecting courses so that at any x between x0 and xo the alloyat 0°K will be more stable as an amorphous solid than in any homogeneouscrystalline form. We have to consider two regimes, one with and the otherwithout solute partitioning, and the transition from one regime to theother, of crystal growth in alloy melts.

tT

L i q u i d u s , Tj ( x )

Limiting Temperature T 0 (x ) ofPartitionless Crystallization

Nucleation onset TemperatureTo'(x)

Glass Temperature/TgU)

Figure 1. Schematic of composition dependences of homogeneous nucleation,onset temperature, To(x), and maximum temperature To(x) forpartitionless crystallization.

Consider a planar front between a crystalline alloy and a thin layerof its melt where heat is extracted through the crystal. The front movesand the impurity redistributes itself in response to an imposed thermalgradient and impurity concentration in the body of the melt. Given theatomic and thermal transport coefficients and the thermodynamic parametersthe front velocity and impurity redistribution, including the conditionsat transition to partitionless growth, can, in principle, be calculated[16].

When impurity is rejected into the melt the velocity of the frontshould at steady state, and if the liquid solution is dilute, satisfy therelation:

D.u = - î- (grad c) .

i

where c. = the impurity concentration in the melt at the interface i.(grad c)j = the concentration gradient in the melt at i.

Dj = the impurity diffusivity in the melt at i.

The upper limiting speed of the front imposed by this relation is

D.

With values D^ 2 * 10~5cm2sec~ ' and X ^ 2 x lutein, typical of liquid melts,u £ 10 meters/sec, which is more than 2 orders of magnitude below themaximum velocity possible f^4000 meters/sec) in collision limited growth.The impurity transport in melts is thermally activated and its rategenerally falls to negligibly low values as the temperature is loweredthrough the glass transition range.

?n the partitionless regime the formal analysis of the growth rate issimilar to that for pure metals, but with TQ(x~) replacing the crystal-meltequilibrium temperature. However, two limiting microscopic growth mechan-isms may be distinguished. In one, crystallization proceeds with no changein impurity near-neighbor coordination, i.e. without reordering of thecompositional short range order (C-SRO) of the melt. Then the growth couldbe collision limited and k-jX might, as in pure melts, approach the soundspeed. In the other limit where some essential reordering of the C-SRO isnecessary for crystallization, k^X would be approximately equal to Dj/X.As when partition occurs, this growth should be thermally activated andbecome negligible at temperatures well below T

Kinetic analyse; [16,17] indicate that the transition from growth withto that without partition should occur within about one order of magnitudeof growth velocity centering at u^Dj/X. In this transition, the actualsolute distribution coefficient goes from near equilibrium to a level nearunity corresponding to complete solute trapping by the crystal.

Nucleation

Experience indicates that in the labile range the temperature To(x)at measurable nucleation onset in alloy melts roughly parallels theliquidus Tp (x). This behavior was accounted for by Thompson and Spaepen[18] by assuming that

(a) The scaled crystal-melt tension, a, as in Spaepen's model, ismainly topological in origin and so only weakly dependent on x.

(b) The crystal nucleus is always in interaction with melt ofaverage composition.

If Tβ(xj falls sharply with x and if, as is often observed, Tg(x)varies l i t t l e with x, TQ(x) on a course parallel with Tg(x) would crossT (x) at some composition. However, provided k always scales with D j ,a! x increases, T'(x) will fall below its parallel course and drop tozero before reaching i ts projected intersection with Tg(x).

Glass Formation

From the foregoing i t follows that a necessary condition for metalglass formation in melt quenching and, indeed, in condensation on sub-strates containing crystal nucleants, is That some redistribution ofimpurity, either by partition or short range reordering, be thermodynami-cally demanded for crystallization [20,5,8,7]. Then the interfacialrearrangement frequency has to scale with Dj and so would be thermallyactivated and suppressed to negligible levels at temperatures below To.

In addition to the impurity redistribution requirement the meltquench rate between Tofx) and Tp(x) must exceed the critical value, -f_r,for bypassing homogeneous crystal nucleaticn. Applying simple nucleationtheoi/ with the scaled tension a^0.6, as indicated by the nucleationresistance of pure metals, and assuming that kj scales with reciprocalshear viscosity as described by the Vogel-Fulcher equation. Spaepen andthe writer [21] related -T to the scaled glass temperature, Tra- Therelation indicates that -fcr should fall sharply with increasing Trg to106 deg/sec at T r g ^l/2 and 1 deg/sec at T r g ^ 2 / 3 . Since Tg(x) isrelatively insensitive to composition Trg(x) should increase and -T c r

decrease with a falling liquidus T#(x) so that -fcr should be near aminimum at minimum T^(x).

These considerations indicate that those alloy melts with T r g >_ 0.4and in which crystal growth is reconstructive and thermally activatedshould be capable of forming glass when quenched at the highest rates nowrealizable.

Boettinger [22] has demonstrated that in certain alloy systems wherepartitioning in crystallization is required, the compositions most favora-ble for glass formation are, indeed, those falling in the estimated x0 tox0 range over which the alloy is more stable in amorphous solid than inhomogeneous crystalline form. In contrast, recent, experiments of Linand Spaepen [23] clearly show that some alloys with T (xl well above Tg(x)can be quenched to glasses. Apparently, in these alloys the kineticresistance to short range reordering, without partition, retarded crystalgrowth sufficiently for T to be reached without appreciable crystalli-zation. Some observations of Bagley and the writer [24] indicated aninteresting transition, with a composition change of only 0.1 at.%, fromthermally activated partitionless to partitioning growth of Ni P crystalsin amorphous films. The partitionless growth resulted in a monophasicpolyhedral crystal morphology while growth with segregation requiredtemperatures more than 75° higher foT an equivalent growth rate and led toa dendritic morphology.

Grain Structure of Rapidly Solidified Alloys

When partitionless crystal growth occurs without thermal activition,it appears that homogeneous nucleation in advance of the crystallizationfront should be as unlikely as in pure metals and so the grain morphologyof the alloy should resemble that of pure metals crystallized undersimilar conditions.

When growth is thermally activated, there are conditions where fineequiaxed grain structures should develop. In the crystallization of thinmolten alloy layers, attended by impurity redistribution, the constitu-tional undercooling can become very large, even with high thermalgradients into the solid at the interface. Thus, interface shape insta-bilities might develop so that crystallites could be injected into themelt by the "meltback" mechanism [25]. However, the minimum times foronset of shape instability or the liquid thickness needed for initiationof the meltback mechanism have not been evaluated. Such evaluationsmight indicate a minimum thickness of molten alloy film required forforming an equiaxed grain structure.

Under the conditions that growth is thermally activated and Trg(x) isrelatively high, in sufficiently rapid quenches the interface temperaturemay be lowered to levels below To(x) so that the frequency of homogeneousnucleation in advance of the interface becomes appreciable. The alloymelt might then crystallize to an equiaxed grain structure or, at asomewhat higher quench rate, undercool to a glass in which a high numberdensity of microcrystallites are embedded. The latter structure wouldevolve to an equiaxed microcrystalline one upon reheating. The individualcrystallites might be mono- or bi-phasic depending on the alloy composi-tion. Thompson et at. [26] have suggested that the number density of cry-stallites formed in treatments of this type can be limited considerably bythe transient times needed to reach steady state nucleation. Experience,indeed, indicates [27] that certain glass forming alloys crystallize toequiaxed fine grained structures when quenched at intermediate rates. Alsothere is considerable experience which shows that glasses of such alloysform similar structures, often composed by biphasic grains, upon reheating.

ACKNOWLEDGMENTS

Much of this paper was written during the author's stay at theNational Bureau of Standards [November 1983) as a visiting scientist inthe Center for Materials Science. The Harvard phase of the research wassupported by ONR Contract N00014-77-C-0002 and MRL NSF Contract DMR80-20247. The author is indebted to W.J. Boettinger and S.C. Coriell(NBS) and M.J. Aziz, A.L. Greer and F. Spaepen (Harvard) for useful dis-cussions.

REFERENCES

1. D. Turnbull, Met. Tra->s. 12A, 695 (1981).2. D. Turnbull, Ann. Rev. Mat. Sci. (ed. R.A. Huggins, R.H. Bube and

D.A. Vermilyea) 13_, pp. 1-7 (1983).3. See discussion in M. Volmer, "Kinetik der Phasenbildung," pp. 200-05,

Steinkopff, Dresden (1939).

10

4. R.E. Cech, Trans. A.I.M.E. 206, 585 f!956).5. D. Turnbull, Contemp. Phys. 10, 473 (1969).6. D. Turnbull and M.H. Cohen, "Modem Aspects of the Vitreous State"

(J.D. Mackenzie, ed.) !_, p. 38, Butterworth's, London (1960).7. F. Spaepen and D. Turnbull in "Laser Annealing of Semiconductors"

(J.M. Poate and J.W. Mayer, eds.), pp. 15-4], Academic Press, NewYork (1982).

8. D. Turnbull and B.G. Bagley in "Treatise on Solid State Chemistry"(ed. N. B. Hannay) 5_, p. 513, Plenum, New York (1975).

9. See N. Bloembergen, "Laser-Solid Interactions and Laser Processing"(S.D. Ferris, H.J. Leamy and J.M. Poate, eds.), A.I.P. ConferenceProc. 50_, pp. 1-10 (1979).

10a. C.J. Lin and F. Spaepen in "Chemistry and Physics of Rapidly Solidi-fied Materials" (B.J. Berkowitz and R.O. Scattergood, eds.), pp.273-280, A.I.M.E. Conference Proc. (1983).

10b. See also S.C. Coriell and D. Turnbull, Acta Met. 30 , 2135 (1982).11. K.F. Kelton, A.L. Greer and C.V. Thompson, J. Chem. Phys. 79_, 6261

(1984).12. H. Biloni and B. Chalmers, Trans. Met. Soc. A.I.M.E- 2S3, 375 (1965).13. J.C. Baker and J.W. Cahn, Acta Met. 17_, 575 (1969).14. J.C. Baker and J.W. Cahn, "Solidification", pp. 23-58, Am. Soc.

Metals, Metals Park, Ohio (1970).15. W.J. Boettinger, S.R. Coricll and R.F. Sekerka, in pres.-;, Mats. Sci .

and Eng. (1984).16. M.J. Aziz, J. Appl. Phys. 53_, 1158 (1982); see also for references to

earlier literature.17. M.J. Aziz, Appl. Phys. Lett. 43_, 552 (1983).18. C.V. Thompson and F. Spaepen, MRS Symposia Proc. (ed. G.E. Rindone)

9_, p. 603, Amsterdam, North Holland (1982); see also, Acta Met., 51_,2021 (1983).

19a. F. Spaepen, Acta Met. 23, 729 (1975).19b. F. Spaepen and R.B. Meyer, Scripta Met. H), 257 (1976).20. D. Turnbull, "Solidification", pp. 1-22, Am. Soc. Metals, Metals

Park, Ohio (1970).21. F. Spaepen and D. Turnbull, "Rapidly Quenched Metals" (N.J. Grant

and B.G. Giessen, eds.) pp. 205-229, MIT Press, Cambridge. Mass.22. W.J. Boettinger, "Proc. of 4th Intl. Conf. on Rapidly Quenched

Metals" (T. Masumoto and K. Suzuki, eds.) J_, pp. 99-102, JapanInst, of Metals, Tokyo (1982).

23. C.J. Lin and F. Spaepen, Appl. Phys. Lett. 4J_, 716 (1982).24. B.G. Bagley and D. Turnbull, Acta Met. 18, 857 (1970); see Figs. 5

and 7.25. K.A. Jackson, J.D. Hunt, D.R. Uhlmann and T.P. Seward III, Trans.

Met. Soc. A.I.M.E. Z56_, 149 (1966).26. C.V. Thompson, A.L. Greer and A.J. Drehman, "Proc. of the 4th Intl.

Conf. on Rapidly Quenched Metals" (T. Masumoto and K. Suzuki, eds.)J_, pp. 743-746, Japan Inst. Metals, Tokyo (1982).

27. A.J. Drehman, Ph.D. Thesis, pp. 100-102, Division of Applied Sciences,Harvard University, Cambridge, Mass. (1983).

11

SOME TECHNIQUES FOR THE STUDY OF ATOMIC MOTIONS WITHAPPLICATIONS TO CERAMIC MATERIALS

A.S. Nowick

Henry Krumb School of MinesColumbia University, New York, N.Y. 10027, USA

ABSTRACT

A review is presented of the use of the techniques of a.c. impedancemeasurements, d ie lec t r i c re laxat ion aril anelast ic relaxat ion to studytechnological ly in terest ing ceramic mater ia ls . I l l us t ra t i ons are given ofthe appl icat ion of these methods to the study of so l id e lec t ro ly tes aswell as quartz crystals used fo r frequency con t ro l .

INTRODUCTION

Knowledge of rates of atom movements plays an important ro le in theunderstanding of most mater ia ls . In the case of ceramic ( ion ic ) mater ia ls,the migrat ing species are of ten charged ions, and then t he i r migrationrates can be studied by e l ec t r i ca l means. Such studies are of par t icu larinterest i n the case of materials that are r e l a t i v e l y good ion ic conduc-to rs , the so-cal led so l id e lect ro ly tes or "superionic conductors" whichare of in te res t as e lec t ro ly tes for so l i d -s ta te batteries and fuel ce l l s .Such appl icat ions are of su f f i c i en t in teres t tha t in recent years therehas been an outpouring of review books [1 - 3] and conference proceedings[4-7] in t h i s area. Materials studied include those that ccd- jc t by mi-grat ion of a l ka l i ions, s i l v e r ions, protons, oxygen i o " arn f k j . i n eions, among others.

The present paper w i l l review three techniques that have been ac-t i ve l y used i n recent years i n the author's laboratory, w i th i l l u s t r a t i o n sof the kind of informat ion, both basic and appl ied, derivable from them.The techniques are: (1) a.c. impedance measurement, (2) d i e l e c t r i c relax-a t ion , and (3) anelastic re laxat ion . We w i l l see that valuable informa-t ion about technological ly important systems has been obtained usingthese techniques.

A.C. IMPEDANCE MEASUREMENT AND ANALYSIS

An ideal ion ica l l y conducting material should be represented elec-t r i c a l l y as a capacitance C in paral le l wi th a resistance R, the formerrepresenting the d i e l e c t r i c properties of the medium and the l a t t e r i t sconduction. Real materials are not so simple, however, due to a number ofcomplications. F i r s t , there i s the problem of introducing and dischargingthe conducting ionic species at the respective electrodes. Second, onefinds another e f fec t , loosely cal led the "grain-boundary" e f f e c t , which isdue to blocking of car r ie rs at internal in ter faces wi th in the e lec t ro l y te .As a resu l t o f these factors the actual equivalent c i r c u i t of a sample canbe extremely complex. A r e l a t i v e l y simple representation i s shown in

12

Fig. l a , consisting of three R-C units in series with each other, one rep-resenting the bulk or l a t t i c e behavior, another the grain-boundary effectand the f inal one, the electrode effect. In order to study such an equiv-alent c i r c u i t , a.c. measurements are made over as wide a frequency range("frequency window") as possible (usually 1 Hz to 106 Hz). The sample canthen be represented by i t s complex admittance, Y* containing both real andimaginary parts (representing the current in-phase with and 90° out-of-phase with the applied voltage). One may write

Y* = G(w) + iuC(u) (1)

where G(ui) is the effective conductance and C(CD) the capacitance both, ingeneral, functions of the frequency. We may also introduce the reciprocalquantity, Z* = 1/Y*. called the complex impedance, which also has real andimaginary parts

Z* = Z'(u>) - iZ"(u) (2)

A convenient and widely used analysis is the complex impedance plot inwhich Z" is plotted as a function of Z

1 with to as parameter. For the case

of the equivalent circuit shown in Fig. la, such a plot yields three semi-circular arcs, one for each R-C unit, as shown in Fig. lb. In this plotthe frequency increases as we go from right to left (arc l to arc 3). Inthis way, it becomes possible to separate out the true lattice conductionfrom the grain-boundary and electrode phenomena, and to study eachseparately.

(a)Rμ, R.

—wv

0

c.

50 Lac to.c.

500 •

(W

5067

•

©

•c

/

Vo

•

1

02: 6 V. 1^0}

0.92 »10"Zalm

t.O » 10"*a!m

Fig. l a . Equivalent c i r - u i t repre-senting la t t ice [I), grain boundary(gb) and electrode (e) effects.

Fig. l b . Schematic diagram of com-plex impedance plot correspondingto c i r c u i t of part (a). Arrowindicates direction of increasingfrequency.

Fig. 2. Examples of compleximpedance plots for CeO2:6%Y2O3 at three different temper-atures. Arcs are labeled tomatch Fig. lb . From Ref. [ 8 ] .

13

Actually, the experimental "frequency window" is usually not wideenough to display all three effects at any one temperature, but by cover-ing a range of temperatures, we can see all of them. Figure 2 shows anexample for the case of Y 3 + doped CeO2- This material is an oxygen-ionconductor through the migration of oxygen-ion vacancies which are intro-duced into the lattice as charge compensation for the lower valerit cationdopant. Such ionic conductors may be used as the solid electrolyte inhigh-temperature fuel cells or in oxygen sensors. From Fig. 2, we seethat at Tow temperatures (e.g. 178 °C) arcs 2 and 3 appear while at hightemperatures arcs 1 and 2 are obtained. Further we see that only arc 1,the electrode arc, has a strong dependence on partial pressure of oxygenof the ambient gas. These results show, however, that the behavior ismore complex than that predicted from the schematic diagram of Fig. lb.Firstly, the arcs are not full semi-circles but are somewhat depressed (asshown by the sloping lines drawn in Fig. 2). This result can be inter-preted as meaning that each arc corresponds to a narrow distribution ofR-C circuit elements rather than to single values. For simplicity weignore this fact for the present. Secondly, it is found that the elec-trode contribution can involve more than a single arc. This will be dis-cussed below. Nevertheless, the model of Fig. lb can serve as the basis;T extracting appropriate parameters. The two intersection points R,~and R23 are clearly related to the equivalent-circuit parameters by

R23and

R,12so that R12 -

R23 gives Rinto a conouctivit

<gb

(3)

(4)

•y, a,

a = d/AR

i\ . . Each of these resistances can be convertedin the usual way by the relation

(5)

where d is the thickness and A the area of the sample. The results maythen be plotted on a conductivity plot, log oT vs. 1/T, as shown in Fig. 3.Here the lower curve (solid triangles) is obtained from R,, and the uppercurve (open triangles) from R23. In addition, results of Tour-probe d.c.measurements, which are plotted as open circles, show excellent agreementwith R,2, thus firmly establishing that both arcs 2 and 3 are due to the

Fig. 3. Conductivity plots for aCeO2:l% Y2O3 sample. Upper curve,obtained from the intersection R23represents the lattice conductivity.Lower curve, consisting of a.c.measurements from R12 intersection,as well as d.c. 4-probe data, isdominated by the grain-boundaryeffect. From Ref. [9].

2.4 ZJo

14

electrolyte. Both curves can be fitted to the conventional Arrhenius-typerelation

al = B exp (-H/kT) (6)

in which H is the activation enthalpy.

We have obtained useful information from each of the three arcs.From arc 3 we have obtained the lattice conductivity as a function oftemperature, composition and dopant [9,10]. Considering that oxygen-ionvacancies are introduced as charge compensation for the dopant ions, onemight expect a monotonically increasing conductivity as the dopant concen-tration increases. It is found instead that a goes through a sharp maxi-mum as a function of dopant concentration as shown for Y2O3 dopant inFig. 4. At the same time, the activation enthalpy goes through a minimum.These results show that strong defect interactions suppress the conductiv-ity beyond a concentration of 4-6 mole % dopant. The study of defectinteractions in systems of high concentration is a subject of great inter-est and one which requires further study. Similarly, it was found thatfor different trivalent cation dopants of the same composition, the high-est conductivity and lowest activation enthalpy H occur for a dopant whosesize is closest to that of the host cation [10]. Such a result means thatstrain energy terms play an important role in determining H, a result thathas recently been verified by computer simulation calculations [11]. Froma practical viewpoint, these studies show how best to optimize the latticeconductivity with respect to dopant species and dopant concentration.

In most applications of solid electrolytes, one wishes to achieve themaximum possible overall d.c. conductivity. Accordingly the grain-boundaryeffect cannot be ignored. In fact, Fig. 3 shows that the overall conduc-tivity can be as much as lOOx lower than the lattice conductivity due tothe grain-boundary effect. Similar effects are prevalent for other solidelectrolytes, e.g. in β-alumina which is the electrolyte in the much pub-licized sodium-sulfur battery [12]. In the present case, particularlybecause Hgb » r'-i (see Fig. 3), we can only interpret the electrical

(ev) .

O-β -

0.6

(il-cm)',-1

05 .1

Fig. 4. Variation of conductivitya (et 182 °C) and of activationenthalpy with composition, forCeOo - YoOq solid solutions. FromReff [9]f

J

15

effects in terms of the existence of a continuous (very poorly conducting)blocking layer, presumably associated with grain boundaries. In order toverify the presence of such a layer and to determine its origin, we exam-ined thinned (ion-milled) samples by STEM (scanning transmission electronmicroscopy) in combination with microanalysis by EDAX (Energy DispersiveX-ray Microanalysis) and EELS (Electron Energy Loss Spectroscopy) [13].A number of microstructural features were observed in this study, but theone most relevant seemed to be the presence of "thick boundaries", asshown in Fig. 5, i.e. layers of SOOA thickness which seem to be contin-uous. These were generally found to be an amorphous phase which had Si asits major cationic constituent. This silica-type phase did not surroundeach grain, however, but appeared around a small agglomerate of grains.In order to determine whether the presence of Si as an impurity is indeedresponsible for the grain-boundary effect, we have very recently succeededin preparing doped ceria samples from starting materials that were essen-tially silicon free. Our measurements show that in such a material thegrain boundary effect is virtually eliminated.

The remaining arc (arc 1 in Figs, lb ar. 2) is that due to electrodeeffects. Actually electrode processes are very complex and rarely giverise to just a single arc [14]. For the case of oxygen-ion conductors, wehave carried out extensive studies with porous Pt-paste electrodes as wellas a variety of others [15]. The simplest behavior (i.e., a single de-pressed arc) was observed for freshly prepared Pt-paste electrodes at rel-atively high oxygen partial pressures. With the aid of an auxiliary tech-nique, called the current-interruption method [15], it was possible toshow that, in this simplest case, the electrode process is controlled by acharge-transfer mechanism, in which an adsorbed oxygen adatom 0a(j under-goes the reaction

°ad 2e" + vo = + Vadwhere VQ and Vacj are la t t ice oxygen vacancy and vacant adsorbed s i te , re-spectively. The cathodic reaction goes from left to right and the anodic

Fig. 5. STEMmicrostructure of asintered CeO2:6%Gd2(h sample show-ing thick bound-aries" of anamorphous silicaphase. From Ref.[13].

16

in reverse. At low oxygen pressures or when the electrode has aged underhigh temperature or high current, an increase in Z" on the low-frequencyside of the arc takes place. This is related to a transport limited pro-cess, or "concentration polarization."

DIELECTRIC AND ANELASTIC RELAXATION

Because of important similarities we will treat these two techniquestogether. In both cases an alternating field is used, electric and stressfields, respectively. We measure the fractional energy loss per cycle inthe form of the quantity tan 6, where 6 is the angle by which the response(polarization or strain, respectively) lags behind the applied field. Inboth cases the simplest behavior takes the form of the well-known Debyepeak as a function of frequency:

tan 6 = A • U)T/(1 + U2T

2) (7)

where A is a measure of the strength of the relaxation process and T isthe relaxation time. In the case of dielectric or anelastic relaxationsdue to defect pairs (or higher clusters), A is proportional to the defectconcentration and to the square of the electric or elastic dipcle strengthof the defect, respectively. On the other hand, the reciprocal x"

1 is

related to the frequency of the controlling ionic migration process and istherefore thermally activated:

T-1 = v

0 exp (-H

r/kT) (8)

where Hr is the activation enthalpy for relaxation. Because of Eqs. (7)

and (8), it is possible to observe tan 6 either as a peak in frequency atconstant temperature or a peak in temperature at constant frequency. Formore complex processes, tan 6 must be written as a summation over expres-sions of the type Eq.(7),and the corresponding peak is then broader than asimple Debye peak [16].

For dielectric relaxation an alternative to the a.c. tan 6 measurementis available. It is called the TSDC (thermally stimulated depolarizationcurrent) method (also known as ITC) [17]. This method offers an extremelyhigh sensitivity and therefore can detect the reorientation, under elec-tric field, of electric dipoles at concentrations as low as 1 ppm. It hasbeen applied to ceramics of the type already discussed [18], but for lackof space we will not be able to cover this work here.

To illustrate the use of these relaxation techniques we turn insteadto a different problem, that of the frequency instability of quartz crys-tals. It is well known that α-quartz, due to its piezoelectric property,can be fabricated into resonators that vibrate at a fixed frequency. Acommon application is to control time pieces (i.e. watches). But applica-tions of such resonators to controlling satellites and guidance systemsdemand far greater frequency stability, often to as much as 1 part in 10

9

or 101 0, even in radiation environments. Frequency instabilities under

irradiation are found to be closely linked to impurities present in thesecrystals, one of the most important of which are alkali ions, e.g. Na

+ or

Li+ [19]. These alkalis are present in the crystal as compensation for the

impurity A l3 +, which sits on Si"

+ site and is deficient one positive charge.

In recent years, techniques have been developed for replacing alkalis byH

+ or one alkali with another by a process of electrodiffusion or

17

50 -

* 30

1 10<-> oo -10s2 -30-50 -

QUARTZ OSCILLATORSACCUMULATED OFFSET

10 MeV ELECTRONS

1

—

Z-GROWTHD t f f P SYNTHETIC

\- Q BBS El A AAAAAAtt_ ° O O O

O

oo

°o

1 1

SWEPT Z-GROWTHSYNTHETIC

^ N A T U R A LI

104 105 ' 10w

RADS (Si)10'

Fig. 6. tffect ofelectron irradiationon the room-tempera-ture frequency change,6f, of 5 MHz quartz-crystal resonators.From Ref. [21].

"sweeping" [20]. In this process the crystal is treated at elevated tem-perature under a strong d.c. electric field in an appropriate environment.Also, methods of hydrothermal growth have made possible the production ofsynthetic quartz crystals that are often far lower in impurity contentthan natural crystals. Figure 6 shows the effect of electron irradiationat room temperature on three quartz crystals, a synthetic and a natural,both as grown, and on the same synthetic crystal after sweeping to replacealkali ions with H+. The scale of frequency changes in this graph is verylarge compared to the best present-day requirements. Nevertheless, thefigure shows that H+ sweeping greatly reduces the frequency change thattakes place under irradiation. Thus we see that alkalis contribute to 6f.

One reason for a frequency offset at room temperature can be theoccurrence of an anelastic relaxation peak at lower temperatures. The

* .f0-100 '» '«> INIMIWJJO

Fig. 7. Anelastic loss (in-ternal friction, Q"1) of a Na-doped quartz crystal as a func-tion of temperature. FromRef. [20].

18

formal theory of anelasticity [16] shows that the presence of in anelasticpeak given by Eq. (7) produces a frequency depression

Sf/ f = - A/2 (9)

at all temperatures above the peak, where 5f is the change in frequencyrelative to a defect-free crystal. In other words, to understand a fre-quency offset at room temperature requires that we know about relaxationpeaks that occur below room temperature. Such peaks are indeed prevalent.For quartz crystals containing Na+, a prominent pair of anelastic peaks isfound at low temperatures, as shown in Fig. 7. A comparable pair of peaksis observed in dielectric relaxation [22]. These have low activationenergies (Hr = 0.05 and 0.14 eV, respectively) and are attributed to theNa+ being bound to a substitutional Al 3 +, with the Na+ residing in the in-terstitial channel. No such peaks due to Al-Li pairs have been found inspite of a careful search for them [23], presumably because the Li + sitson a 2-fold axis of symmetry. Irradiation produces important changes inthese defects. Electron-hole pairs generated serve to free alkali ionsfrom the Al 3 + and to create Al-h (aluminum-hole) pairs instead. The mi-gration of the alkali can be observed as conductivity enhancement immedi-ately after irradiation [24]. It is not yet clear which defects serve asthe traps at which the alkali ions terminate, but such trapping sites prob-ably also capture one or more electrons. The Al-Na peaks are eliminatedor reduced by the irradiation and, instead, new large peaks can be observedboth by anelastic and dielectric loss measurements. Figure 8 shows suchdielectric loss peaks for both Na+- and Li+- containing crystals. Similaranelastic peaks have been observed [25]. It is not yet clear as to whetherthese peaks are due to Al-h centers or to alkali centers which are createdby the trapping of Li+ or Na+ freed by the irradiation. Nevertheless,

120

100

80

60

4 0

20

.TO5

-

-

A.

T0Y0

ni 1/

\ \

1IT////

i •

SUPREME

\ .

\

t

tIRRAD.) •

/

jj^Li SWEPT

' '" No SWEPT

_^ H SWEPT

J 1 1 L

Fig. 8. Dielectric loss of Li-swept, Na-swept and H-sweptsynthetic quartz followingroom-temperature X-irradiation.From Ref. [24].

0 4 8 12 16 20 24 28 32100/T

19

Eq. (9) shows that the elimination of a low-temperature anelastic relax-ation peak by irradiat ion results in an increase in frequency, while thecreation of a new peak results in a frequency decrease. Thus, i t is clearthat the solution to questions concerning frequency instab i l i t ies at roomtemperature l ies in a better understanding of these low-temperature relax-ation phenomena.

In conclusion, i t was intended to show how the methods describedherein give insight into ionic migration in ceramic materials. These meth-ods can be usefully complemented by other widely used techniques such asdiffusion and NMR measurements. Together, they give the type of insightinto defect migration mechanisms that make i t possible to eliminate prac-t i ca l problems or to determine conditions of maximum performance.

ACKNOWLEDGMENTS

The author wishes t o acknowledge the c o n t r i b u t i o n s o f h is present andformer assoc ia tes : D rs . Da Yu Weng, H. J a i n , J . Toulouse and R. Gerhardt -Anderson, and the suppor t received f rom both the U.S. Department o f Energyand the U.S. A i r Fo rce , RADC.

REFERENCES

1. S. Geller, ed., "Solid Electrolytes", Topics in Applied Physicsvol. 21, Springer-Verlag, Berlin, 1977.

2. P. Hagenmuller and W. van Gool, eds., "Solid Electrolytes", AcademicPress, New York, 1978.

3. E.C. Subbarao, ed., "Solid Electrolytes and Their Applications",Plenum Press, New York, 1980.

4. W. van Gool, ed., "Fast Ion Transport in Solids", North-Holland,Amsterdam, 1973.

5. G.D. Mahan and W.F. Roth, eds., "Superionic Conductors", Plenum Press,New York, 1976.

6. P. Vashishta, J.N. Mundy and G.K. Shenoy, eds., "Fast Ion Transport inSolids", North-Holland, Amsterdam, 1979.

7. J.B. Bates and G.C. Farrington, eds., "Fast Ionic Transport in Solids",North-Holland, Amsterdam, 1981.

8. Da Yu Wang and A.S. Nowick, J. Solid State Chem. 35 (1980) 325.

9. Da Yu Wang, D.S. Park, J. Griffith and A.S. Nowick, Solid State Ionics2 (1981) 95.

10. R. Gerhardt-Anderson and A.S. Nowick, Solid State Ionics 5 (1981) 547.

11. V. Butler, C.R.A. Catlow, B.E.F. Fender and J.H. Harding, Solid StateIonics 8 (1983) 109.

12. R.W. Powers and S.P. Mitoff, J. Electrochem. Soc. 122 (1975) 226.

20

13. R. Gerhardt-Anderson, A.S. Nowick, M.E. Mochel and I. Dumler, Proc.Conf. High Temperature Solid Oxide Electrolytes, ed. F.J. Salzano,Brookhaven National Laboratory 1983, vol. 1, p. 225.

14. J.R. Macdonald, in "Electrode Processes in Solid State Ionics",M. Kleitz and J. Dupuy, eds., p. 149, Reidel Publ. Co., Dortrecht-Holland, 1976.

15. Da Yu Wang and A.S. Nowick, J. Electrochem. Soc. 126 (1979) 1155, 1166;127 (1980) 113; 128 (1981) 55.

16. A.S. Nowick and B.S. Berry, "Anelastic Relaxation in CrystallineSolids", Academic Press, New York, 1972.

17. C. Bucci, R. Fieschi and G. Guidi, Phys. Rev. 148 (1966) 816.

18. Da Yu Wang and A.S. Nowick, J. Phys. Chem. Solids 44 (1983) 639.

19. J.C. King and H.H. Sander, Rad. Effects 26 (1975) 203.

20. D.B. Fraser, in "Physical Acoustics" vol. 5, W.P. Mason, ed., p. 59,Academic Press, New York, 1968.

21. B.R. Capone, A. Kahan, R.N. Brown and J.R. Bucfrnelter, IEEE Trans, onNucl. Sci. NS-13 (1966) 130.

22. D.S. Park and A.S. Nowick, Phys, Stat. Sol. (a) 26 (1974) 617.

23. J. Toulouse, E.R. Green and A.S. Nowick, in "Proc. 37th Ann. Symp. onFrequency Control", 1983, p. 125, U.S. Army EP.ADCOM.

24. H. Jain and A.S. Nowick, J. Appl. Phys. 53 (1982) 485.

25. J.J. Martin, L.E. Halliburton and R.B. Bossoli, in "Proc. 35th Ann.Symp. on Frequency Control", 1981, p. 317.

21

THE CHEMISTRY ON A SUBNANOMETER SCALE OF

RADIATION-INDUCED PRECIPITATION AND SEGREGATION

IN FAST-NEUTRON IRRADIATED TUNGSTEN-RHENIUM ALLOYS

Roman Herschitz and David N. Seidman

Cornell Un.versity, Department of Materials Science and Engineeringand the Materials Science Center, Ithaca, New York 14853-0121

ABSTRACT

The phenomena of radiation-induced precipitation and segregation have

been investigated in W-10 at.% Re and W-25 at.% Re alloys, employing

the atom-probe field-ion-microscope technique. The W-10 at.% Re alloy

is subsaturated with respect to the solvus line of the primary solid

solution (3 phase), while the W-25 at.% Re alloy is supersaturated

with respect to the same solvus line. The specimens had been irradia-

ted in the Experimental Breeder Reactor II to a fast-neutron fluence

of 1.4x10 neutrons cm"2 (E>0.1 MeV) at 575, 625 and 675 C. This

corresponds to 8.6 dpa and an average displacement rate, for the iiwo

year irradiation time, of 1.4x10" dpa s ~ . The results of the present

investigation show a very significant alteration of the microstructure

of both alloys as a result of the fast-neutron irradiation. In the

case of the W-10 at.% Rβ alloy coherent, semicoherent and possibly

incoherent precipitates with the composition ^WRe and a disc-shaped

morphology — one or two atomic planes thick — were detected at a

number density of n<10^ cm , and a mean diameter of o>57 A. For the

W-25 at.% Re alloy coherent, semicoherent and incoherent precipitates

with the composition ^WRe3 were detected; the precipitate's number

density is vLoJ-7 cm"

3 with a mean diameter of 40 X. None of the WRe

precipitates or the WRe3 coherent precipitates were associated with

either line or planar defects or with any impurity atoms. Therefore,

a true homogeneous radiation-induced precipitation occurs in these

alloys. The semicoherent WRe3 precipitates were associated with He

atoms; that is, these precipitates may have been heterogeneously nuc-

leated. In the W-25 at.% Re alloy a two dimensional WRe3 phase has

been observed at a grain boundary. A physical argument is pr:^rnted

for the nucleation of WRe or WRe3 precipitates in the vicinity of

displacement cascades produced by primary knock-on atoms. It is

suggested that in both cases the first step in the nucleation of a

precipitate is due to the formation of tightly-bound mobile mixed

dumbbells which react to form an immobile di-rhenium cluster. Possible

sequences of point-defect reactions which can lead to either WRe or

WRe3 cluster are detailed. The further growth of a cluster (WRe or

WRe3> into a precipitate is most likely driven by the irreversible

vacancy: self-interstitial atom annihilation reaction, as suggested

recently by Cauvin and Martin.24 Point-defect mechanisms for all the

other observations are also discussed.

INTRODUCTION

Over the last few years there has been a rapid growth of interest inthe phenomena of radiation-induced (as opposed to accelerated) segre-gation and precipitation.1-4 Different types of irradiation — elec-trons, ions or neutrons — can induce significant segregation of alloy-ing elements either toward or away from grain boundaries, voids or freesurfaces. Radiation can also cause the heterogenous or homogeneousprecipitation of a phase in subsaturated solid solutions and it canalter the phase stability of alloys. Radiation-induced segregationand precipitation are of paramount technological importance, sincethey play a crucial role in the nucleation and growth of voids and havea strong effect on the physical properties of metals alloys used inthe fuel cladding and core structure of the fast-breeder reactor, aswell as in the materials used for the first wall of fusion reactors.These phenomena are also of considerable fundamental interest.

The study of W(Re) alloys is of technological importance, as they areused in thermocouples for the measurement of temperature in nuclearreactors. As a result of an exposure to a fast-neutron flux thedecalibration of W(Re) thermocouples occurs. ' The alloys W-10 at.%Re and W-25 at.% Re are of particular interest in understanding thephenomenon of radiation-induced precipitation, as the former alloy issubsaturated with respect to the solvus line of the primary solidsolution (6 phase), while the latter alloy is supersaturated withrespect to this solvus line — it is in the 3 plus a phase field.7'Sikka and Moteff^ and Williams e_t al_- have identified the crystalstructure of radiation-induced precipitates in fast-neutron irradiatedW-25 at.% Re alloys using transmission electron microscopy — forspecimens which had been irradiated at 1100 C and higher — and itcorresponds to the y.-phase which has the composition WRe3- Williamset al . 10 also investigated fast-neutron irradiated W-5 at.% Re andW-ll at.% Re alloys. For specimens which had been irradiated at 1100 Cand above all the precipitates analyzed by electron diffraction wereconsistent with the x-phase crystal structure. Whereas Williams e_tal_- were unable to obtain interpretable electron diffraction patternsfrom any of the specimens which had been fast-neutron irradiated at900 C or lower.

In this short summary paper we present the results of an extensiveatom-probe field-ion microscope (FIM) study of radiation-inducedprecipitation and segregation in fast-neutron irradiated W-10 at.% Reand W-25 at.% Re alloys. ' Our atom probe FIM allows us to deter-mine the chemical identity of all the elements in the periodic tabled3"16

In addition, the atom-probe FIM has a lateral spatial resolution, forchemistry, of a few-tenths of a nanometer and a depth resolution whichis determined by the interplanar spacing of the region being analyzed.

We found very significant alterations of the microstructure between575 to 675 C. In the case of the W-10 at.% Re alloy precipitateswith the composition ^WRe (o phase) were detected at a number densityof VL0l6cm~3. They were not associated with linear or planar defectsor with any impurity atoms; i.e., a true homogeneous radiation-inducedprecipitation occurs in this alloy. Coherent, semicoherent and inco-herent precipitates were detected. For the W-25 at.% Re alloy coherent,semicoherent and incoherent precipitates with the composition

23

(X phase) were detected at a number density ^1017cm~3. Thecoherent precipitates were not associated with either line or planardefects, or with any impurity atoms. This strongly suggests that thecoherent WRe3 precipitates were a result of a homogeneous radiation-induced process. The semicoherent and incoherent MJRej precipitateswere found to be associated with He atoms; i.e., they may have beenheterogeneously nucleated. In addition we found evidence for a two-dimensional t»WRe3 phase at a grain boundary; this phase is the resultof a radiation-induced segregation process.

EXPERIMENTAL DETAILS

Wire specimens of W(Re) alloys were irradiated to a fast neutron-fluence of i-4xlO22 neutrons cm"2 (E>0.1 MeV) at elevated temperatures(575, 625 and 675°C) in Experimental Breeder Reactor II (EBR-II) atRichland, Washington. This corresponds to 8.6 dpa for row 7 of EBR-II.Hence, the average displacement rate for the two year irradiation timeis 1.4xlO~7 dpa s"1.

The wir3 specimens were electroetched into sharply-pointed FIM speci-mens. Next the specimens were analyzed chemically by the atom-probetechnique at an ambient pressure of ^4x10"^ Torr, with the specimensmaintained at 45 K. A pulse fraction (f) of 0.15 was used for all theanalyses. The quantity f is the ratio of the pulse voltage to thesteady-state dc voltage. A constant pulse frequency of 60 Hz wasemployed. The average field-evaporation rate — average number ofions evaporated per field-evaporation pulse — was equal to 0.02 ionspulse . Using these experimental conditions we were able to obtaingood agreement between the nominal Re concentration, and the Re con-centration as determined by the atom-probe technique for unirradiatedalloys. These experimental conditions were used in all of our chemicalanalyses.

The specimens were imaged employing 3He as an imaging gas. The reasonfor.using -%e, rather than He gas, was to minimize the concentrationof He present in the atom probe and therefore, to make it possibleto identify *He atoms which have had their origin in the neutron-irradiated specimens.

The basic mode of displaying the data in the present experiment is inthe form of an integral profile. A Re integral profile is obtainedby plotting the cumulative number of Re events versus the cumulativenumber of W plus Re events. The average slope of such a plot corres-ponds to the average Re concentration of the volume analyzed, sincethe cumulative number of all events detected is proportional to depth.In analyzing a particular precipitate the slope of an integral profile

lower limit to the actual Re concentration inpitate (<cgP >*), as in most cases the dimensions of the analyzedcylinder are greater than the size of a precipitate. The superscriptppt stands for precipitate, the subscript u on the bracket means anuncorrected value and the superscript * implies a corrected value. Therelationship between <cgpt> and <cPPfc>* for different possible preci-pitate morphologies is presented in Appendix A of Herschitz and Seidman.

24

EXPERIMENTAL RESULTS: W-10 AT.% Re

Four radiation-induced precipitates were detected and analyzed, whereasno voids were found in the W-10 at-% Re alloy. The density of theradiation-induced precipitates is equai to ilO an" ; it was determinedfollowing the procedure used by Brenner and Seidman. '

The following summarizes the main experimental results:

(1) Coherent, semicoherent and possibly incoherent precipitateshave been observed. The number density of precipitates is<vLOl6cm~3.

(2) The observed precipitates are disc shaped — one or twoatomic planes thick. And their mean diameter is '- 57 A.

(3) The composition of the radiation-induced precipitates corres-ponds to 'MtfRe; that is, <cPPfc>* 52 at.% Re. This resultindicates that the WRe precipitates in the W-10 at.% Realloy are radiation resistant in the temperature range 575 to675 C — in the presence of a fast-neutron flux.

(4) The precipitates were not associated with either linear orplanar defects, or with any impurity atoms; i.e. a truehomogeneous radiation-induced precipitation occurs in thisalloy.

(5) No voids were detected in this alloy. This indicates thatthe addition of 10 at.% Re to W suppresses void formation,as voids have been detected in pure tungsten — which hadbeen subjected to fast-neutron irradiation.

EXPERIMENTAL RESULTS: W-25 AT.% Re

Six precipitates, three voids, a grain boundary and a region immediatelyadjacent to a grain boundary were chemically analyzed by the atom probetechnique.

The following summarizes the main experimental results for this alloy:

(1) Coherent, semicoherent and incoherent precipitates have beenobserved. Their n\diameter is ^40 A.observed. Their number density is M.0cm and the mean

(2) The precipitates observed have either a disc shaped or sphe-rical morphology.

(3) The composition of the radiation-induced precipitates corres-ponds to "WRe^; i.e., < c ^ >* is approximately equal to 75at.% Re.

(4) The coherent precipitates (^WRe3) were not associated witheither linear or planar defects or with any impurity atoms;i.e., a true homogeneous radiation-induced precipitationoccurs in this alloy.

(5) The semicoherent and incoherent precipitates were associatedwith ^He atoms; i.e., heterogeneous precipitation may haveoccured in this alloy.

25

(6) Voids at a number density of M.0 cm and a mean diameterof 90 A have been detected in this alloy. No significantRe enrichment or depletion at these voids had occured.

(7) Formation of a two-dimensional ^WRe3 phase has been observedat a grain boundary.

DISCUSSION

W-10 at.% Re alloy

The fact that the precipitates in the subsaturated alloy are not asso-ciated with either structural defects or with any impurity atoms indi-cates that a true homogeneous radiation-induced precipitation occursin this alloy. Experimental evidence for homogeneous radiation-inducedprecipitation has been presented recently by Cauvin and Martin in thecase of Al (Zn) alloys, 8'19 by Brager et_ al^.20 in the case of a 316stainless steel, by Mukai and Mitchell for a Ni (Be) alloy, ^ andKinoshita and Mitchell and Wahi arid Wollenberger23 for Cu (Be) alloys.Theoretical treatments of this physical phenomenon have been consideredby Cauvin and Martin and Maydet and Russell; ° the latter authorsonly considered the possibility of the nucleation of incoherent preci-pitates, whereas Cauvin and Martin also considered coherent precipi-tates .

We now describe a possible sequence of plausible events which can resultin the homogeneous micleation of WRe (the o phase) precipitates, in asubsaturated alloy which is subject zo irradiation with fast neutrons. '

The primary source of radiation damage, in the case of fast neutrons, isthe displacement cascade. Each displacement cascade is created by aprimary knock-on atom (PKA) with a mean recoil energy of 4 keV. In thecase of pure tungsten it is known from FIM experiments that a displace-ment cascade, created at 15 K, consists of a vacancy-rich core ( 2 to30 at.%) surrounded by a distribution of self-interstitial atoms (SIAs),which is created by the replacement collision sequence mechanism.The concentration of SIAs on the periphery of a displacement cascadecan be as high as 'VL to 3 at.%. Since the radiation damage is highlylocalized in the displacemenr cascade — the point defect supersatura-tion in between the displacment cascades is initially negligible — itis probable that the nucleation of a WRe precipitate occurs in itsvicinitv The absolute efficiency of this nucleation process is aslow as the final density of radiation - induced WRe precipitates is^lO-Lfc>cm~J> — which is significantly less than the number density ofPKAs that produce displacement cascades.^

Employing the known properties of point defects in W and W(Re) alloysit can be demonstrated that plausible first steps in the nucleation ofa WRe (a phase) precipitate involve the migration of tungsten SIAs toRe atoms to form mobile mixed dumbbells — in the immediate vicinity ofa displacement cascade — which in turn react to form an immobile di-Recluster.11 The di-Re cluster can then grow by the accretion of mixeddumbbells, and pure tungsten or rhenium SIAs. Specifically, the forma-tion of a WRe cluster is envisaged to occur via the following possiblereactions: (a) two mixed dumbbells react to form an immobile di-Recluster; (b) the di-Re cluster reacts with a pure tungsten SIA to

26

form a WRe2 cluster; and (c) the WRe2 cluster reacts with a second

tungsten SIA to form W2Re2 (or 2WRe) cluster. During the course of

the two year irradiation the displacement cascades dissolve slowly (see

Appendix B of Herschitz and Seidman-'--'-) and they provide the vacancies

which can result in the shrinkage of clusters. Recent experiments by

Averback and Ehrhart^8 on Ni-1 at.% Si also suggest strongly that point

defect clustering and trapping reactions occur in the vicinity of dis-

placement cascades.

Further specific details of the growth or shrinkage of a cluster are

difficult to state, but they can be rationalized in terms of the

Cauvin-Martin model for radiation-induced metastability. The physical

basis of the Cauvin-Martin model is that the irreversible vacancy-

SIA annihilation reaction drives solute clusters towards a larger solute

content and hence to precipitation.

A plausible mechanism for the suppression of void swelling, in this

alloy, involves the dominance of vacancy: SIA recombination over the

destruction of these point defects at a biased sink — the dislocation.

This is possible, in particular, by the recombination of vacancies

wieh SIAs which are trapped in immobile clusters involving SIAs and

rhenium atoms. This strong recombination process prevents the accumu-

lation of a sufficient number of vacancies for the nucleation and

growth of voids. This mechanism for the suppression of voids is con-

sistent with the mechanism suggested for the homogeneous nucleation of

WRe (o phase) precipitates.

W-25 at.% Re alloy

A very striking observation is the detection of coherent precipitates

with the composition WRe3, in the W-25 at.% Re alloy, which are not

associated with either structural defects or impurity atoms. The latter

observation strongly suggests that they were homogeneously nucleated.

The basic problem is to explain the radiation—induced precipitation of

WRe3 if the x phase is not in thermal equilibrium with the primary Cβ)

solid solution between 575 to 675 C.

The x phase may also be nucleated by the homogeneous nucleation mecha-

nism suggested in the previous subsection for the a phase in the vici-

nity of displacement cascades. The formation of a WRe^ cluster is en-

visaged to occur via the following reactions: (a) two mixed dumbbells

react to form an immobile di-Re cluster; (b) the di-Re cluster reacts

with another mixed dumbbell to form a Re3 cluster; and (c) the Re

3

cluster reacts with a pure tungsten SIA to form a WRe3 cluster. Once

again the further growth (or shrinkage) of this elementary cluster can

be understood to occur via the Cauvin-Martin model**4 for radiation-

induced metastability.

A necessary condition for the formation of a large number density of

WRe3 precipitates is that the nucleation current of these immobile

WRe3 clusters be considerably greater than the nucleation current of

immobile WRe clusters — in the W-25 at.% Re alloy. In the previous

subsection we suggested three point defect reactions which can lead

to the formation of a W2Re

2 (or 2WRe! cluster. If it is assumed that

these nucleation reactions, as well as the three reactions postulated

above for the nucleation of a WRe3 cluster, are in detailed balance

27

then it is readily shown that t

3c <* c cWRe3 W-W Ii

and2 2

CW Re " CW-wcii '

where cw_w is the concentration of tungsten SIAs and the constants ofproportionality are the products of the rate constants of the individualpoint defect reactions that lead to WRe3 of fr?2Re2 (2WRe) . The aboveequations show that c^g should be greater than c w R e . This isbecause the value of c^ grows more rapidly — in the direct vicinityof a displacement cascade — in the supersaturated W-25 at.% Re alloythan in the W-10 at.% Re alloy; note that cw_w is initially the samein both alloys and is in the range 10~2 to 3*10~2 at. fr.

This model for the homogeneous nucleation of WRe3 precipitates, from asupersaturated W-25 at.% Re alloy, shows qualitatively that the nuclea-tion current of immobile WRe3 clusters is greater that of W2Re2 (2WRe)clusters. However, the model does not explain why in the subsaturatedW-10 at.% Re alloy the Re concentration of the precipitates stops at^50 at.% Re. This question stands as an unsolved problem at present.

4Another very interesting observation is the detection of He atoms insidesemicoherent or coherent participates. The detection of 4He atoms inthis alloy is at first glance somewhat surprising, as the cross sectionfor 4He production in pure tungsten is quite small. The cross sectionfor the production of Tie atoms on rhenium atoms is not available,however L.R. Greenwood (Argonne National Laboratory, private communica-tion) estimates that the 4He production rates on Re and W should bequite similar. Thus, for a fluence of 8.6 dpa the estimated 4He con-centration is "^4x10" 3 appm in W-25 at.% Re; this assumes the displace-ment threshhold energy of a Re atom is identical to that of a W atom— 52 eV.23 This suggests that the 4He atoms were most likely producedon the impurity atoms present in this alloy. The elements B, C, N, Oand S all have rather large cross sections for the production of Heatoms (L.R. Greenwood, private communication). The absence of Heatoms in precipitates in the W-10 at.% Re alloy ^ may simply be a resultof a lower level of impurity atoms in this particular alloy.

Herschitz and Seidman 2 discuss four possible mechanisms which canresult in the detection of 4He atoms in the WRe3 precipitates. Twoof the mechanisms imply that the WRe3 precipitates were heterogeneous lynucleated and two that they were homogeneously nucleated. There is noobvious way to distinguish among these four mechanisms. Hence, we areleft with the distinct possibility that the semi or incoherent WRe2precipitates were heterogeneously nucleated.

In the voids analyzed by the atom probe technique we were unable todetect 4He atoms. The specimen temperature was 45 K during the chemicalanalyses/ hence, we can rule out the possibility of 4He atoms diffusingout of the voids as they were dissected, since the measured mobilityof He atoms at 45 K is extremely small.30"32 The most likely reasonwe were unable to detect 4He atoms in the voids was simply that thevolume fraction of the void analyzed was small — typically much less

28

than 0.1, and since the absolute number of 4He atoms per void is notexpected to be very large, the probability of detecting one ^He event isvery small.

The corrected Re concentration in a grain boundary was determined tobe "W5 at.% Re. This value corresponds to WRe 3 (x phase). The grainboundary Re concentration was found to fall to the bulk value (25 at.%Re) in *v>4 A. Thus, the x phase forms along a grain boundary and wehave an example of a two-dimensional phase, which is the result of anon-equilibrium radiation-induced segregation process. The atomicmechanism for the formation of this phase can be explained in terms ofthe migration of mixed dumbbells to the grain boundary. This mechanismis consistent with our suggestion that the WRe3 precipitates form in aW-25 at.% Re alloy, subject to a fast-neutron radiation field, as theresult of a homogeneous nucleation process which involves mixed dumb-bells reacting to form an immobile di-Re cluster.

ACKNOWLEDGEMENTS

This research was supported by the U.S. Department of Energy. Additionalsupport was received from the National Science Foundation through theuse of the technical facilities of the Materials Science Center atCornell University. We wish to thank Mr Robert Whitmarsh for enthusias-tic technical assistance, Dr Alfred Wagner (now at Bell Laboratories)for preparing the specimens for irradiation, Dr Martin Grossbeck (OakRidge National Laboratory) for arranging for an irradiation in EBR-II,Dr Robert S Averback (Argonne National Laboratory) and Dr Avner Brokman(Hebrew University) for useful discussions, Dr L R Greenwood (ArgonneNational Laboratory) for kindly performing calculations for us employingthe ENBF/B-V code, and Dr Georges Martin (Centre d1Etudes Nucleairesde Saclay) for useful questions and comments on the more extendedversion of this manuscript.

a) Presently at R.C.A. Corp., Astroelectronics Division, Princeton,New Jersey 08540.

b) Presently on a leave of absence at the Hebrew University ofJerusalem, School of Applied Science and Technology, BergmannBuilding, 91904 Jerusalem, Israel.

REFERENCES

1. J. Nucl. Mater. £3_, 1 (1979).2. J.R. Holland, L.K. Mansur and D.I. Potter, editors, Phase Stability

During Irradiation (Metallurgical Society of AIME, Warrendale,Pennysylvania, 1981).

3. F.V. Nolfi, Jr., editor. Phase Transformations During Irradiation(Applied Science, London, 1983).

4. K.C. Russell, Phase sLability Under Irradiation, to appear inProgress in Materials Science (Pergamon Press, Oxford).

5. W.E. Browning, Jr. and C.E. Miller, Jr., in Fourth Symposium onTemperature: Its Measurement and Control in_ Science and Industry(Reinhold Publishing Company, 196 2), Vol. 3, part 2, pp. 271-276.

6. R.L. Shepard and B.H. Montgomery, Oak Ridge National LaboratoryReport No. 5108, Oak Ridge Tennessee, November 1976, pp. 418-427.

29

7. J.M. Dickinson and L.S- Richardson, Trans, of the ASM j> , 758 (1958).8. M. Hansen and R. Anderko, Constitution of Binary Alloys (McGraw-Hill,

New York, 1958), pp. 1153-1154.9. V.'K. Sikka and J. Moteff, Metall. Trans. 5_, 1514 (1974).10. R.K. Williams, F.W. Wiffen, J. Bentley, and J.O. Stiegler, Metall.

Trans. A 14A, 655 (1983).11. R. Herschitz and D.N. Seidman, Cornell Materials Science Center

Report No. 5014 (1983). To appear in Acta Metallurgica (1984) .12. R. Herschitz and D.N. Seidman, Cornell Materials Science Center

Report No. 5015 (1983). To appear in Acta Metallurgica (1984) .13. A'. Wagner, T.M. Hall, and D.N. Seidman, Rev. Sci. Instum. 46,

1032 (1975) .14. T.M. Hall, A. Wagner, A.S. Berger and D.N. Seidman, Scripta Metall.

1(3, 485 (1976) .15. T.M. Hall, A. Wagner and D.N. Seidman, J. Phys. E: Scient. Instrum.

10_, 884 (1977) .16. A. Wagner, T.M. Hall and D.N. Seidman, J. Nucl. Mater. 69_ and 70,

413 (1978) .17. S.S. Brenner and D.N. Seidman, Radiat. Effects 2A_, 73 (1975).18. R. Cauvin and G. Martin, J. Nucl. Mater. J33_, 67 (1979) .19. R. Cauvin and G. Martin, Phys. Rev. B 23_, 3333 (1981).20. H.R. Brager and F.A. Garner, J. Nucl. Mater. T3/ 9 dq78) .21. T. Mukai and T.E. Mitchell, J. Nucl. Mater. 105, 149 (1982).22. C. Kinoshita and T.E. Mitchell, Electron. Micros. £, 236 (1980) .23. R.P. Wahi and H. Wollenberger, J. Nucl. Mater. 113, 207 (1983) .24. R. Cauvin and G. Martin, Phys. Rev. B 23_, 3322 (1981).25. S.I. Maydet and K.C. Russell, J. Nucl. Mater. 64_, 101 (1977).26. L.A. Beavan, R.M. Scanlan and D.N. Seidman, Acta Metall. 19_, 1339

(1971) .27. C.Y. Wei and D.N. Seidman, Philos. Mag A 43_, 1319 (1981).28. R.S. Averback and P. Ehrhart, to appear in J. Phys. F: Metal Phys.

(1984) .29. M.I. Current, C.-Y. Wei and D.N. Seidman, Philos. Mag. A 47_, 407

(1983) .30. A. Wagner and D.N. Seidman, Phys. Rev. Lett. £2_, 515 (1979) .31. J. Amano, A. Wagner and D.N. Seidman, Philos. Mag. A 44_, 177 and 199

(1981) .32. J. Amano and D.N. Seidman, Cornell Materials Science Center Report

No. 4153 (1982). To appear in J. Appl. Phys. (1984).

30

SOME ASPECTS OF THE CRYSTALLIZATION OF AMORPHOUSCOPPER-RICH COPPER-HOLMIUM THIN FILMS

L. Shikhmanter^, M. Talianker^'-1 and M,P. Dariel^b>C')

a) Materials and Process Engineering, Engineering Division,Israel Aircraft Industries Ltd., Lod, Israel.

b) Department of Materials Engineering, Ben-Gurion Universityof the Negev, P.O. Box 653, Beer-Sheva, Israel.

c) Nuclear Research Centre-Negev, P.O. Box 9001, Beer-Sheva,Israel.

ABSTRACT

Transmission electron microscopy was used to study the compositiondependence of the crystallization of amorphous copper-rich Cu-Ho thinfilms. The phases formed during the crystallization of these films cor-respond to those predicted by the equilibrium phase diagram. In addition,the crystallization of a hexagonal, phase which so far has not been repor-ted in rare earth-copper systems, takes place. Some aspects of the mor-phology of the precipitates crystallizing within an amorphous matrix arediscussed.

1. INTRODUCTION

In recent years amorphous alloys have attracted considerable atten-tion. Amorphous metals not only exhibit technologically promising proper-ties, they also constitute convenient materials for the study of the phy-sical aspects of the amorphous-crystalline phase transformation in solids.In our previous studies we have observed that the crystallization of rareearth (R)-Cu and R-Ag amorphous films of near-equiatomic composition re-sults in the formation of an equilibrium intermetallic compound [1-3]. Inorder to investigate the composition dependence of the crystallizationprocess we have also studied the crystallization of amorphous copper-richCu-Ho thin films.

2. EXPERIMENTAL METHODS

Amorphous thin films, approximately HOoR thick, were prepared byvacuum evaporation at 10"G Torr of an appropriate crystalline alloy ontoa carbon film substrate held at room temnerature. The initial Cu-Ho crys-talline alloys were prepared by arc-melting mixtures of the holmium andcopper that was followed by homogenization anneal in evacuated quartzcapsules.

The crystallization process was observed in situ during the variousannealing treatments by using a JEOL 200B transmission electron micros-cope equipped with a hot stage attachment.

31

In the course of this study thin amorphous films having the compo-sitions 68 at % Cu-Ho, 74 at % Cu-Ho and 78 at % Cu-Ho were investigated.

3. RESULTS AND DISCUSSION

The crystallization of the amorphous 68 Cu-Ho film begins at 190°Cand results in the formation of lense-shaped crystals (Fig.l) of theorthorhombic C£Cu2-type H0Q12 intermetallic compound.

As the original crystal grows new, differentlyoriented HoCu2 crystallites emerge from its ex-ternal surface, Fig.2a. The new crystalliteshave a twin-like orientation with respect tothe original parent crystal. A schematic re-presentation of this system is shown in Fig.2b.The twin planes are the [103] planes and the[002] planes contain the direction of the fastgrowth of the HoCu2 crystals in the amorphousmatrix.

Cu-Ho amorphous film afterFig.l. The structure of a thin 68 atcrystallization at 190°C.

The crystallization of amorphous 74Cu-Ho film starts at 260°C andleads to the appearance of the crystalline phase which consists of alarge number of parallel thin plates, Fig.3a. The selected area diffrac-"tion (SAD), pattern taken from this phase can be indexed in terms of ahexagonal unit cell with parameters a = 11.45A and c = 16.82A. As far aswe know, there exists no previous report on a stable hexagonal phase atroom temperature in the 20-25 at % R-Cu concentration range. Franceschi,in his studies of the R-Cu (R = Nd, Gd, Dy) [4r6] systems, suggested onthe basis of the observed thermal effects, the existence of a high-temperature R2CU7 phase. According to Franceschi, this phase cannot beretained at room-temperature, its structure therefore had not been de-termined. Possibly, the hexagonal phase, which we observed in the courssof the crystallization of 74Cu-Ho films corresponds actually to a com-'-ound with the same stoichiometry, namely H02CU7

In contrast to crystallization taking place at 260°C, the isother-mal annsal of the amorphous 74Cu-Ho film at 280°C results in the simul-taneous crystallization of two different morphologies, denoted A and B,of the crystalline phases as shown in Fig. 3b.The analysis of the SAD-pattern, taken from this sample, revealed thatthe A-type crystalline phase is identical with the hexagonal phase, ob-tained during the isothermal anneal at 260°C, and that the B-type phasehas a tetragonal structure with parameters identical to the parametersof the Dy2Cu9 compounds [4].

32

002103

Fig.2 a - The H0C112 crystal grown in the amorphous matrix withemerging twins; b - The schematic representation ofthe crystal shown in (a).

Fig.3. The crystalline phases formed in the amorphous 74 at % Cu-Hofilm during isothermal crystallization at a) 260°C; b) 280°C.

33

The crystallization of amorphous 78Cu-Ho thin films begins at 240°C and resultsin the emergence of elongated rhombic-likecrystals (Fig.4) having the same tetragonalstructure as was determined for crystalsformed during the crystallization of 74Cu-Ho film at 280°C.

The detailed morphology of the preci-pitates crystallizing within the amorphousmatrix is the result of highly complex in-teractions involving the volume change as-sociated with the crystallization, the re-lative values of the elastic constants ofthe precipitate and the matrix and the va-lues of the various precipitate matrix in-terfacial energies.

Fig.4. The crystal formed in an amorphous 78 at % Cu-Ho film duringisothermal annealing at 240°C.

The crystals of RCu and RAg phases with cubic CsC£-type structure,formed during the crystallization of amorphous films, have a circularshape 1-3 . Such a shape was also observed for the crystallization ofot-Fe crystals (bcc type structure) from Feeo(Ci-xBx)2o amorphous alloys[7], Y-FeNi crystals (fcc-type structure) from the FeôNiuoBao amorphousalloy [&] and crystals of a Pd-rich fee phase from the Pds3Sii7 amor-phous alloy [9].

We may assume that high symmetry cubic crystals possess a relati-vely low anisotropy of their elastic constant. Neglecting the possibleanisotropy of interfacial energy terms in the first approximations, wemay define our problem as that of determining the morphology of isotropic precipitates in an isotropic medium. Elastic theory [1O] predictsthat if the precipitate is stiffer than the matrix (higher value of theelastic constants) it will display a spherical shape. Such a sphericalshape will transform into a cylindrical one in a thin film, when thedimensions of the precipitate are of the order of the thickness of thefilm. This indeed was observed for the growth of the equiatomic CsC£-type RCu and RAg compounds.

Crystals with a lower than cubic symmetry crystallizing within theamorphous matrix may possess, in principle, a higher degree of anisotro-py of their elastic constants. Again neglecting surface energy effectswe may expect such crystals to display departures from the sphericalmorphology exhibited by the cubic-type crystal. This indeed was observedfor the crystals appearing in the copper-rich Cu-Ho amorphous matrix.The crystals with orthorhombic, hexagonal and tetragonal structures areformed during crystallization of these amorphous films. All these crys-tals have an elongate.", rorm (Figs. 1-4). A similar elongated shape ofthe crystals was also observed for crystallization of orthorhombic Fe3Bphase from the Fe7s-B25 amorphous alloy 111], and for the hexagonal Te-Se tshase within amorphous Te-Se films [12],

34

REFERENCES

1. L. Shikhmanter, M. Talianker and M.P. Dariel, Thin Solid Films,90, 51 (1982).

2. L. Shikhmanter, M. Talianker and M.P. Dariel, J. Phys. Chem. Solids44_, 745 (1983).

3. L. Shikhmanter, M. Talianker and M.P. Dariel, Bulletin of the IPS,29_, 80 (1983).

4. E. Franceschi, J. Less Common Met., 87^ 249 (1982).

5. M.M. Carnasciali, G.A. Costa and E.A. Franceschi, J. Less-CommonMetals 92^ 97 (1983).

6. M.M. Carnasciali, S. Cirafici and E. Franceschi, J. Less-CommonMetals 92^ 143 (1983).

7. K. Shimomura, P.H. Singii and R. Ozaki, J. Mat. Sci. 15_, 1175 (1980).

8. K. Muller, M. Von Heimendahl, J. Mat. Sci. 17_, 2525 (1982).

9. M. Scott, Rapidly Quenched Metals III, (Edited by B. Cantor), Vol.1,198, Metals Society, London, 1978.

10. D.M. Barnett, J.K. Lee, H.I. Aaronson and K.C. Russel, ScriptaMetal1., 8, 1447 (1974).

11. J.S. Vermaak and J. Petruzzello, J. Appl. Phys. 53_, 6809 (1982).

12. M.M. Carnasciali, S. Cirafici and E. Franceschi, J. Less-CommonMet. 81, 115 (1981).

35

DENDRITIC GROWTH AND DENDRITE ARM SPACING IN THE

SOLIDIFICATION OF STEEL

M. Bamberger & I. Minkoff

Department of Materials Engineering

Technion, Israel Institute of Technology

Haifa, Israel

INTRODUCTION

Dendritic growth is characteristic of materials growing under conditions

of thermal or diffusional instability. While non-metallic materials may

be observed to grow bounded by low index crystallographic planes,

metallic growth is invariably unstable for even small values of under-

cooling. The crystals grow in a characteristic manner shown in Figure 1.

The difference between the case for metals and that for non-metals is in

the relatively small values of surface energy for metals between solid

and liquid phases. The higher values for non-metals act to stabilise

the crystal faces in growth.

Figure 1 shows the characteristic parameters of a metallic dendrite -

the tip radius P and the spacing between the dendrite branches X2. This

is termed the secondary spacing while the primary spacing denotes the

distance between the primary arms of growing dendrites.

A2 is an important parameter in cast metals since it determines the

scale of the segregatxon and hence the homogeneity of the material and

the degree of separation of the inclusions. The tip radius P is an

inverse function of the growth rate R, decreasing with an increase of T?.

A number of papers have attempted to correlate X with p. However the most

recent analysis (1) tends to separate the two phenomena of radius of the

tip and instability of the surface behind the tip. The tip is stable in

growth while the branching phenomenon is an instability in growth which

operates by its own mechanism in conditions generated along the surface.

In addition to the instability mechanism, X2 may be dependent on

dissolution mechanisms, subsequent to growth. These are not dealt with

here.

A MODEL FOR BRANCHING INSTABILITY

The problem becomes one of determining an appropriate model for the

spacing between the branches. One such model proposed by Adams (2) is

shown in Fig.2. The arms are taken as cylindrical and grow radially

into a liquid characterised by a solute distribution which was deter-

mined by experimental methods. Adams determining that there was a

constant ratio of solute between that in the liquid at the surface of

the branch Cγ and that in the liquid at the centre between branches Co.

He determined X by heat flow calculations.

In the model which we have adopted, the value of CjjCo is also taken as

constant but we refer this to a physical requirement of the system which

is the optimum spacing of arms for the most rapid dissipation of the

36

undercooling. We suggest that the spacing is such that the initialvalue of solute concentration at the centre between dendrite arms doesnot vary by more than 17C from the solute concentration of the bulkliquid, and this occurs at a sharp minimum. If there were no sharpminimum, the arms would be too far spaced for thft maximum dissipation ofthe undercooling. If the value of solute concentration at the centrevaries too much from Co, the arms are too closely spaced and their rate ofradial thickening would be reduced. We can then calculate X which istwice the thickness of the boundary solute layer for a ratio of CL/COtaken to be 1.005. We have calculated the values of dendritic growthvelocity R for steel cast on copper chills of varying thickness, usinga computer programme and then obtained ^2- We have compared thesecalculated values with values experimentally obtained on cast steelplates. The calculated values were obtained from the following rela-tionship resulting from the model:-

A = . l n kR ln l-ko °

The experimentally measured values of X are in good agreement with thepredicted values. In the equation, ko is the distribution coefficient(for carbon in y iron) and D is the diffusion coefficient (for carbon inliquid iron).

EVALUATION OF X

The evaluation of X is made as follows:

The distribution of solute ahead of a growing dendrite arm is given by:

_ ElCL = Co [1 + £±2 exp

D ]o

Co is the original concentration of solute in the liquid. We thenevaluate the distance L at which C^ = 1.005 C o and suggest that this isthe limiting initial concentration between dendrite arms. This wasconfirmed by experiment.

We then have:

D , 0.005 .L = " ¥ ln T^- ' k°

and for ko = 0.45, X2 = 2L = 11 . D/R.

EVALUATION OF R

The solidification rate, or rate of growth, R was determined from thetemperature distribution in the cast plates. This was considered as aone-dimensional problem with all the heat being extracted in thedirection of the chill. The calculations give at each instant of time,the progress of the solidification front. The influence of the chillwas included in the programme by simultaneously calculating the

37

temperature field in the chill itself. Thus the different thicknessesof chill resulted in different values of R.

COMPARISON OF X (CALCULATED) WITH X (EXPERIMENTAL)

Figure: 3 shows the comparison between calculated and experimentalvalues for different plate and chill considerations. This correlationis good for the 75 mm thick cast plate but shows some deviation for the50 ram plate, it is suggested due to departure from the boundary condi-tions as the system becomes smaller in dimensions. Figure 4 shows thevariation of X with plate and chill thickness.

REFERENCES

1. S.C. Huang, M.E. Glicksman. Acta Met. 29_, 1981, 717.

2. P.K. Rohatgi, C M . Adams. Trans AIME 239, 1967, 1729.

0.05

0.01

10 20 30 40(mm) 10 20 30 40 50 60(mm)

Distance from chill - — »

Fig 3

X Spacing vs. Distance from ChillCalculatedMeasured

A, (mm)

t0.35

0.30

0.25

0.20

0.15

0.10

0.05

0.02

50mm thick plate

» 50mm chill. x 100mm chill

wafer cooled

75mm thick plate

10 20 30 40 50 10 20 30 40 50 60 70Fig 4 Distance from chill ( mm ) — »

K Spacing vs. Distance from Chill

J LFig 1

Characteristicparameters ofdendrite p andX2

Fig 2Model of Rohatgi andAdams (2)

39

TRANSITION FROM FIBROUS TO LAMELLAR MORPHOLOGYIN UNIDIRECTIONAL SOLIDIFIED Ni-W EUTECTIC

S.F Dirnfeld and Y. ZutaDepartment of Materials Engineering

Technion, Israel Institute of TechnologyHaifa 32000 Israel

INTRODUCTION

Observations of transition from fibrous to lamellar structure in uni-directionally solidified (UDS) eutectics were reported. In a Ni-Weutectic by Kurz and Lux!, in an Au-Cu eutectic by Livingstone2 and inAl-A^Ca eutectic by Street et al.3. This phenomenon was reported as anirregular one2, and partially explained by Kurz and Luxl on the basis ofchanges in interfacial free energy. Hunt and Chilton^ attributed thetransition to certain given percentage of inclusion elements which forcethe lamellar to change their orientation and thus change the interfacialenergy per unit area of interphase boundaries. Their explanation doesnot show any relationship with growth rate-R. Jackson and Hunt5 gave amathematical criterion for predicting the preferred structure, thesystem with a volume fraction of the strengthening phase greater than32 pet yields a lamellar structure, while at lower volume structure afibrous structure tends to be generated.

The purpose of this paper is to analyze the dependence of the morphologyof UDS eutectic alloys as a function of inter-fibrous or inter-lamellarspacing and to explain the transition from fibrous structure to alamellar one with increasing the solidification rate. Experimental datafrom the Ni-W eutectic are presented.

EXPERIMENTAL

Specimens of Ni-W eutectic alloy were solidified unidirectionally by aspecial apparatus" where the rate of solidification - R was controlledby the velocity of the furnace and a temperature gradient - G of 10.5°°C/mm at the solid-liquid interface was created. The as grown compositestructures of UDS Ni-W eutectic at relatively slow R (2 to 10 mm/hr)consists of W fibres in a Ni-W solid solution matrix. The strengtheningW phase has a volume fraction of about 6 pet. Fig.l shows the trans-verse section of a UDS specimen growth at 8 mm/hr (G/R = 1.312), thestructure consists of parallel fibers having a uniform distribution. Byfaster R, transition from fibres to lamellar has been observed, thefaceted fibres are transformed into lamellae and two kinds of morphologyexist simultaneously. At a growth rate R = 16.8 mm/hr (G/R = 0.628).most of the structure are lamellar, Fig.2 shows the transverse micro-structure.

ANALYSIS OF MORPHOLOGY

Now the relationship between the volume fraction of both phases and thegeometry of either the lamellar or fibrous structures, will be discussed.Let us consider a cube of unit edge dimension (unit volume) within which

40

Fig.l. Fibrous structure,transverse section

40 #

Lamellar structuretransverse section

the second α-phase appears either as fibers or lamellae in the γ-phasematrix, Fig.3. The number of existing fibers in the unit volume is

nf = 1/X^f, where Xf is the inter-fibrous spacing. The number oflamellae should otherwise ben£ = 1/X^, where X& is the inter-lamellar spacing. The volume ofthe α-phase in the unit volumecube is equal to the volume frac-tion F of this phase. Dividing itby the specific number of fibersor lamellae, we have the averagevolume of a fiber or lamella, thus:Vol. of a fiber = F/n

£ = FX

2

f (1);

Vol. of a lamella = F/n£ = FX& (2).Therefore: df = 2v T/TT'Xf (3)and d& = FXo (4)

1

Fig.3. Transverse cross-section ofunit volume containing fibers orlamellae.

Disregarding the fiber end interfaces or the lamella end surfaces, weobtain the specific lamella-matrix boundary area - Sf:

= 2/A0 (5) S

£ = n

f 7 2/X

f- /irF (6)

It is seen that S« depends only on the interlamellar spacing, and Sf on

the interfibrous spacing and the volume fraction F. There exists acertain volume fraction of the α-phase F

e, for which the specific inter-

facial area for both morphologies is the same. By SJJ, = Sf, we have

( 7 )e h *

and when Xf = X^ the volume fraction of the α-phase will be Fe

or approximately 32%.0.318

41

Assuming that the crystallographic orientation relationship is the same

between the phases in both morphologies, the lamellar and fibrous one,

we have the following expression for the energies:

Es = a S^ (8) ; E

s = a S

f + A(r) (9)

where Es£ and ESf stand for the interfacial energies for the lamellar and

fibrous structures respectively, a is the surface tension for a planar

interface and A(r) the surface tension due to the curvature of the fiber

boundary. The latter would not exist in lamellar interfaces, but in

fibrous structure the A(r) term will increase with decrease of the fiber

radius. This can be expressed by:

O Sf 1

E = OS + — i - = SO(1 + ±-) (10)

Sf t r

£ 1 r

f

where r is the fiber radius.

Substituting eq. (5) and (6) for Esj,j Egf respectively, we obtain

2a

Af

E. =£L (11)

. E = 2

Substituting the value for r_ from eq. (3) we get:

Eq - £ ( 1 3 ) ; E = 2 |

S£ h

Sf X

f \\

The expression for Egf compared with the expression of Eg^ shows that bydecreasing the interfibrous spacing Aj (increasing the solidificationsrate-R) increase considerably Eg

f of the fiber morphology (the second

term of eq. (14) is inversely proportional to x£). However, increase ofthe volume fraction-F of the α-phase will not affect the interfacialenergy of the lamellar structure, but will affect it in the fibrous one.

During eutectic solidification, separation into two solid phases takesplace, each phase with its own concentration of the two components. Inorder to obtain it, diffusion in the melt must take place to build upthe necessary distribution. In this section, the energies needed for itin different morphologies - lamellar or fibrous - will be evaluated.Therefore, we look for some diffusion parameters that will depend on thegeometrical features only, such that will show the differences in energyneeded in the different structures. We return to the unit cube model ofthe material, used in the preceding section, and observe the diffusionfield around a lamella or a fiber. The diffusion path equals the inter-lamellar or interfibrous spacing. Let us put the lamells or fiber atthe midpoint of such a spacing, which represents the width of thediffusion field. In the case of a fibrous structure the width of thediffusion field is the same in both orthogonal directions around a fiberhaving a square form. In the lamellar case the diffusion field lengthin the direction normal to the spacing was taken so that the area of thefield equals that of the fibrous case, see Figs. 4a and b.

In the lamellar case, since the length normal to the spacing direction

was A|/\£, the area of the a phase is {\y\i)A%. In the fibrous case

42

r - * 11

fa.r

\\V

the area of the a-phase is irdf/4. Sub-stituting the valuesof d£ and df fromequations (3) and (4),we get the following:

2

% ; -r^=FXf (15)

Fig.4. The diffusion field a) around a lamellab) around a fiber, c) equivalent arounda fiber.

Since the areas ofthe diffusion-fieldconfigurations areequal in both cases,that of the matrixphase-y is also thesame. In the lamel-lar case the diffusiontook place only inthe spacing-widthdirection (in themelt in front of thesolidification inter-face) so that the

representative geometrical parameters of the process are the lamellarhalf thickness d£/2 and the diffusion field half dimension \%/2 (seeFig.4a). The situation around a fiber is different, and the diffusion-field configuration is more conveniently treated in terms of a circle,with the same diffusion distances around the fiber in each direction,provided this circular configuration has the same area as the square.Thus, the diameter of the circular diffusion-field will become 2Xf//if,and the representative geometrical parameters will be df/2 - the fiberhalf-diameter, and Xf/i/iF - the diffusion field dimension. Since theamount of both phases in the equivalent diffusion-field for bothmorphologies is the same, it is hypothesized that the total energy neededfor the diffusion is proportional to the average diffusion distance forboth components A and B for build-up of the concentration distributionsin the solid state. A mechanistic approach is suggested, whereby theaverage diffusion distances for the components are X^ and XJJ. Thesedistances are between the "centers of gravity" for the concentrationfonfigurations (within the diffusion-field element) for the melt and forthe two-phase solid states. This approach accounts for the concentrationdifferences as well as for the geometrical configurations. In order tocalculate the "center of gravity" points, the static rule of mements willbe considered. It states that the distance of that point of a given areafrom another chosen arbitrary point equals the sum of moments of thepartial areas about that same chosen point, divided by the sum of thepartial areas. Let the following designations be used:I/j. and L2 are the width of the phases a and y respectively; C^ and C2 -weight concentrations of component A in phases a and y respectively and(1-Ci), (I-C2) - the same for component B in both phases (Fig.5). Thus,be the rule of moments, one gets:

43

Cβ

1

1-C,

Co

1-C,

——

- L , -

T T

(a) (y> (a) (y)

Fig.5. Concentration profile

of components A and B.

XA =

and

L l + L 2 L l C l / 2

+ L 2 C 2 ( L l + L 2 / 2 )

L 2 C 2 + L 1 C 1

- L

l( 1-

CJ/2

+ L2

( 1-

C2

) ( Ll

+ L2/2

) Ll

+ L2

"TJ T f "i p ^ _LT ("\ C \ O2 2 1 1

(17)

for the lamellar structure Li=d£/2 and

L2=X9-d!

;/2 see Fig.5. If we use the d£

value from eq.(4), then Li=FX£/2 and

L2=(l-F)X£/2, therefore we get the

average diffusion distances:

and

X£ (F -F)(C

2-C

1)

(18)

(19)

However, for the fibrous structure L =dJ2 and L =Xf//rT - j— ; using d

f

from eq.(3), one gets:

= Xf/F (!-• = V^TTT • X, (20)

therefore:

XAf =

2/rr

(21) ; xB f =2/F

(22)

Now, assuming that the total energy for diffusion within a diffusion

field element is proportional to the average diffusion distance, the

analysis for diffusion in both morphologies are given by:

( 2 3 ) EDf -

Where: PA and Pg are some energy functions incorporating the diffusion

coefficients and concentration parameters for components A and B, and n

is the number of diffusion field elements within the actual cross section

of the specimen; E ^ and EDj are the total required energies for

diffusion for the lamellar and fibrous morphologies respectively. The

source of these energies in the system is the amount of supercooling

below the solidification temperature^. It should be recalled that the

system itself will choose the morphology, lamellar or fibrous, which

needs the lower amount of supercooling. Let us denote E£ for the energy

needed to obtain lamellar structure and E£ Tor the fibrous one:

+ E,.c (26)=

E

si(25) ; E

f =

44

Substituting eqs. (13) and (14), we get:

A (F2-F)(C -C ) X (F*-F)(C -C )

X. (F-/F)(C_-C ) A (F-/F)(C -C )E 20 /iJ + 2TO + n{ j_E 2_J^ + p r_J L _ ^£ Af X^ A 2/ir v¥(C2-C1)+C2

B 2/irf X 2 1 2 y ^ ^

(28)CONCLUSIONS AND DISCUSSIONS

From the above expressions (27) and (28) the following conclusions can bedrawn

i) Increase of the interface spacing in both morphologiesreduces the interfacial surface energies.

ii) The fibrous interfacial surface energy, compared with thelamellar one, increases considerably as the interfibrousspacing X_ increases.

iii) Decrease of the interphase spacing in both morphologiesreduces the average diffusion distances, and thereby theenergy needed for diffusion.

Following these conclusions, the transition from one morphology toanother by changing the rate of solidification can be explained asfollows:

When an eutectic system solidifies, the rate of solidification dictatesthe maximum available average diffusion distance. A lower rate willincrease the average diffusion distances and vice versa. Eutecticsystems will seek to reduce the surface energies in both morphologies.In the fibrous structure the surface energy is reduced also by increas-ing the curvature (eq.lO), since the number of fibers decreases withincrease of the interface spacing and therefore their radii increasefor a constant volume fraction. To sum up, increase of the interphasespacing, A. or A-, reduces the surface energy terms in the totalexpressions; on the other hand, it increases the energy needed fordiffusion, and the two increases may cancel each other out, in whichcase decrease of the solidification rate would not increase the inter-face spacings. This accoun ~ for the upper bound on these spacings(and thereby on the thickness of the lamellae or fibers) existing indifferent eutectic systems. With a given rate of solidification, whichdictates the maximum average diffusion distances, the structure callingfor lower energy expansion will be preferable. Increase or decrease ofthe solidification rate may bring about a situation favoring the secondmorphology, and transition to that structure will then take place.

It seems that at relatively low solidification rates the total energyfor the fibrous structure is lower than that for the lamellar structure.But when the solidification rate increases the dimension of the fibresand X decrease and the curvature component of the surface tensionincreases until the total energy needed for the fibrous structureexceeds that for the lamellar, and at that point the latter is favored.

45

This finding is supported by the micrographs (R=14 mm/hr) which show astructure composed of both fibers and lamellar, since at that solidifi-cation rate the total energy is the same for both morphologies. Thefibers have square or rectangular cross-sections; the lamellae exceedthe fibers through joining of several fibers end to end. Certainlamellae are seen to havt an "appendage" at the end, as though a fiber isattached there (Fig.2).

The model shown in this section is applicable to other systems, and mayaccount for transition phenomena from a fibrous to a lamellar structurereported as unexplained irregular effects by Livingstone^ in the Au-Cusystem and by Street^ in the Al-Al^ C a system.

REFERENCES

1. W. Kurz and B. Lux, Met. Trans, l^ (1971), p.329.

2. J.D. Livingstone, J. Appl. Phys. 41 (1970), p.192.

3. K.N. Street, F.C. St. John and G. Piatti, J. Inst. Met., f5 (1967)p.326.

4. J.D. Hunt and J.P. Chilton, J. Inst, of Metals, 91 (1963), p.338.

5. K.A. Jackson and J.D. Hunt, Trans. AIME, 236 (1966), p.1129.

6. S.F. Dirnfeld and D. Schwam, Proc. 2nd Ristf International Conf. onMetallurgy and Materials Science, Roskilde, Denmark, 1981, p.225.

46

THE MICROSTRUCTURE OF RAPIDLY SOLIDIFIED Fe-Cu ALLOYS

A. Munitz and Z. Livne

Nuclear Research Center-Negev P.O.B. 9001, Beer-Sheva, Israel.

ABSTRACT

rhe impact of cooling rate on the microstructure of Fe-Cu alloys wasinvestigated. A variety of solidification techniques v;ere employed,in orderto achieve cooling rates of several order of magnitude variation. Ourexperimental results Teveal an extension of copper solid solubility iny iron. The cooling rates have a drastic influence on the microstructuTe.\t cooling rates of about 10* °C/s a dendritic solidification mode couldbe observed, while at high cooling rates the iron has g. circular appearanceembeded in a copper rich matrix. The similarity of the microstructureobtained by electron beam surface melting to that obtained in molten pocketsby explosive bonding of Fe-Cu plates, suggests a similarity in the coolingrates (i.e. 105 °C/s) (1).

INTRODUCTION

Explosive bonding of metals is a method for producing large area bondingbetween metal plates, particularly when the metal mechnical propertiesdiffer grossly. A well known example is the bonding between large platesof lead and steel, which can not be achieved by any other means. A typicalindustrial application of explosive bonding is the welding of pipelinesfor heat exchangers. In such systems the pipeline has to be corrosiveresisting inside, and highly thermally conducting for the rest of itsthickness.

Steel - Cu explosive-bonded materials are used for protection walls ofsafes. The outer layer is a very hard thin steel layer aimed to preventmechanical penetration. The inner layer consists of a thick copper platefor heat absorption aimed to prevent penetration with a torch.

Explosive bonding involves very fast solidification rates [1). In ourpresent work we report a metallurgical study of fast solidification ratesin Fe-Cu alloys involving explosive bonding, as well as electron beamsurface melting and melt quenching on a wateT-cooled copper plate. Withthese three techniques it was possible to scan over a large range ofsolidification rates.

EXPERIMENTAL

In the present work we have studied specimens of Fe-Cu alloys obtained bythe following methods: A) Resolidification of melt on a water-cooled copperplate: about 10 gr of Fe 50 w/o Cu bulks were melted in a ceramic cruciblein an induction furance. After melting, the material was poured over awater-cooled copper plate. The cooling rates evaluated on the basis ofthe cast material thickness (2) range between 10 and 100 °C/s. B) Electronbeam surface melting and subsequent solidification by self quenching.The melting conditions were as follows: driving voltage 60 kV, current

47

between 5 and 10 mA, electron beam scanning velocity about 500 cm/min.The electron beam focus was about 1 mm beneath the molten surface. Coolingrates as high as 105 't/s were achieved (3,4^. C) Resolidification of moltenpockets in an explosive bonded Fe-Cu plates. A 5 mm thick copper plate ofdimensions 200 mm x 300 mm was used as the flyer plate, and a 10 mm thickiron plate of the same dimensions was used as a parent plate. The flyerplate was placed 1.5 mm parallel to and above the parent plate, the latterwas lying on the ground. The explosive material used T=ras TNT+NH4N03 givingrise to a 2200 m/s detonation velocity.

Selected areas from the above mentioned three solidification techniqueswere cut with a slow diamond wheel. The surface was then mechanicallypolished with diamond pastes up to 1/4 vm, then etched v-th 2gr FeCl3+10 ccHC1 in 95 cc ethanol (Cu etchant) for 10 s. The specimens were examinedwith a scanning electrom microscope (Philips SEM 505) with EDAX options.

RESULTS AND DISCUSSION

In Fig. 1 we show secondary electron images (SEMs) of Fe - 50 w/o Cuquenched on a water cooled copper plate. The figure reveals a dendriticsolidification structure, consisting of primary arms of about 100 pm longand 10 vim wide. EDAX analysis showed, that between 9 to 13 w/o Cu ismiscible in the dendrite arm (see Fig. 1). Observation of Fe-Cu equilibriumphase diagram (5) indicates an upper limit of about 8% miscibility ofcopper in y-Fe phase. In our case first to solidify are primary y-Fedendrites, which dissolve up to 13 w/o copper. The relatively high coolingrates involved (10-100 °C/s) allow this enhanced miscibility of Cu iny-Fe phase. The y-Fe phase is unstable at room temperature. It transforminto a-Fe phase. Indeed no y-Fe phase could be detected by x-ray diffraction.All the sample volume contained a-Fe and e-Cu phases solely. Cu is immisiblein a-Fe. It then precipitates in the dendrite arm as the primary y-Fecrystals transform into the a form. Upon etching with nital, a fine micro-structure is revealed in the primary y-Fe crystals (see Fig. IB). Thedendrite arms are embeded in a copper-rich matrix (68 w/o Cu). It consistsof a primary e-Cu, as well as a-Fe secondary phases. Precipitated copperis expected to exist within the a-Fe, but this detail could not be seenin the present resolusion.

in Fig. 2 we show a secondary electron image of a resolidified surfacemelted by an electron beam. The copper etching of the specimens revealstwo types of structure: A) Big particles (>10 pm) with flat surfaces,within which fine rounded second phase precipitants may be observed.Elemental x-ray micro-analysis of these particles made by the scanningelectron microscope showed, that about 30 to 35 w/o Cu is present. On theother hand, when x-ray diffraction was performed on electron beam treatedsurfaces, only a-Fe phase could be seen. We atribute the high copper con-centration in these particles to the impact of the high cooling rate onthe Cu miscibility in y-Fe. In this case, cooling rates as high as 105 °C/sare effected, which in turn enhance the Cu miscibility up to 35 w/o.Presumably, the big particles first solidify as primary y-Fe crystals.They are then vigorously swept by the convections in the molten pool.When solidification continues the y phase transforms into an ct-Fe phase,and this process is accompanied by Cu segregation. Compared to the previo-usly described solidification of Fe-Cu alloy quenched on copper (Fig. 1),only short cooling times are involved. Therefore, a different type of Cuprecipitation could be observed; (B) Small tiny spheres with diameters

48

between 0.2 vm and 2 ym embeded in a copper-rich matrix.

In Fig. 3 we show a secondary electron image of Fe-Cu alloy resolidifiedas molten pockets by explosive bonding of Fe-Cu plates. Upon proper Cuetching the sample reveals two types of structure: (I) Big flat particles,with surfaces which appear smooth even at high magnifictions. X-ray micro-analysis of this structure showed only 0.9 w/o Cu. These results suggestthat the particles did not melt al all. Apparently, the particles weretorn away from the iron plate by the large melt convection, a result ofthe large turbulent flow induced by the high pressure involved in theprocess. They did not have sufficient time to melt; (II) Tiny round spheres,about 0.1 vm diameter, embeded in a Cu rich phase, which was disolved bythe etchant. The presence of tiny spheres attached to the big flat particleindicates, that the particle envelope has melted, and resolidified.

The similarity between the microstructure of Figs. 2 and 3 suggests, thatcooling rates as high as 10s °C/s could be obtained during resolidificationof molten pockets, consistent with theoretical predictions (1).

SUMMARIZING REMARKS

The microstructure of Fe-Cu alloys were investigated. It was found thatthe major impact of high cooling rates is to enhance copper miscibilityin Fe. The primary Y-Fe transforms into a-Fe, and induces copper precipit-ation. The character of the a-Fe, which transforms from y-Fe particlesis different in the different solidification processes. Microstructuralsimilarities between electron beam molten surfaces and molten pockets inexplosive bonded plates suggest, that cooling rates as high as 105 °C/sare obtained in both cases.

ACKNOWLEDGMENT

The authors wish to thanks Mr. C. Cotler and Z. Barkai for their technicalaid, and to Dr. Z. Burshtein for his editing remarks. Finally, the authorsthanks Prof. B. Weiss for his helpful discussions.

REFERENCES

1) Z. Livne, Structure phenomena in the bond zone of explosively bondedplates, Nuclear Research Center-Negev, NRCN 464, 1979.

2) F.A. Joly and R. Mehrabian, J. Mater. Sci. 9_ , 1446 (1974].

3) S. Kou, S.C. Hsu, and R. Mehrabian. Metall. Trans. 12B , 33 (1981).

4) R. Mehrabian. Int. Metals Rev. 27 , 185 (1982).

5) P.M. Hansen, Constitution of Binary Alloys, 2nd Ed. (McGraw Hill, 1958).

49

Secondary electron images (SEMs) of Fe-50 w/o Cu qunenched onwater cooled copper plate. (A) General view. A typical x-rayspectrum of a dendrite arm is given in the right side; (B) Across-section of dendrite arm after etching with nital „

50

Fig. 2 : Secondary electron image of an electron beam molten surface.An x-ray spectrum of the big particle is given on the right side.

Fig. 3 : Secondary electron image of resolidified molten pocket in explosivebonded Fe-Cu plates. An x-ray spectrum of particle (type I) isgiven on the right side.

51

PHASE STABILITY AND MASSIVE TRANSFORMATION IN

Y 0 -COMPLETELY STABILIZED ZIRCONIA (Y-CSZ)+

A.H. Heuer, R. Chaim* and M. Ruhle**Dept. of Metallurgy and Materials Eng.,

Case Western Reserve Univ., Cleveland, U.S.A.

INTRODUCTIONZirconia (ZrO2) alloyed with Y2O3 stabilizer has recently become animportant material due to the high strength and fracture toughness rela-tive to other stabilized oxide-zirconia systems. The improvement of thesemechanical properties is directly connected to the volume fractions ofZrO2 polymorphs, namely cubic (c), tetragonal (t) and monoclinic (m)phases, within the material microstructure. The presence of high concen-trations of oxygen vacancies needed for electrical charge neutrality andthe sluggish diffusion of the stabilizer cation causes instability of thehigh-temperature polymorphs (c.,_t) under certain conditions. Thusdifferent phase-diagrams for the _t + c_ two-phase field have been reported(1).

EXPERIMENTAL

As-received mixed powders of ZrO2 - 12 wt% Y2O3 were pressed and sinteredat 1600°C for 2 hr. Small specimens were subsequently heat-treated asfollows: 1550°C/l hr; 1400°C/1,5 hr, 1 week; 1250°C/20 hr. These werechecked by X-ray diffractometry using Cu-Ko. radiation and later preparedfor transmission electron microscopy (TEM), by mechanical and ion thinningfollowed by carbon coating. The alectron microscope (JEOL 200-CX) wasoperated at 200 KV.

RESULTS AND DISCUSSIONTEM observation of the sintered specimens revealed c grains with a tweedstructure containing small _t precipitates. Careful observation of thesegrains in very thin areas showed fine modulations in two directions(Fig.l) due to the strain field around the coherent t_ precipitates.Selected area diffraction patterns (SADP) from such £ grains for [lll]c

zone-axes showed the 3 variants of [112]t reflections (Fig.2) indicatingthe presence of conventional primitive t_ structure. A completely differ-ent microstructure was observed for 1550°C heat-treatment correspondingto new peaks in the X-r:-iy profile (Fig.3), the c_ grains have converted totwinned and acicular crystals. The internal microstructure of thesecrystals was characterized by curved antiphase boundaries (APB). Howeverobservation of appropriate diffraction patterns (Fig.5) showed thepresence of only 2 variants of [112]t reflections indicating a differentt_ symmetry, namely body centered tetragonal (BCT), t_'. This phase wasreported recently (2) as a product of quenching 8 wt% Y2O3 PSZ in theskull melting process. The same microstructure was found for heat-treatments at 1400°C for 1 and 5 hrs. where several tweed c grains were

+ To be published in Advances in Ceramics, 1983* Dept. of Materials Engineering, Technion, Haifa, Israel.** Max-Planck Institut fur MetalIforschung, Stuttgart 1, F.R.G.

52

also observed. EXAX microanalysis showed an 13 wt% Y2O3 content in cgrains, where 12 wt% Y2O3 was found for the _t' phase.The rapid phase transition of c to _t' at constant composition suggests amassive transformation. According to the ZrO2~Y2O3 phase diagram (3), atthe above temperatures a drastic composition change from 12 wt% to 3-4 wt%Y2O3 should occur to form the conventional _t precipitates. Thus the new_t* phase could form without composition change (a displacive transforma-tion) at short annealing times, but must necessarily change crystalsymmetry.Indeed this metastable form, which was stable relative to the t_ -> m trans-formation, disappeared with further annealing for 120 hrs. In addition,a new microstructure of conventional _t phase then appeared, containinglens-shaped to colonies which were internally twinned (Fig.6). This t_phase was not stable relative to the t_ •> m transformation. At 1250°C thesame tweed c structure was found, confirming the sluggish nature of dif-fusion process in this material.

ACKNOWLEDGEMENTS

R. Chaim acknowledges the MPI association support in liis stay in Stuttgart.

REFERENCES

1. A.H. Heuer, N. Claussen and M. Ruhle, Advances in Ceramics, 1983, tobe published.

2. V. Lanteri and A.H. Heuer, Advances in Ceramics, 1983, to be published.3. H.G. Scott, J. Mater. Sci., 10, 1527-1535 (1975).

Fig. 1: £ grain with tweed structure insintered material.

53

Fig. 2: [111] zone-axis diffraction fromc_ grains.

Fig. 3: {400} peaks for £,_t»£' phases

54

Fig. 4: APB's in the acicular _t*crystals.

Fig. 5: [111] zone-axis diffractionfrom t1 phase.

Fig. 6: _t lens-shaped colonies,internally twinned.

55

MICROSTRUCTURE EVOLUTION AND ORDERING IN COMMERCIAL MGO-PARTIALLY

STABILIZED ZIRCONIA (MG-PSZ)*

R. Chaim and D.G. Brandon


Technion-Israel Institute of Technology, Haifa, Israel

INTRODUCTION

The polymorphic nature of ZrO2 on cooling from cubic (c_) to tetragonal (_t)

and finally monoclinic (m) structures is now well understood. For

technological applications ZrO2 is usually "alloyed" with the cubic

oxides, for example MgO, CaO, Y2O3, to stabilize the high temperature

polymorphs (£,_t) at ambient temperature. The jt -*• m stress induced

martensitic transformation is accompanied by a 3% volume increase, and is

responsible for the fracture toughness improvement. Thus the mechanical

properties of PSZ are sensitively dependent on the material micro-

structure, a fact which explains the comprehensive microstructural

studies (1-6).

EXPERIMENTAL

Sintered commercial 9 mole % MgO-ZrO2 (PSZ) specimens were heat-treated

under the following conditions: 900°C/20 hr, 1000°C/5 hr, 1100°C/l hr,

1400°C/120 hr. The specimens for transmission electron microscopy (TEM)

were prepared by mechanical and ion thinning methods, followed by carbon

coating. The electron microscope (JEOL CX-100) was operated at 100 KV.


The as-received material was composed of c grains with small ellipsoidal

shaped _t precipitates, dispersed homogeneously in the £ matrix (Fig.I),

Other £ grains contained larger lenticular shaped _t precipitates (Fig.2),

which were transformed to the m symmetry and characterized by internal

twinning. Large m twins were observed as a grain boundary phase (Fig.3),

with a glassy phase at grain boundary triple points. Selected area

diffraction patterns were used to identify the different phases mentioned

above. The £ grains containing small t_ precipitates have related diffuse

scattering intensity (DSI) in diffraction patterns, which has proved to

be associated with the £ matrix (7).

The microstructure after 1400°C annealing was changed completely to £

grains containing m particles in an eutectoid microstructure with an

MgO-rich phase. No changes in the volume fraction of the different

phases was observed after annealing at 900-1100°C, except for appreciable

strain-field contrast around the small £ precipitates. Selected area

diffraction from such areas revealed an ordering phenomena characterized

by strong superlattice reflections (Fig.4). The ordered phase was

identified as β-phase (Mg2ZRsOi2) with rhombohedral structure, located in

the interspace between the small _t particles (Fig.5).

To be published in the Jour. Mater. Sci.

56

This rhombohedral structure is a distorted version of the c^ fluorite

structure, with strings of oxygen vacancies located along a specific

<111>C direction. The β-phase has been reported (8) as a high temperature

phase formed by quenching from 1850°C. Thus the formation of this phase

at lower temperatures has been considered a result of non-equilibrium.

The DSI in PSZ was related previously (7) to short-range ordering of

oxygen vacancies, which produced trigonal relaxation of the c^ fluorite

structure. The slow diffusion of stabilizer cation enables only short-

range diffusion, which causes local enrichment of c_ regions between small

_t particles. The appropriate high saturation concentration enables the

β-phase formation (28.6 mole % MgO), which was assisted by the presence

of ordered oxygen vacancies. This microstructure is metastable and for

long annealing times dissociation should occur (Fig.6), to reach bulk

equilibrium.

The formation of β-phase enhances the stabilization of _t particles by

retarding further t_ growth, but destabilizes it by increasing the

coherency strain field at the t-0 boundaries.

ACKNOWLEDGEMENTS

This work was in part performed at MPI fiir Metallforschung, Stuttgart.

REFERENCES

1. G.K. Bansal and A.H. Heuer, J. Amer. Ceram. F o e , j>B (5-6),

235-238 (1975).

2. R.H.J. Hannink and R.C. Garvie, J. Mater. Sci., Y]_, 2637-2643 (1982).

3. D.L. Porter and A.H. Heuer, J. Amer. Ceram. S o c , 62^, (5-6),

298-305, (1979).

4. L.H. Schoelien, Ph.D. Thesis, Case Western Reserve Univ., Cleveland,

Ohio, USA, 1981.

5. R.H.J. Hannink, J. Mater. Sci., 18, 547-470 (1983).

6. R.H.J. Hannink, J. Mater. Sci., l^, 2487-2496 (1978).

7. R. Chaim and D.G. Brandon, Advances in Ceramics, 1983, to be published.

8. 0. Yovanovitch and C. Delamarre, Mat. Res. Bull., 11, 1005-1010 (1976).

57

Jig. 1; Ellipsoidal _t precipitates in a £

Fig. 2: Lenticular m phase in a £ matrix.

Fig. 3: Large m twins at grain-boundaries.

58

Fie. 4: [lll]c zone-axis diffraction from

ordered β-phase regions.

Fig. 5; D.F. image showing β-phase between

£ particles.

Fig. 6; Dissolution of the ordered β-regions.

59

DEFORMATION INDUCED DECOMPOSITION OH URANIUM-TITANIUM MARTENSITE

* * * **G. Kiiranel , J. Sariel , A. Landau and M. Talianker

* Nuclear Research Center Negev

** Ben-Gurion University of the Negev.

Uranium and titanium form a solid solution at high temperatures at allconcentrations. The crystal structure of this solution is A2 tungstentype as that of the high temperature phases of the pure elements (y uramiumor 3 titanium). The uranium-rich alloys (below 25 at % titanium) decomposeinto a uranium and S-^Ti when cooling from the y solid solution, but rapidcooking of y phase by water quenching results in a metastable a' phase,which is a supersaturated solution of titanium in a uranium. The firstreaction y -*• a+S is an eutectoidic decomposition and the second is amartensitic transformation γ-â

1.

The a1 martensite may transform into a+S during heat treatment above 300 C.

The S phase has an A1B2 crystal structure, which is similar to the to phase

of zirconium or titanium alloys. The formation of u phase from the bccstructure (H-T phases of zirconium and titanium) is v.'ell established andits analogy with the formation of 6 from y uranium has been indicated (1) .

The discovery of the diffusionless transition a1 ->• m in zirconium under

high hydrostatic pressure (2) led us to examine the stability of a1

martensite under stresses. We found that plastic deformation of a1 is

always accompanied by the a' -> a transition.

Geometric analysis shows that the transition a1 -»• 6 is similar to a

1 -*• to

in zirconium, so that the diffusionless transformation a1 *- a+S takes

probably place during plastic deformation of (U-Ti) martensite.

EXPERIMENTAL

Uranium 5% Ti was cast into graphite molds at high vacuum. The cast wasannealed for 24 hours at 850°C after hot working in order to get anhomogeneous titanium distribution. The martensitic structure was obtainedby resolution of the titanium at 800°C and water quenching. Tensilespecimens were machined and the tensile test showed elongation of 25 ± 5%.The specimens, which have been broken in the tensile test were cut for XRDstudies. For comparison, several specimens were subjected to various heattreatments leading to different structures.


The diffraction spectra giving the different lattice unit sizes, whichobtained according to the heat treatment given,are:

a) Slow furance cooling from y phase, giving a mixture of a uranium and6-IJ

2Ti (Figure 1-1).

60

b) Air cooling from y, giving a mixture of a uranium which dissolvestitanium (but unsaturated) with S-U-Ti (Figure l-II).

c) Water quenching giving an a1 martensite, which is suppersaturated solidsolution of titanium in uranium (Figure l-III).

The lattice parameters calculated by least-squares routine are indicatedin the figure.

The diffraction spectrum obtained from the deformed martensite (at theneck of the tensile specimen) is similar to that of figure l-II,but withoutthe 6-phase peaks. It indicates that a transition a' > a occurred duringthe deformation. In spite of the absence of the 6-phase peaks, the excessedtitanium is probably precipitated as 6-phase, giving the complete reactionaf -+ a+tS, since the a phase dissolves no more than 3 at% titanium. As thereaction occurs rapidly (during the tensile test) and at room temperature,it is assumed that it is diffusionless.

A similar transition was found in zirconium. The diffusionless transitiona1 -*• a) in zirconium was explained by a model showing the crystal lographicrelationships between the zirconium a' and u phases-

The similarity of a' uranium to a' zirconium and 6-U2Ti to u phase leadsto a similar model. Figure 2 shows the crystallographic orientationrelationship between a'-uranium and 6-U2Ti:

(loo) [oio]a II(oooi) [iioo]6The lattice distortions between uranium and lUTi are 0.46%, 4.7% and 2.7%along [l00]a , [0lQ]a and [00l]a respectively. The missfit is in the rangewhich plastic deformation can accomodate. Moreover, a comparison of theatomic arrangements in the twolattices reveals that the atomic shuffleson the (100)a plane to form (1100)6 plane are small - 0.13 A°and 0.2 A

G

l [OOl]along [001]^must occur in the directions of ±[l00]titanium in U»Ti.

and [01l]a directions respectively, but shifts of 0.7 A°a to form two layers of uranium and

REFERENCES

1. R.D. Tomlinson, J.M. Silcock and J. Burke, J. Inst. Metals, 1970, 98,154.

2. A. Rabinkin, M. Talianker and 0. Botstein, Acta Met. 1981, 29, 691.

61

Figure 1: Diffraction spectra obtained according to the heat treatment given.

o

1 2

o o

Figure 2 : Atomic positions on (100)(J Cleft), compared with (OOODy T i (right).

t i

63

APPLICATIONS OF ANALYTICAL ELECTRON MICROSCOPY TO MATERIALS

J. I. Goldstein, D. B. Williams and M. R. Notis

Department of Metallurgy and Materials EngineeringLehigh University, Bethlehem, PA 18015 USA

The analytical electron microscope (AEM) incorporates an energydispersive spectrometer (EDS) to detect x-rays from an electron-trans-parent foil. The combination of a thin foil and a focused high energyelectron beam permits a nominal x-ray spatial resolution of ~ 10-50nm. This resolution is two orders of magnitude better than thatobtained with bulk specimens in the scanning electron microscope (SEM) .In addition, quantification of the x-ray data is relatively straight-forward using the ratio method, where the concentration ratio ofelements A and B in the thin foil is directly proportional to the x-ray

intensity ratio:CA / CB " AB <vv (1)

The constant k is independent of concentration but varies with the AEMoperating voltage and may be determined using standards or by calcula-tion from first principles. The AEM instrument is used in numerousmaterials applications. This paper will illustrate a wide variety ofexamples where the instrument is of particular use.

1. Study of Precipitate-Free Zones in Al-Ag (1)The occurrence of precipitate-free zones (PFZ) around grain bound-

aries in aged Al-based alloys has been attributed to solute depletiondue to grain boundary precipitation, or vacancy depletion. In Al-Agalloys two distinct zones free o£ the metastable Y ' precipitate areobserved (Fig. 1): a wide 'grey' PFZ (GPFZ), and a narrow ( ~ 500 nmwide) 'white' PFZ (WPFZ) immediately adjacent to the grain boundary.A typical TEM microstructure of an aged alloy is shown in Fig. 1,

DISTANCE FROM GRAIN BOUNDARY x 100 t&HCt PROM C«*!N SOUNQARV « 2 Q 0 » *

Fig. 1: A typical micro- Fig. 2: (a) AEM concentration profile acrossstructure of a Al-16 wt%Ag sample aged at 433 Kfor 50 hrs.

a grain boundary WPFZ for a sample aged at433 K for 50 hrs. (b) AEM concentrationprofile across a GPFZ in the grain boundaryregion.

64

exhibiting a well-defined WPFZ and GPFZ. Solute concentration profileswere measured using an electron-beam diameter of 5 nm. X-ray data wereanalyzed using equation (1). The results (Fig. 2) indicate that theWPFZ is caused by marked solute depletion (Fig. 2a) and the GPFZ byvacancy depletion (Fig. 2b). The solute content of the GPFZ is equalto that of the bulk. Therefore both solute depletion and vacancydepletion mechanisms explain the formation of PFZ in this system.

2. Early Stage Growth of the Ni Al Intermediate Phase in Ni-NiAl Dif-fusion Couples (2)

The major experimental techniques used during the past decade tostudy the kinetics of interdiffusion and intermediate phase growth havebeen the optical microscope and the electron microprobe. Because bothof these techniques have spatial resolution limits of~1 \M m, it hasnot been possible to examine early growth kinetics. For example thereis little or no information available on the kinetics and/or morphologyof Ni Al layer growth for times < 3 hrs.

Figure 3 is a typical microstructure showing the growth of Ni Alin a Ni-NiAl diffusion couple produced at 1100°C. The protrusions inthe Ni,Al phase are always associated with grain boundaries in theNi Al layers since grain boundary diffusion significantly contributesto Ni Al layer growth below 1100°C.

Fig. 3: TEM micrograph of an Ni_Al layer protrusion associated with asingle Ni_Al grain boundary.

AEM profiles obtained across the diffusion interfaces are shown inFig. 4. To obtain the upper data set in Fig. 4 the data were corrected

Fig. 4: AEM profiles for a specimenheld at 1100°C for 15 minutes.

DISTANCE (/*«)

for x-ray absorption effects (3) . The equilibrium concentrations atthe appropriate two-phase boundaries at 1100°C are shown in Fig. 4.The data obtained at the interfaces appears consistent with theinterface concentrations predicted from phase equilibria. It thusappears that interfacial equilibrium is established even at very shorttimes.

3. Chemical Identification a: Submicron Particles in Steel Weld Metal(4)The complexity of the variables that affect the welding process

make it difficult to determine the relationship between the microstruc-ture and the mechanical properties of welds. Interpretation of thesignificance of microstructural differences has been limited by thelack of localized chemical information. However AEM, in combinationwith optical and scanning microscopy can overcome some of theseproblems.

The identification of submicron particles in pressure vessel steelweldments is of particular interest because of their effect on mechani-cal properties. Fig. 5a shows a TEM image of precipitates in such aweldment. The EDS spectrum of the precipitates, Figure 5b, taken inthe AEM shows large amounts of Si and Mn indicating that the precipi-tates are silicates, responsible for poor impact properties. If suchprecipitates are analyzed in bulk specimens using the SEM the matrix Fewill so dilute the EDS spectrum that silicate identification isimpossible.

(a) (b)Fig. 5:(a) TEM image showing the presence of precipitates character-istic of G80 weldments. (b) EDS spectrum from matrix precipitatesin Fig. 5(a).

4. Low Temperature Diffusivity Measurements (3)Y In Fe-Ni alloys, a knowledge of the diffusivity of Ni in Fe-Ni

(D ) below 800°C is necessary to model the growth of ferrite inaustenite. Experiments that report D values have been carried outonly above 1000°C. Simulation of ferrite growth therefore requires aconsiderable downward extrapolation of high temperature diffusivitydata. Diffusion distances obtained for temperatures < 1000°C and fortimes < one month are smaller than the spatial resolution of the EPMA.Because of the improved resolution of the AEM, diffusion profiles thatare 40 times as small as those necessary for the EPMA technique can bemeasured at temperatures down to 750°C.

66

Fig. 6 shows three profiles in an Fe-10.4 wt% Ni/Fe-15.5 wt% Nidiffusion couple. In Fig. 7 the measured D* values are compared withother high temperature data. The diffusivities agree well with theextrapolated values of the high temperature data.

3 o"i i 96S1 !

'a"*| - «ENe-raaiiRrrn rieret

Fig. 6: Measured Niconcentration profilesacross the weld interface.

Fig. 7: Variation of measured Ni diffusiv-ity with temperature. The open circle ( © )is the measured D ^ in Fe-12.5% Ni-0.15% P,and the filled circles (•) are thosemeasured in Fe-12.5% Ni.

In summary the AEM can be used for a large number of materialsapplications. With improved specimen preparation techniques higherbeam currents and higher operating voltages, we may eventually haveeven better x-ray resolution than is available on our present AEMinstruments.

Acknowledgment sWe wish to acknowledge the support of the National Science Founda-

tion through Grants Nos. DMR 79-23278, 80-23955, EAR 82-12531, theDepartment of Energy through Contract No. EY 76-5-02-2408 and theNational Aeronautics and Space Administration through Grant NGR39-007-043.

References1. Merchant, S. M., Notis, M. R. and Williams, D. B., 1982, Solid-

Solid Phase Transformations, The Metallurgical Soc. A1ME,Warrendale, PA, p. 733.

2. Glitz, R., Notis, M. and Goldstein, J. I., 1982, Solid-SolidPhase Transformations, The Metallurgical Soc. AIME, Warrendale, PA,p. 691.

3. Goldstein, J. I., Costley, J. L., Lorimer, G. W. and Reed, S. J. B.3

1977, SEM/1977 I, 0. Johari, ed. Chicago, p. 315.4. Sankar, J. and Williams, D. B., 1981, SEM/1981/I, SEM, Inc., AMF

0'Hare (Chicago), IL, p. 159.5. Narayan, C. and Goldstein, J. I., 1983, Met. Trans., 14A, 2437.

67

FIELD ELECTRON AND ION EMISSION FROMZIRCONIATED AND Zr FREE W CATHODES

* **J. Pelleg and J.L. Fink

Materials Engineering, Ben-Gurion Universityof the Negev, Beer Sheva, Israel.

**atBel l L a b o r a t o r i e s , Murray H i l l ,

N . J . 07974 , U.S.A.

ABSTRACT

Electron Beam Exposure Systems (EBES) for patterning future submicrondevices have been one of the major fields uf investigation in currentyears [1]. The successful operation of EBES cathodes is not fully under-stood. The cathodes are basically single crystal emitters having a <100>orientation, and they are processed before use. The term given to thisprocessing is activation [2], and it basically consists of heating theemitter in an oxygen ambient, after Zr has been deposited on it bysome desirable technique. The purpose of activation is the lowering ofthe work function (WF) of preferential planes of the cathode by theads orb ant, in this case Zr T3-8], to achieve strong emission from theseplanes during uninterrupted use for thousands of hours. These are {100}planes. However, reports indicate [6,9] that other planes also areassociated with the lowering of the work function and thus with brightemission.

In this work the emission characteristics of zirconiated W tips isreexamined. Field electron emission microscopy (FEM) is supplementedby field ion microscopy (FIM).

EXPERIMENTAL

The system could be operated in FIM or FEM mode due to the reversepolarity of the power supply.Since the fields necessary for ionizationof the imaging gas is about an order of magnitude larger than the nega-tive fields required for electron emission, it was essential to work withion emitters having radii in the 500-1000 A range. W single crystalemitters having <100> orientation were prepared by electropolishing.

The Zr was evaporated in situ onto the W cathodes. Before the evaporation,the Zr source assembly was outgassed for an extended period and inaddition it has been kept always at a sufficiently high temperatureduring the stand-by period.

Evacuation of the microscope was performed by a combination of a diffu-sion pump, Ti sublimation pump and ion pump. The diffusion pump wascooled by a liquid nitrogen trap, which was kept permanently cold to

*~~ftork performed in Bell Laboratories, Murray Hill, N.J. 07974, U.S.A.while on sabbatical leave.

68

prevent oil contamination of the chamber. The best residual pressureachieved was 1.7x10" torr. Frequent changes in polarity of the powersupply were performed on turning to FEM experiments or vice versa, andtherefore it was desirable to keep as short as possible the time intervalbetween the two modes of operation, and therefore FEM runs were performedin the 1x10" - 5x10" torr range. For FIM experiments, He gas purifiedthrough a heated vycor leak was introduced to the chamber and the ionpump valve turned off just before this operation. Only the diffusionpump was used to maintain a dynamic pressure of 1.9-2.3x10 torr forthe imaging. The tip was cooled by liquid nitrogen. Heating and flashingof the tip for very short times, and its field evaporation (FE) wasperformed during the experiments as needed.


Attention in this work was directed toward obtaining additional informa-tion by FIM on the effect of adsorbed Zr on W tips, to determine the condi-tions favouring this adsorption to be localized on certain planes and tosee if it corresponds to the observed FEM emission from {100} crystalplanes. As can be seen in Fig. 1 such a confinement of the emission isindeed possible under certain conditions even if a cold cathode is theemitter. In the absence of Zr the pattern of the FEM micrograph shownin Fig. 2 was obtained. Unlike in the presence of Zr, where only {100}planes are emitting, in the absence of Zr a smaller central (100) plane,and planes in the vicinity of {111} faces are emitting. Flashing of thetip at a higher temperature is essential to arrive at the pattern seenin Fig. 1, but it is not yet clear if it is sufficient. The purpose ofthe flashing is to provide sufficient thermal energy for the adsorbedatoms to rearrange themselves by overcoming barriers in their way to thepreferential sites, in our case, to {100} planes. There is a greatprobability to find deposited atoms at a random distribution in meta-stable sices without this rearrangement, as can be seen in Fig. 4. Comparethis FIM pattern with a Zr free pattern seen in Fig. 3. It is not yetclear that in order to obtain the desired emission shown in Fig. 1, thepresence of Zr and flashing of the tip are sufficient. There are indi-cations that this unique emission in <100> oriented emitters might beassociated with blunting and faceting of the tip. These processes mightoccur during the application of a heated cathode, since during theoperation of cold cathodes the pattern of Fig. 1 was repeatedly reproducedon tips which has undergone blunting and faceting.

CONCLUSION

The results of this work indicate that the patterns obtained by thermal-field electron cathodes can be reproduced also in cold cathodes. Thepresence of Zr is essential for achieving it. FIM patterns indicate thepresence of Zr atoms -after flashing- on certain planes and not on others.No conclusive evidence emerges from the work if faceting is an essentialprerequisite to obtain confinement of the emission to the {100} planes.It can be safely concluded that in the absence of Zr no such confinementof the emission is possible.

69

REFERENCES

1.2.3 .4 .5 .6 .7 .3 .9 .

D.R. H e r r i o t t , J . Vac. S c i . T e c h n o l . , 20 (1982) 7 8 1 .R. L iu ,unpub l i shed work, 1983.M. Good and E.W. M i l l e r , Handbuch der Phys ik , 21 (1956) 176.

Shredn ik , Sov. Phys. - S o l i d S t a t e , 1 (1956) 1137.Shredn ik , Sov. Phys. - S o l i d S t a t e , 3 (1961) 1268.Fur se i and S.A. Shakerova, Sov. Phys. Techn. P h y s . , 11 (196u, 827.Swanson and L.C. Crouser , J . Appl. Phys . 40 (1969) 4741 .Swanson and N.A. Mar t in , J . Appl. P h y s . , 46 (1975) 2029.

Fig. 1. Field-emission pattern of zirconiated W[100] t ip. 0.65 KV; 9.1xlO"IU torr .

Fig. 2. Field-emissionQfrom Zr-free W [100] t ip .0.84 KV; 1x10 torr.

70

Fig. 3. FIM pattern of the tip before5Zr evaporation,7.48 KV; He pressure 2.2x10" torr.

Fig. 4. FIM pattern. Distribution of the Zr onthe tip before flashing. 7.75 KV; thepressure 2.2x10" torr.

71

THE TRANSMISSION ELECTRON MICROSCOPY OF NBi-5 BRAZED JOINT

OF INCONEL 718

B. Grushko, 0. Botstein, B.Z. Weiss

Dept. of Materials Engineering

Technion, Haifa 32000, Israel.

The structure of a brazed joint of Inconel 718 with BNi-5 filler metal

was investigated. Transmission electron microscopy, X-ray diffractometry

and EDSA were used in the study of the diffusion zone and of the over-

heated joint.

The diffusion zone in the base material (fig.l) can be divided into two

sub-regions. In sub-region I* the main precipitate is a Nb-rich

G-phase, while in sub-region II, the depth of which can be directly

related to the width of the gap, the precipitates are chiefly carbides of

Nb and Ti (fig.l, 2).

In an overheated brazed joint the filler metal penetrates into the grain

boundaries of the base metal (fig. 3) which results in the formation of

new phases. The dominant one was identified as a hexagonal Laves phase

(Cr,Ni,Fe,Si)2 (Nb,Ti,Mo). It forms an eutectic with the γ-phase in the

grain boundaries (fig. 4a,b).

A fine precipitation of y" and (Nb,Ti)C phases was revealed in the y

present in the grain boundaries (fig. 4c), while in the central parts of

affected grains only (Nb,Ti)C precipitation was found (fig. 5)•

A "binary" phase diagram for Inconel 718 - BNi-5 system is proposed

(fig. 6).

Acknowledgement

The authors would like to express their appreciation to the Wolf

Foundation and the Technion Research Fund for their financial support.

*(close to the actual liqud-solid interface)

72

Fig.l. The diffusion zone SEMimage, deep etchingXI,500

Fig.2. The carbide precipi-tation dark fieldX66,000

Fig.3. The structure of the overheated braze joint.XI,500

73

a) b)

Fig.4. TE microscopy of the grain boundary

region (the overheated joint).

a) The y-A, eutectic with y" precipitation

c) in γ-phase, bright field, X50,000

b) Diffraction pattern of Ai Z.A.=[2423].

c) Diffraction pattern taken from γ-y".

Z.A.=[001]y.

1350

1260

1090

F i g . 5 . Diffraction p a t t e r nof the MC p r e c i p i t a t i o n inthe y matrix in thevicinity of the grainboundary. Z.A.[lll]y.

(Ifi-718) %BNi -5 BNi-5

Fig.6. The scheme of the phase diagram.

Inconel 718-BNi-5 system.

74

SIZE EFFECT IN RADIATION INDUCED CCGREGATION (RIS)

* **L. Kornblit and A. Ignatiev

*Materials Engineering Department ,

Ben-Gurion University

**Physics Department,

University of Houston

A net correlation has been previously established between thevolume misfit factor of a solute atom in a binary metallicalloy and the direction of its radiation induced migration[1]. Following the results obtained by H.W . King [2], whichestablished that the atomic volumes of allotropes changelittle in the transition points, the volume misfit factor wasdefined as the mean atomic volume misfit of the solute withrespect to the solvent, normalized to the concentration ofthe solute [3l. In Ref. 3 a brief summary of reported observa-tions of radiation induced segregation (RIS) in binary alloysis given containing 26 alloys. The prevailing majority ofthese alloys (23) obey the rule stating that a positive volumesize misfit results in a negative (away from the sink) solutetransport and vice-versa. There are however, three exceptions,(Al-Ge, Ni-Ge and Cu-Fe; the first element in the couple isthe solvent). It is difficult to see why the mean atomicvolume misfit rule is violated only for these three alloys,hence here we redefine the volume misfit concept in termsmore applicable to atomic diffusion.

From pseudopotential theory it is known that (at least forsimple metals and alloys) the total energy of the metal(alloy), E, can be presented in the following form

E = E (V) + Z if,

°where E is that part of the energy which is sensitive to themagnitude of the volume V, but does not depend implicitly onthe structure, i.e. is independent of the coordinates of the"pseudoatoms".' The remaining part of energy is the structuresensitive part and is presented as a double sum over all atompairs in the crystal, ij>.. is the interaction energy betweenpseudoatoms i and J. Usually E (V) is the dominant term. Fromthe minimal energy requirement a relation exists between thevolume dependent energy E and the structure sensitive energyterm. Within this frame-work it is easy to understand whyenthalpies of allotropic transformations for an element aremuch less than the enthalpy of sublimation and why the atomicvolume differences during ullotrcpjc transformation are so

75

small. The explanation is that E prevails within the totalenergy expression therefore the tendency in the crystal isto change structure rather than volume. The close-packedbehaviour of metals and alloys is also a result of the^ domi-nation of E ( V).

o

When one characterizes the atomic behaviour of an alloy fromthe standpoint of total energy, (e.g. density, stacking faultcreation) the appropriate characteristic will be the atomicvolumes of the solute and the solvent. But for RIS and soluteatom transport what matters is, in fact, the diffusivity ofthe solutes. Diffusion, however, is a process which leavesthe volume of the crystal unchanged but depends strongly onthe mutual dimensions of the solute and solvent atoms. There-fore, for diffusion processes one should utilize a volumemisfit parameter, which follows from the dimensions of theatoms rather than from the mean volumes of the atoms in theunit cell [h]. This can be done by defining the volume mis-fit parameter as ^ /r )5~1, where rând r s o l y are themetallic radii of the solute and solvent atoms, respectively.For close-packed, or nearly close-packed structures (fee, bee,hep) these radii can be taken simply as half the nearest-neighbour distance or half the bond length of the correspond-ing elements [5l •

We list in Table I the misfit parameter based on solute/solvent metallic radii and compare it to that based on meanatomic volumes. It can be seen in Table I that the discre-pancies noted on ref. 1, with respect to volume misfit basedon mean volume per atom and direction of solute segregationunder RIS are eliminated by using the volume misfit based onatomic size.

REFERENCES

1. L.E. Rehn in Metastable Materials Formation by IonImplantation, '('S'.'T. Picraux and V.J. Choyke , eds. )Elsevier(1982) p. 17.

2. H.W. King in Alloying Behaviour and Effects in Concentra-ted Solid Solutions, (T.B. Massalski, ed.7 Gordon andBreach Tl963) p. 85-

3. L.E. Rehn and P.R. Okamoto, Phase Transformation andSolute Redistribution in Alloys durinr Irradiation, (F.V.Nolfi, Jr., ed.) Elsevier (1983).

It. See, for example, W .B . Pearson: The Crystal Chemistry andPhysics of Metals and Alloys, '//iley-Ir:terscience (1972),p . liHi.

5. Handbook of Chemistry and Physics, 5^th Edition, CRC cress(1973) p. F197.

76

Table 1. Volume misfit parameters based on atomic size (r . , rsolv^ o r

mean atomic volume (King) determined for a number of solute/solventsystems. Note that the discrepancies in the direction of segregation(under RIS) predicted from the King volume misfit parameter are removedwhen the atomic size volume misfit parameter is used.

Alloy

Pd-CuPd-FePd-MoPd-NiPd-W

Al-GeAl-SiAl-Zn

Cu-AgCu-BeCu-FeCu-Ni

Ni-AlNi-AuNi-BeNi-CrNi-GeNi-MnNi-MoNi-SbNi-SiNi-Ti

Ti-AlTi-V

Fe-CrMg-Cd

r s o l v ^

1.3775

1.4315

1.278

1.2458

1.4478

1.241151.59855

so l 1 •*

1.2781.21151.362551.24581.37095

1.22491.175851.3397

1.44471.11301.241151.2458

1.43151.442051.11301.24901.22491.365551.362551.4501.75851.4478

1.43151.3112

1.24901.4894

Volume Mis f i t %

< r so l / r so lv> - 1

- 2 0- 2 7- 3- 2 6- 2

- 3 7- 4 5- 1 9

+44-34- 8- 7

+52+55- 2 9+1- 5+ 32+ 31+58- 1 6+57

- 3- 2 6

+2-19

Direc t ion of Volume Mis f i t %Segregat ion After King [2](from Ref. 1)

+ -19+ -12+ -4+ -14+ -4

+ +13+ -16+ -6

+44+ -26+ +5+ -8

+ 15+64

+ <0+10

+ +15+23+22+ 21

+ -6+29

+ -20+ -15

+4+ - 2 1

77

APPLICATION OF ADVANCED COMPUTER GRAPHIC DISPLAY TECHNIQUES TOMATERIALS PROBLEMS: PHASE EQUILIBRIA AND DIFFUSIONAL GROWTH

M. R. Notis, S. K. Tarby and J. I. Goldstein

Department of Metallurgy and Materials EngineeringLehlgh University, Bethlehem, PA 18015 USA

Phase diagrams and the understanding they provide of equilibriumphase relationships in multicomponent systems are of major importanceto both basic education and to advanced research in all areas ofmaterials science and engineering. Binary systems are reasonably smallin number and simple enough so that phase relations can be fullyvisualized in two-dimensional displays. However, most materials ofengineering interest are ternary or higher order systems and theirphase relations are significantly more difficult to visualize. Whilemajor advances have been made in the past decade concerning the calcu-lation of phase diagrams from thermodynamic data, the development ofmethods for the graphic display of these phase diagrams has limited ourability to utilize the information generated and hence to understandand solve applied problems involving phase equilibria. We have recentlyused a wide variety of CAD (Computer Aided Design) systems in to displayboth binary and ternary phase diagrams and have found these systems tobe of extreme importance in both educational and research areas.

1. Phase Diagrams—Background

The phase diagram for a binary system, which consists of twocomponents, is easily portrayed in two dimensions (again, assumingpressure fixed). For this situation, Gibbs phase rule (F + P = C + 1)indicates a unique composition for each of the phases in equilibriumwithin a two phase field at constant temperature. Therefore, knowledgeof temperature enables the determination of equilibrium phase composi-tions, and knowledge of the nominal composition of the system allowsthe calculation of the mole fraction of each of the phases by applica-tion of the reverse lever law to the composition tie-line. Computerapproaches to display binary phase diagrams have been developed morethan a decade ago, and there is currently an extensive program organizedthrough the American Society for Metals and the American CeramicSociety to tabulate data and generate binary diagrams. The NationalBureau of Standards has technical responsibility for this program.

If another variable is added, either as a third component, or bymaking the pressure variable, the use of a third dimension is requiredto accurately portray the equilibrium phases present. However, it iscommon to display these three-dimensional diagrams in two-dimensionalsectional views, and hence visual perception of proper phase relationsvery often limits the utility of these phase diagrams. This problem

7S

has generated a host of approaches to the display of ternary phaserelations. These include the use of coordinated sets of two-dimensionalternary diagrams, transparent overlay projections, solid models made ofplaster of Paris or soap, and the use of transparent sheets of glass orplastic to represent isothermal planes. Also, a film series on phasediagrams is available from Pennsylvania State University which demon-strates some simple binary and ternary computer generated diagrams,Finally, two dimensional isothermal sections of ternary diagrams aregenerated by computer in the phase diagram development program of theNational Bureau of Standards.

2. CAD Systems and Phase Diagram Display

As will be described in the following section, CAD computersystems have been shown to be ideally suited to phase diagram display.The CAD facilities include the following computer-graphics systems:

Digital Equipment Corporation VAX 11/780Evans U Sutherland PS300 Dynamic Display System (E/S)ApplicoTi DECSYSTEM 11/34IBM 4341 Model Group 2 (vm) with IBM 3250 terminals.

Turnkey software systems include McAuto Unigraphics (McDonnell DouglasAutomation Company), SDRC (Structural Dynamics Research Corporation),and Applicon's Solids Modeling package. The Applicon software runs onits own hardware (PDP11/34), and the VAX11/780 is used to support theMcAuto and SDRC software. The software on the VAX11/780 is operationalin an interactive mode using VAX-11 Fortran. In addition, the softwarepackages are capable of utilizing a finite element generator and MOVIE.BYU graphics. The VAX11/780 processor has six CAD work-stations (eachhaving a Tektronics 4014, alphanumeric screen, function keyboard, anddata tablet) and six VS11 color terminals.

Most of our work has been with the Tectronics 4014 using Uni-graphics software; this system has an extremely high resolution displaybut is currently limited to a 'green screen.' A new Unigraphics systemwith high resolution color capability is currently being installed. Sofar, the majority of our color display work has been on the VS11terminals. Both of the above systems (Tectronics 4014 and VS11) arelimited in terms of the speed of dynamic display. We have performedlimited work with the high speed E/S system and we are developingprocedures to shift data and program files developed on the Unigraphicssystem over to the E/S system. We are also working with both Appliconand IBM 4341 display systems and are comparing the input and outputadvantages and problems for all systems. Because some systems haveunique advantages for either data input or display output we are tryingto develop universal file transfer software so that we can takeadvantage of the benefits of each particular system.

Data input to CAD systems may be accomplished in a number of ways:by entering individual data points numerically into the system from thekeyboard, by using equational forms for multiphase surfaces fromprograms written for a specific system, or by digitizing devices fromhard copy of a phase diagram. For example, diagrams published in Bull,of Alloy Phase Diagrams may be laid upon the digitizer pad and tracedover to input to the computer in just a few minutes each.

79

Once a phase diagram is available, there are a number of specificfeatures or uses of the diagram that might be of interest. Theseinclude:- the number of phases present at a given alloy composition and

temperature- the maximum solubility of the components in a given phase at the

temperature of interest- the equilibrium compositions of coexisting phases in multiphase

regions (i.e., the "tie line" compositions)- the weight fraction or atom (uole) fraction of phases present

(i.e., using the lever rule)- the ability to follow the course of crystallization (solidificatiir.x)

and tc understand the nature of the phase transformations whichoccur (i.e., reaction types)

With this in mind, the following present capabilities of the CADsystem have been established:- multimode data input with direct access to interactive

computation (Fortran) and simultaneous display; data may be inputusing any temperature scale and in either weight or atomic percentfor composition

- "wireframe" of solid model in three dimensions by high resolution(single color—green) display

- color graphics display of wireframe model- color graphics display of solid model with color shading to enhance

3-D perspective- "exploded view" capability in wireframe or solid model (i.e., "solid

puzzle" view)- any arbitrary surface view or rotation as chosen- any isothermal section or composition cut- split screen, multiview capability (four separate views

simultaneously)- selection of specific phase diagram area and blow-up (expanded view)- color coding of individual lines- heightened intensity or variable line width for individual lines

(future)- hidden line styling (solid, dash, dot, alternating pattern, etc.)- hidden line removal- output to terminal, multicolor printer/plotter, or to videotape.

These capabilities can be demonstrated with a number of examples.First, and easiest, binary phase diagrams can be displayed and comparedwith experimental data. Figure 1 shows the most recently published'provisional' version of the Al-Ag phase diagram (1). Superimposed onthis diagram is the low temperature experimental data for the equilib-rium solvus line, obtained by analytical electron microscopy (2). Theexperimental data can be 'curve-fit' by a least squares quadraticspline, with the simple push of a button, and this curve compared tothe published diagram. Also shown in Figure 1 is the data for themetastable GP Zone solvus (2). In the color version of this figure,the diagram may be displayed in one color and each of the data sets ina separate color thereby making multiple data sets quite easy tocompare.

A similar example, this time involving experimental diffusion

so

studies, is shown for the Ni-Al system in Figure 2. The interfacialcomposition data obtained from a recent study (3) concerned with thegrowth of the Ni Al phase formed during interdiffusion in a NiAl vs Nidiffusion couple, is shown superimposed on the Ni-Al phase diagram. Weare now in the process of generating a ternary phase diagram for theNi-Al-Cr system (4) in order to display the ternary diffusion dataobtained from a research program now in progress.

The final example (Figure 3) shows a hypothetical ternary eutecticsystem with partial solid solubility of the three components. Thistype of display can be formed from a two dimensional figure withisothermal contour lines, such as are found in textbooks or journals.The perspective view shown in Figure 3 allows easy visualization of theisothermal intersections with the liquidus and solidus surfaces.

The interactive capability allows complete interfacing of displayand calculation features of the system. For example, the system iscapable of calculating the weight fraction or atom (mole) fraction ofeach phase in a two-phase field, given the temperature of interest andthe nominal composition of the alloy. This is accomplished by firsttyping the nominal composition of the alloy and the desired tempera-ture. The computer then locates this composition as a point in theGibb • triangle base-plane (Figure 3) and projects a vertical linethrough i' parallel to the temperature axis (>.~axis). Selection of thetemperature identifies the isothermal plane of interest. Further,tie-lines may be drawn to determine the composition of phases inequilibrium at the temperature of interest. The tie-lines may be inputfrom experimental data, or, with the simplifying assumption that thetie-line is directed to the ternary end-member corner, it may becreated by extending a line from the ternary corner in the isothermalplane through the intersection of the vertical (composition) line andthe isothermal plane. In any event the intersections of the tie-linewi h, for example, a liquidus surface on the one end, and the solidussurface on the other end, may easily be determined. Softwarecapability built into the system determines the length of tie-line fromthe composition vertical to either the liquidus or solidus surface andthe ratio of this length to the total tie-line length gives thefraction of each phase present under equilibrium conditions. Thecomposition of each ternary phase is also printed out.

In summary, CAD computer display system- can be used for datastorage and display of complex phase equilibria. The interactive modeallows for rapid comparison cf the stored phase diagrams with new inputdata. The dynamic display capability allows for the visualization ofrelationships not easily discernible from two-dimensional staticfigures such as commonly found in textbooks and journals.

References1. R.P.Elliot and F.A.Shunk, Bull. Alloy Phase Diagrams, _1_, 36 (1980).2. S.M.Merchant, M.R.Notis and D.B.Williams, Met. Trans., 14A, 1825

(1983).3. R.Glitz, M.Notis and J . I .Goldste in , p.691 in Solid-Solid Phase

Transformations, H.I.Aaronson, et a l . , eds . , Met. Soc, A1ME (1982).4. S.M.Merchant and M.R.Notis, J . Mat. Sci. & Eng., to be published.

81

Acknowledgment

The authors thank Mr. Jeff Roeder and Ms. Cathy Curtin for theactual generation of the computer f igures.

Ag A I

oLo>aa>

00 10 20 30 40 50 60 70 80 90 100

Weig^t P e r c e n t Aluminum

Figure 1. The Ag-Al system (1)with experimental data (2) super-imposed .

1600

1S00

1400

1300

1200

1000

900

BOO

700

600

M Ni

/TX• : / / : \ — -

LΑ

-•

•

/

\

00 10 20 30 40 50 60 70 80 90 100

A.tom ' c Percent

Figure 2. The Al-Ni system withexperimental data (3) superimposed.

B

Figure 3. Schematic diagram of ternary eutectic system showing iso-thermal planes and tie-lines.

82

COERCIVITY, SQUARENESS RATIO AND MICROSTRUCTURE IN Co-WTHIN FILUS

U. Admon(a), M.P.Dariel^, E. Grunbaum , G.Kimmel(al and J.C.Lodder^

(a) Dept. of Metallurgy (b) Dept. of Physics, Nuclear Research Center -Negev, Israel. (c) Dept. of Electron Devices and Materials, Tel-AvivUniversity, Ramat Aviv, Israel, (d} Dept. of Electrical Engineering,Twente University of Technology, Enschede, The Netherlands.

ABSTRACTThin Co-W magnetic films can be electrodeposited in a variety of

non-equilibrium structures. By a proper choice of the depositionparameters it is possible to obtain alloys displaying a wide range ofmicrostructures and, consequently, of magnetic properties. Co-W films,200-500A thick, were electrodeposited under various plating conditions.The coercivities and squareness ratios of these films were determinedby magnetometry, their microstructures by transmission electron microscopy.A close correlation between the magnetic properties and the respectivemicrostructures has been established.

INTRODUCTIONCobalt is readily electrodeposited from aqueous solutions as a pure

metal or in the form of various binary, ternary and even quarternaryalloys. Over 55 such alloys have been reported in the literature (1).Tungsten, on the other hand, cannot be electrodeposited as a pure metalfrom aqueous solutions neither can it be co-electrodeposited with otherelements, with the exception of iron, cobalt or nickel (2). In the caseof Co-W, deposits with up to 65 wt./o (37 at./o) tungsten were obtained.Electrodeposited thin films often exhibit non-equilibrium phasestructures which depend strongly on the deposition process parameters.This is particularly true in the case of cobalt based alloys for whichthe presence of the non-equilibrium fee phase is attributed to thesluggishness of the fee to hep transformation.

Omi et al. (3) studied thick (20mn) Co-W films by X-ray diffraction.Their deposits contained varying proportions of a crystalline hep phaseand a non-crystalline phase corresponding to the C03W stoichiometry.

Rachinskas (4) Polukarov f5) and Armyanov and Vitkova (6)determinedthe magnetic properties of thick (over lpm) Co-W films at various platingconditions. They reported coercivities in the range of 200-600 0e, andsquareness ratios of 0.6-0.8.

EXPERIMENTALThe cobalt-tungsten films were electrodeposited onto coppei-coated

microscope slides. The slides had been pre-coated with a thin layer offormvar. The specimens for the magnetic measurements were prepared bycutting the slides into 1 r.m squares.

The specimens for transmission electron microscopy were preparedby cutting 2 mm- squares on the coated surface, floating off the squares

acidic pH,g/l60

0-5

50

30

-

-

0.1

H S04

basic pH,g/l80

0-80

-

-

350

50

-

NH4OH

83

by dissolving the formvar, and finally selectively dissolving the copperas described elsewhere (7). The following solutions were used forelectrodeposition:

CoSo4•7H20

Na2WO4•2H70

MgSO4-7H2O

H3 B O3

Rochelle Salt-4H90

(NH4)2S04

Sodium Lauryl Sulfate

pH correction

The bath temperatures used were 22°, 50 , and 85 C, at current density of10±l mA/cm2. The deposits had a gray metallic appearence with increasingbrightness at elevated bath temperatures. At pH=2 the solutions wereunstable and tended to precipitate. The deposits were dull, and it wasnot possible to obtain uniform samples for the magnetic measurements.

The magnetic measurements were carried out using a Foner typevibrating sample magnetometer at the Twente University of Technology.The microscopic examinations were done at 100 kV on a JF.OL-JEM 7Atransmission electron microscope at the NRCN laboratories.


The detailed correlation between the microstructure of the deposits(grain morphology, phase constitution and texture) and the depositionparameters is beyond the scope of this paper, and will be reportedelsewhere. However, as the magnetic properties are determined primarilyby the microstructure, the main microstructural features that have beenobserved will be described here.

Two types of structure were obtained for different depositionparameters, and are shown in Figure 1. In most cases, the deposits werecrystalline (Figure la and lb), consisting of the hep and the (unstable)fee phases in various proportions. The hep crystallites showed a varyingdegree of texture, with anCoO.l] fiber axis perpendicular to the filmplane. However, Co-W films deposited from basic baths at roomtemperature and salt weight ratios above 80/10 (these numbers, and otherswhen quoted, give the C0SO4 to Na2WC>4 salts weight ratios in (g/1) in theplating baths) were amorphous, or composed of crystallites of a verysmall grain size, as can be seen fTom their diffuse electron diffractionpatterns (Figure lc).

Figure 2 shows schematic diagrams in which the crystallography andtexture of the deposits, obtained from the electron diffraction patterns,are represented as a function of the bath composition, temperature andacidity. It should be noted that the corresponding tungstenconcentrations in the films ranged from 7 to 34 wt/o. Its dependence onthe plating conditions will not be described here in detail.

84

The typical grain morphology of the crystalline deposits is shownin the transmission electron micrographs of Figure la and lb. The hepand fee grains are qui-axed and their size is in the range of a fewhundred Angstrom. In deposits from acidic baths the mean grain size, aswell as its spread, increased with the bath temperature (Figure la). IΓ.deposits from basic baths and above room temperature large singlecrystals, a few thousand Angstrom in diameter were evenly dispersed ina population of small crystallites, a few hundred Angstrom in diameter(Figure lb). The large crystals often showed planar defects and asubgrain structure, probably due to the grain structure of the coppersubstrate. A similar effect was observed in the amorphous films (Fig.lc)

The coercivity, He, and the squareness ratio, R=Ir/I

s (I

r and I

s

are the residual and saturation magnetization, respectively), obtainedfrom the hysteresis loops, are summarized in Figure 3 and 4. Thefollowing conclusions may be drawn:

(a) The Co-W films have a medium to high coercivity. It ranges from100-600 Oe in the crystalline films and decreases to 20-30 Oe inthe amorphous films.

(b) He is low for films which contain a high proportion of feecrystallites, and increases with increasing content of hepcrystallites. For amorphous films He is even lower. Thisstems from the differences in magneto-crystalline anisotropy,which is the highest for the hexagonal phase and zero forthe amorphous phase (compare Figure 2 with Figure 3).

(c) He decreases with increasing perfection of texture ([00.l]perpendicular to the film plane) . This can be seer, atpH=8.5; T=85°C when the bath composition changes from 80/20to 80/80 (see Figure 3), and at pH=6 when the bath temperatureis increased from 50° to 85°C. This is consistent with thefact that the[00.l]axis, which is the easy axis of magnetizationof the hep crystallites, approaches the normal to the filmplane.

(d) He increases when the population of large cr>stals increases.This occurs at pH=8.5 at higher temperatures and tungstenion concentrations in the bath. However, this effect issmall compared to the phase composition effect. This resultis in full agreement with those of Armyanov and Vitkova (7)but only in partial agreement with those of Rachinskas (4).

(e) The squareness ratio, R, increases in deposits of higher feephase contents, and reaches values as high as 0.9. This canbe attributed to the presence of four independent easy axesof magnetization in the fee phase, in contrast to the singleaxis in the hep crystallites. Thus, when the external fieldis being removed there is a greater chance for the magnetizationvector to jump into an easy direction close to the field direction.The theory (8) predicts R values of 0.866 for fee and 0.5 for thehep phases. Our results agree, within the experimental error,with these values.

(f) The amorphous deposits have a high value of R, as might beexpected. However, R being less than unity indicates that thereis some degree of short range order in the films. This assumptionis supported by Omi et al. (3).

(g) R decreases when the degree of texture perfection ([00.l]perpendicular to the film plane) increases.

85

(h) R decreases when the spread in grain sizes increases. This wasobserved for baths with increasing tungsten ion concentrations,and agrees with the results of Rachinskas (4).

(i) There is a crystallographic and magnetic azimuthal symmetry inthe plane of the films. This was deduced from the symmetry ofthe diffraction rings and confirmed by in-plane magneticmeasurements at various directions.

CONCLUSIONSThere is a strong influence of the deposition parameters

(particularly the temperature, pH, and composition of the bath) on themicrostructure, and hence on the magnetic properties. By a carefulchoice of these parameters it is possible to facilitate the productionof magnetic alloys with a wide range of properties.

ACKNOWLEDGEfENTS

The authors wish to thank Dr. T. Wielinga of the Twente Universityfor several helpful discussions, and Mr. B. Yusov of the NRCN for thepreparation of the specimens.

REFERENCES

1. Krohn, A., C.W. Bohn, Plating, March 1971, 237.2. Sastry, B.S.R., Metal Finishing, Oct. 1965, 86.3. Omi, T., H. Yamamoto and H.L. Glass, J. Electrochem. Soc, 119

1972, 168.4. Rachinskas, V.S., in "Electrodeposition of Metals" Proc. 10th

Lithuanian Conf. of Electrochemists, Dec. 1968, translated fromRussion by the Israel Program for Scientific Translations,Jerusalem,1970, p. 51.

5. Polukarov, Yu. M., in "Electrodeposition of Alloys", V.A. Averkin,ed., Moscow 1961, translated from Russian by the IPST, Jerusalem,1964, p. 52.

6. Armyanov, S., S. Vitkova, Surface Technology, 1_, 1978, 319.7. Admon, U., A. Bar-Or and D. Treves, J. Appl. Phys. 44_, 1973, 2300.8. Chikazumi, S., "Physics of Magnetism", John Wiley, 1964, Ch. 12.

86

Typical microstructures of Co-W thin films (for the notationsee text) :

(a) crystalline, uniform grain size (pH=4; 85°; 60/5).(b) crystalline, non-uniform grain size (pH=8.5; 85°; 80/40).(c) non-crystalline (pH=8.5; 22°; 80/40).

T,°C20 ""'40

S.W.R.Fig. 2: Dependence of phase composition and texture on the deposition

parameters:

(a) acidic baths, S.W.R. = 60/5(b) basic baths, S.W.R. = 80/0 to 80/80S.W.R. is the Salts Weight Ratio (CoSO4 to Na2W04) in the baths.

Dots-fcc, bars-hep, shaded area-non-crystalline.The [00.l] texture is illustrated by the degree of alignment ofthe bars.

87

22°

T,°C

S.W.R.Fig. 3: Dependence of the coercivity, He, on the deposition parameters:

(a), (b) same as Figs. 2(a), 2(b), respectively.

80/0 30/10 80/2(} 80/ , 0 R0/R0

S.W.R.

Fig. 4: Dependence of the squareness ratio, R=Ir/js,on the depositionparameters:

(a), (b) same as Figs. 2(a), 2(b), respectively.

88

STRUCTURAL ANALYSIS OF HIGH VACUUM, HIGH TEMPERATURE

BNi-5 BRAZED JOINTS OF INCONEL 718 SUPERALLOY

B. Grushko, B.Z. Weiss

Dept. of Materials Engineering

Technion, Haifa 32000, Israel.

INTRODUCTION

In the present paper the results of structural investigations of the

brazed joint of Inconel 718 by BNi-5 filler metal are reported.

Interdiffusion of elements between the base material and the brazing joint

may lead to the formation of new phases (alloys) in the intermediate zone,

which have a lower melting temperature than the parent material, resulting

in mobility of the "real" liquid-solid interface. It is obvious that

compositional elements of filler and base metals are present on both

sides of the former "real" liquid-solid interface, while "location" of

these elements is controlled mainly by thermodynamic and "environmental"

factors as well as gap clearance.

It is therefore reasonable to assume that the post-brazing structure of

the "brazing-influenced zone", which includes the brazing joint and some

adjacent areas of the base material, is influenced mainly by three

factors, viz. temperature, time, and gap width (TETIG).(1).

THE EXPERIMENTAL PROCEDURE

Three methods were used for the purpose of identification, namely: optical

and SEM microscopy, energy dispersion analysis (EDSA) , and X-ray

diffractometry. Specimens (fig.l) were brazed in a vacuum resistance

furnace in accordance with thermal regimes described in fig.2. X-ray

diffraction measurements were conducted on section α-a of the specimen II.

A deep etching technique was applied in order partially to dissolve the

surrounding y matrix. The etchant was 5% Nital electrolyte applied for

5-8 minutes with a potential of 6V and a current density of 1.3 mA/mm2.

For metallographic investigations the specimens were etched with aqua

regia. In addition, a thermal etching technique was used, in which the

specimens were heated to 700eC, kept for 10 min at that temperature, and

then cooled in air.

RESULTS

The depth penetration of the liquid into the base material,

measured in the JS-μ section (see fig.3), was found to average 30 um for

II-A ) specimens, 60 ym for II-C specimens, and 440 Um for specimen II-D.

*)

II, III are specimen configurations (fig.l); A,B,C,D are the thermal

regimes according to fig.2.

89

X-ray diffraction results on deep-etched specimens are shown in fig.4.

In the II-A and II-C specimens, lines of phase 6-and G-phase were

observed (γ-phase was dissolved). In specimen II-D only G-phase lines

were observed in addition to the lines of the γ-matrix, which was, in

this case, only partly dissolved. Because of the directional solidifi-

cation, quantitative analysis was practically impossible. Rotation of

the specimens did not change the intensity of the 0 lines, while peaks

of the G-phase disappeared at certain orientations of the specimens and

reappeared in others. This indicates a uniform distribution of the 6

compound, most probably as a fine constituent of the eutectic. The

meta.llographic examinations showed the presence of the dominant γ-solid

solution, coarse irregular eutectic of the G-phase and y of a flower-like

morphology, irregular eutectic of 0 and γ (the 0 and G-phase could be

distinguished with great certainty by the "EDS analysis).

In the II-C specimen two fine eutectics could be observed, being coloured

differently by thermal etching. Approximate calculations of the

eutectics1 compositions led to the following results: the "bright"

eutectic: 40% G and 60% 0-phase; the "dark" eutectic: 60% G, 30% 0, and

10% γ-phases. In specimens II-A and II-C the γ-phase was found to

contain small quantities of Fe, Nb, and Mo. In the specimen, II-D, the

γ-phase contains significant quantities of the base metal elements, but

the concentration of Si was found to have considerably decreased. EDS

point analysis showed that the Nb replaces Cr in G and its content in the

G-phase increases when the heating regime is changed from II-A to II-D,

which also results in an increase of the lattice parameter from 11.12A to

11.22A (see fig.4).

The 9-phase does not show any presence of Nb. The results of metallo-

graphic studies and EDS analysis of the microstructure in the wide-gap

joints, (Specimen III brazed by thermal regime A, b>0.25 mm) were

basically similar to those described previously for specimens IIA and

IIC. As the gap narrows (<0.1 mm), a joint with a centerline eutectic

is formed, as shown in fig.5. The structure consists of large bright

G-phase particles and small gray areas of 0. The composition of the

G-phase is very similar to that obtained in II-D specimens despite the

lower brazing temperature applied.

Two effects were observed as a result of the gap's narrowing: a) the

volume fraction of 6 decreases; b) Nb content in the G-phase increases.

In joints with gaps narrower than "V50 ym only the presence of che G-phase

could be observed,in joints 30 ym G-phase is presented as a discontinu-

ous chain of small particles.

DISCUSSION

On the basis of the results obtained it can be concluded that the phase

formation in the BNi-5 brazing of Inconel 718 can be basically described

by the Ni-Cr-Nb-Si quasi-quarternary system. Titanium can be considered

as an element interchangeable with niobium, while nickel and chromium

are sometimes replaced by iron. This approximate phase diagram (fig.6)

was constructed by using four ternary diagrams known from the literature

(2,3,4,5). A few alterations, based on experimental results, were

introduced; e.g. the TT-phase (of the Ni-Cr-Si system2) was not included

90

since it was not observed experimentally. The absence of the TT-phase

suggested the formation of an eutectic between the G- and the γ-phases,

which was subsequently confirmed experimentally.

Occlusion of isomorphous phases, such as NbgNii6Si7, CrgNi^Siy (G) and Xj_-

hexagonal Laves phases, leads to the creation of regions in which only a

single one of those phases is present. It should be assumed that the

phase distribution and compositions at room temperature are very close

to those shown in the diagram (fig.6). The initial location of the BNi-5

alloy is represented by point (1) . This point is approximately situated

at the boundary of the region y+Q and the region y+9+G. The influence

of the base metal on the structure of the NBi-5 brazing can be illustrated

by shifting the figurative point from position (1) to position (2), which

represents the approximate location of the base material - Inconel 718.

(see fig.6b; the straight line approximation was used), Initially, the

Y-6-G structure is present in the molten alloy. Gradual shifting along

the line (1-2) (see fig.6b) shows that once the 6 phase disappears, the

structure should consist only of G- and γ-phases, and finally only the

γ-phase remains. Experimentally this situation is simulated by varying

the width of the gap. The wider the brazing gap, the closer the phase

constitution is to position (1), and the narrower the gap, the closer

the constitution is to position (2). The predicted changes in volume

fractions of the G-and 6-phases were in practice shown to occur by

changing the width of the gap.

The wide spectrum of brazing experiments, conducted in different thermal

conditions and for different brazing gaps, showed that all the phases

that appear in the constructed phase diagram were actually observed in the

brazing joint. T-iis means that, although the brazing process does not

seem to follow conditions of equilibrium, yet the formation and outward

appearance of the different phases can be predicted from the phase

diagram.

ACKNOWLEDGEMENTS

The authors would like to express their appreciation to the Wolf

Foundation and the Technion Research Fund for their financial support.

REFERENCES

1. R. Johnson, Weld. res. suppl. 1981, Vol.60, pp.185-193.

2. E.I. Gladyshevskii and L.K. Bornsevich., Russian J. Inorg. Chem. 1963,

Vol.8, pp. 997-1000.

3. E.I. Gladyshevskii et al., Inorg. Mater., 1969, Vol.5, pp. 1882-1883.

4. H.J. Goldschmidt and J.A. Brand, J. Less-Common Metals, 1961, Vol.3

5. L. Kaufman and H. Nesor. Proceeding of the conference on In Situ

Composites. Sept. 5.8. 1972 Lakerville. Conn. Vol. Ill, Publ.

NMAB-308-III, Washington D.C. 1973.

91

11Fig.l. The specimens. Fig.2. The thermal cycles of brazing

Fig.3. The B-β section

of specimen Il-C, aqua

regia, X100.

Fig.4. Diffraction pat-

terns for type II speci-

mens .

Fig.5. Micrograph of

III-A specimen, thermal

etch., X600

(a)

Ni-Cr

(b)F i g . 6 . Schematic diagram of the Ni-G-Si-Nb system for

800-1000°C(a) general view; (b) te rnary s e c t i o n .

92

PHASE RELATIONS IN THE Cu-Nd SYSTEM ON THE Cu-RICH SIDE

1 2 2 3C. Laks , J. Pelleg , and L. Zevin '

Israeli Military Industries, Ramat Hasharon, P.O. Box 1044,2 3Materials Engineering Department and Institutes for Applied Research,Ben-Gurion University of the Negev, Beer-Sheva, Israel

Major inconsistencies in RCu systems (R = rare earth element) occur inthe 75-88 atom % Cu range (1-7). No data are available for the NdCusystem. The objective of this communication is to compare the existingRCu systems and decide which of them behaves similar to the NdCu system.

EXPERIMENTAL

Alloys of Nd-Cu were prepared with 63-93 atom % Cu by arc melting andwere annealed for homogenity. The specimens were examined by X-raydiffraction, differential thermal analysis (DTA), and metallography.Standard techniques were used for metallographic examination by opticalmicroscopy and scanning electron microscopy (SEM) with energy-dispersiveanalysis (EDAX).


Table 1 listt- •-. - '• ntermetallic compounds identified in this work togetherwith informa-":u available frore the literature, and Fig. 1 is a tentativediagram of the phase relations in the system under consideration. Fivecompounds were identified, NdCug being the richest in Cu. A previouslyunknown NdCu4 compound was also detected. Most of the peaks could be in-dexed on the basis of an orthorombic cell, but definite structure deter-mination has to await work with single crystals. The absence of certainsystematic extinctions rules out the possibility of the Pnnn space groupsuggested for RCU4 type couponds (1). It is interesting to note that inthe LaCu system LaCu4 was not detected, unlike in other systems (1,5)where R was a light rare-earth element. This difference between the LaCuand NdCu systems merits further investigation.

CONCLUSIONS

An orthorombic NdCu4 compound was detected in the NdCu system and in thisregard it resembles the CeCu and SmCu (7) systems. No definite structuredetermination was possible by the technique used, and complete evaluationof its structure will require single-crystal work.

93

Table 1. Crystallographic data for phases in the Cu-Nd System

Phase Crystal system and Lattice parameters (A)structure type a b c

Reference

NdCU-o

NdCu

NdCu

NdCu

NdCu

Orthrombic;

Hexagonal;

Unknown

Orthoronibic

Orthorombic

CeCucb

CaCu,-

; CeCu2

; FeB

8.092

7.952

5.104

5.097

-

4.387

4.384

7.32

7.302

5.

5.

7.

7.

4.

4.

062

044

-

-

-

059

096

55

569

10.105

10.203

4.107

4.112

-

7.420

7.417

5.59

5.578

8

This work

This work

This work

This work

9

This work

10

This work

REFERENCES

1. T.B. Rinehairaner, D.E. Etter, J.E. Selle and P.A. Tucker, Trans.Metall. Soc. AIME, 230 (1964) 1193.

2. S. Carifici and A. Palenzona, J. Less-Common Met., 53 (1977) 199.

3. E. Franceschi, J. Less-Common Met., 87 (1982) 249.

4. A. Iandelli and A. Palenzona, J. Less-Common Met., 25 (1971) 333.

5. K.H.J. Buschow, Philips Res. Rep., 25 (1970) 227.

5. L.A. Ofolubkov, N.M. Shibanova, Yu. G. Sakonov and G. Ya. Fedorova,Russ. Metall. N6 (1977) 147.

7. K. Kuhn and A.J. Perry, Met. Sci. J. 9 (1975) 339.

8. K.H.J. Buschow and A.S. Van Der Goot, J. Less-Common Met., 20 (1970)309.

9. A.R. Storm and K.E. Benson, Acta Crystallogr. 16 (1963) 701.

10. A. Iandelli and A. Palenzona, in K.A. Gschneidner and L. Eyring(eds.), Handbook on the Physics and Chemistry of Rare Earths. NorthHolland, Amsterdam, 1979.

94

noo \r

1000

s>2 900

2S.S00

700

600

0

\

\

So <5 ota t3 5>

o1

917*

866°829" 843*

. .

0 10 15 20 25 30 35Atomic percent Nd

Fig. 1. Partial Cu-Wd diagram.

95

EFFECT OF SMALL ADDITIONS ON GRAIN REFINEMENT OF A 14 CARATAu-Ag-Cu-Zn ALLOY

*M. Fishman, L. Gal-Or, G. Iram

Center for Noble Metals, Institute of Metals, Technion, Israel

* Dumax Corp.

INTRODUCTION AND EXPERIMENTAL PROCEDURE

The process of jewellery fabrication by cold working inevitably in-volves intermediate annealings aimed at restoring alloy ductility by re-crystallization of the work-hardened structure. On the other hand, theannealing treatment can lead to a coarse grained structure as a result ofexcessive grain growth following recrystallization. Due to the fact that,plastic flow in crystals is orientation dependent on further deformation,,especially if operations like deep drawing or bending are involved, thisstructure can manifest itself undesirably by producing a surface effectknown as "orange peel".

The purpose of this work was to find an addition (or a combination ofadditions) to a 14 carat Au-Ag-Cu alloy containing 6% Zn that would actas an effective grain refiner of the re crystallized structure in a widerange of annealing temperatures and, preferably, as refiner of the caststructure as well.

The following additions and their combinations were investigated (wt % ) .Ir 0.02, 0.05; Zr 0.04, 0.05, 0.1; Co 0.05, 0.1, 0.15, 0.2, 0.3; B 0.005,0.01, 0.1; 0.1 Zr + 0.005 B; 0.15 Co + 0.005 B; 0.1 Co + 0.05 Zr +0.005B.Alloys were prepared by induction melting in graphite crucible under pro-tection of argon. The solidified structure was investigated on samplescut from both bottom and top sections of the ingots. The rest of the in-gots was cold worked into wire 1 mm in dia with intermediate annealingsat 680°C. The final annealing treatment preceding determination of re-crystallized grain size was performed at 600, 650, 680, 700 and 750°C atvarious time intervals from 0.5 min to 2 h.

THE INFLUENCE OF SMALL ADDITIVES ON THE ALLOY STRUCTURE

Ingots of the unmodified 14 carat Au-Ag-Cu-Zn alloy showed a well pro-nounced dendritic segregation [fig- 1 left) with dendritic branches beingcoarser in the upper part of the ingots. The grain size of the annealedmicrostructure depends drastically on annealing time and temperature. Forexample, the rise of temperature from 600 to 750°C results in 20-fold in-crease of average grain size after 15 min annealing.

Boron and cobalt introduced separately had no grain refining effect oneither cast or annealed structure of the alloy. Zirconium has a slightgrain refining effect on the recrystallized structure that is more markedat higher Zr content. Cobalt and zirconium are much more effective whenintroduced in combination with boron. Combined alloying with 0.1% Co,0.05% Zr and 0.005% B has a strong and reproducible grain refining effecton the recrystallized structure at all annealing temperatures. However,alloys with the complex Co-Zr-B addition have a higher hardness and lowerplasticity than the unmodified alloy.

Iridium produces drastic grain refining of the cast structure (Fig.l)and also substantially reduces the recrystallized grain size (Fig. 2B).Both effects enhanced in the alloy with higher Ir content. A fine grainedrecrystallized structure in iridium containing alloys is retained after

96

threefold remelting. The grain refining effect of iridium on the caststructure was even enhanced by remelting. Iridium does not reduce al loyduc t i l i ty . Tensile tes ts on 1 mm wire showed an increased alongation andunchanged s t rength. In order to verify whether the .grain refinementachieved by addition of 0.05% I r i s adequate to prevent the formation of"orange peel", the Ericksen t e s t has been carried out on sheets of un-modified and Ir-containing alloy annealed for 15 min a t 720°C. As can beseen in Fig. 3 a d is t inc t "orange peel" pattern develops on the surface ofthe unmodified alloy whereas the Ir-containing alloy gives a sa t i s f ac -tor i ly smooth surface.

For the two alloys that demonstrated a s ignif icant grain refining ef-fect, the one with 0.05% I r and the other with the complex Co-Zr-B addi-tion, attempts have been made to reveal the d is t r ibu t ion of additives inthe alloy by the use of transmission electron microscopy and electronprobe microanalysis. However, none of the techniques has proved to be ef-fective.

GRAIN REFINEMENT DURING SOLIDIFICATION

Possible mechanisms of the effect of minor additions on grain ref ine-ment during so l id i f ica t ion can be summarized as follows:1. Minor additions of highly active elements which may react with theatmosphere, crucible , other minor components of the al loy or with some un-controlled impurities can form minute solid pa r t i c l e s of oxides, carbidesetc . On further cooling, these pa r t i c l e s provide s i t e s for so called"heterogeneous nucleation" of the alloy c rys ta l s , thereby diminishing theextent of supersaturation needed for so l id i f ica t ion and increasing thenumber of growing crys ta ls . This mechanism has been experimentally con-firmed by X-ray of TiC par t ic les in an Al alloy with 0.1% Ti (1).2. If an alloying addition forms with the major component of the alloy aphase diagram with an eutectic (or pe r i t ec t i c ) point s i tuated very closeto the pure major component, even small amounts of the alloying componentgive an alloy of hypereutectic (or hyperperitectic) composition. Inthese alloys a more or less wide temperature range ATp exis ts between theliquidus and eutect ic (per i tec t ic) l ines where primary crystals of theminor alloying component (or i t s compound with the major component) wi l lprecipi ta te . Due to a very low concentration of the minor addition veryfine par t ic les are formed and when the eutect ic (per i tec t ic ) temperatureis reached these precipi ta tes wil l provide centers for heterogeneousnucleation and growth of the c r y s t a l l i t e s of the (essent ia l ly pure) majorcomponent.

An additional factor that accounts for a more effective grain re f ine-ment in alloys in comparison with pure metals (1) i s a concentrationgradient formed in the melt in front of the growing c rys ta l s . Due to adifference in composition between sol id and l iquid phases at a givencrys ta l l iza t ion temperature and a limited diffusion r a t e in the melt, thelayer adjacent to the growing surface will become enriched with the com-ponent that preferably remains in the liquid phase. This compositionalchange means a more or less sharp decrease, in accordance with the l iquidslope, of the equilibrium freezing etemperature of the melt in the layeror, in other words, a retardation of crystal growth a t a given coolingra t e . Due to t h i s , a greater number of potent ial centers of hetero-geneous nucleation i s activated on further cooling and be t t e r grain r e -finement i s achieved. The wider the separation between solidus andliquidus curves and the steeper the liquidus slope, the greater re tarda-tion of growth ra te occurs (2) . The well pronounced dendri t ic segregation

97

observed in the 14 carat jewellery alloy studied (Fig. 1 left) indicatesthat a significant compositional gradient develops during i t s solidifica-tion which enhances grain refinement i f this effect i s produced by minoradditions according to one of the two mechanisms mentioned above. The ob-served behaviour of the alloy with additions of I r , B, Zr, or Co duringsolidification i s in agreement with the corresponding binary phase dia-grams. Compositions of the alloys containing B, ZT and Co fall well intohypoeutectic range so that no precipitation of B, Co or ZrAuj can occurduring solidif ication. Contrary to that , alloys with 0.02 and 0.05% Irare h>pereutectic and grain refinement induced by I r precipitation at thef i rs t stage of solidification can be expected. The enhanced effect ob-served at the higher Ir content i s likely due to a greater numbers of I rprecipitates formed in this alloy. I t can be also assumed that the posi-tive effect of remelting is due to a more uniform Ir distribution achievedby remelting.

RETARDATION OF GRAIN GROWTH DURING ANNEALING

The addi t ives s tudied can be divided i n t o two groups according to t h e i rbehaviour at t h e annealing temperatures :i) iridium that is practically insoluble in the alloy and probably is pre-

sent in the form of fine part ic les formed during crystal l ization,i i ) the rest of additives which at the concentrations used are in the

solid solution. According to theoretical predictions (3), both finelydispersed part icles and solute atoms can affect the migration of grainboundaries and thus restr ict grain growth though by different mechanisms.

The driving force for grain growth in a recrystall ized structure is theexcess of energy associated with grain boundaries and i t is inversely pro-portional to the radius of the boundary curvature. Fine particles causea local increase in boundary length and thus create a drag of boundarymotion. At a given volume fraction f of the particles this effect is in-versely proportional to the average part icle radius T_. Moreover, sincethe average boundary curvature and hence the driving force diminish withthe increase of average grain s ize, at some stage the particles can pre-vent further growth. The upper l imit for the grain size D in the pre-sence of part icles is Dm = 4/3 (r/f) . For the alloy withm0.05 wt.% Ir fis 3 x 10 and the effect of the additive depends on the degree of finessof Ir part icles . If we assume the i r size to be 0.1 ym, the maximum grainsize Dm in this case is about 0.2 mm. It means that precipitates of thisaverage size cannot prevent appreciable grain growth during prolonged an-nealings (as i t indeed occurred in the Ir-containing alloy afte 1 h an-nealing at 750°C), but the retardation of boundary migration they producecould result in the grain refinement that was observed after 15 min an-nealings .

The retarding effect of very small amounts of soluble additives ongrain boundary migration in pure metals is a well established fact ( 4 ) . I tis explained as an effect of solute atoms segregation along grain bounda-ries accompanied by energy gain and described by the following equation:CΓ, = CQ-exp(-E/RT), where C« and CQ are boundary and bulk concentrationsor the solute atoms respectively, E is the interaction potential betweenthe solute atoms and the boundary, T is the annealing temperature. Be-cause the boundary region is characterized by a less regular and lessdense atomic structure, size misfit between the solute atoms and thematrix la t t ice i s the most evident reason for the boundary segregation.I t is likely, however, that other factors such as difference in valencyor type and strength of interatomic bonds between host and foreign atoms

98

can play some role, too, in formation of the foreign atoms atmospherearound grain boundaries. When such a boundary is caused to move, the in-teraction of foreign atoms with the boundary creates a drag and, unlessthe driving force for boundary movement is large enough, the boundary iscompelled to carry the atmosphere of foreign atoms along with i t . In thissituation, the boundary velocity is controlled by the rate at which foreignatoms can diffuse behind the boundary (3,4") . Obviously, in spite of re-tardation of grain growth no upper limit can be imposed by boundary se-gregation of solute atoms on the grain size at long time annealings. Theother difference between effect of finely dispersed particles and that ofboundary segregation is that the lat ter diminishes with the increase ofannealing temperature as a result of both increased diffusion mobility ofimpurities and decreased tendency for grain boundary segregation as fol-lows from Eq. 2.

However, the grain refining effect of soluble additives (or the lack ofit) as observed in this investigation is not consistent with the theore-tical predictions. Of soluble additives introduced in various amounts(0.005-0.1% B, 0.04-0.1% Zr, 0.1-0.3% Co) only Zr gave a moderate grainrefinement whereas B,which represents two extremes of the size misfit inrespect to the alloy lattice, had no effect at a l l . At the same time thelowest amount of boron used (0.005%) can have significant effect, whileadded together with 0.05% Zr and 0.1% Co. (Fig. 2 ) . It is evident thatboundary segregation as described by Eq. 2 cannot account for the effectbecause the total bulk atomic concentration of the additives in this alloyis lower than in many alloys where the same additives were introducedseparately without appreciable grain refining effect. Another observa-tion inconsistent with the effect of the atmosphere of solute atoms at theboundaries is a low rate of grain growth observed at 750°C.

The strong enhancement of grain refining observed when complex alloy-ing by B, Zr and Co is used suggests that there must be an interaction be-tween the additives resulting in a limited mobility of grain boundaries.A.s noted by Losch (5) , chemical interaction between two solutes is mostlikely to occur at grain boundaries where i t is promoted by a higher con-centration of solute atoms resulting from their grain boundary segrega-tion. Because boron forms borides with both Zr and Co, we can expectthat i t is precipitation of boride particles at the grain boundaries thatreduces mobility of the boundaries and prevents extensive grain growth inthe alloy containing additions of B, Zr and Co.

REFERENCES

1. A. Cibula J . I n s t . Metals, 1949, 76_, 321-359.2. W.A. T i l l e r J . Metals, 1959, 11_, 512-514.3. R.W. Cahn "Recovery and R e c r y s t a l l i z a t i o n " , Ch.19 in : "Physical

Metal lurgy", ed. by R.W. Cahn, Elsevier North-Holland Publ. Co,pp. 1129-1197.

4. P. Gordon, R.A. Vandermeyer in " R e c r y s t a l l i z a t i o n , Grain Growth andTextures" , ed. H. Margolin, ASM, 1963, pp . 205-266.

5. W.H.P. I.osch S c r i p t a M e t . , 1977, 11, 889-892.

99

Fie. 1: Effect of additions of 0.05% I r on the as-cast microstructure.

Fig. 2: Effect of small additives on the recrystallized structure an-nealed for 15 min at 750°C. A) unmodified alloy; B) addition c0.05% I r ; C) addition of B + Zr + Co.

Fig. 3; Surface of the alloy (left) and the same alloy with the additionof 0.05% Ir deformed by Ericksen cupping t e s t .

100

REINVESTIGATION OF THE Pr-Ga SYSTEM IN THE 66~TU0 AT % RANGE

J. Pelleg - Ben-Gurion University, Beer-Sheva

D. Dayan, G. Kinunel - Nuclear Research Centre-Negev

INTRODUCTION

The phase diagram of Pr-Ga system has beer, the subject of severalinvestigations(1,3). The existence of the compounds PrGa and PrGa2 inthis system has been reported by all investigators. However the exis-tence of some of the previously compounds, i.e. Pr3Ga(l-2), PrsGa3(2,3),PraGa2(1-2,4) and Pr2Ga(3) and their structure are controversial. Never-theless, the topic of this communication is concerned with the Ga-richside of the Pr-Ga system in the 67-100 at. % Ga range. In previous work(5,6,12) it has been established that some of the RGa2-type compounds(R is a rare earth) have a wide range of homogeneity wMle other not(7).This can be understood on the basis of a model suggested by Pelleg andZevin(6). Similarly, the possible existence of an additional intermetal-lic compound on the gallium rich side has been pointed out(8). This hasbeen confirmed(9) and the structure has been identified as RGa6.

In order to clarify the predictions based on the Pelleg and Zevinmodel(6) regarding solid solubilities of Ga in PrGa2 we undertook astudy using x-ray diffraction, differential thermal analysis, metallo-graphy, SEM and diffusion couples techniques. Particular emphasis wasplaced OTi determining the crystal structure of the intermediate phasesand the evaluation of the change of the lattice parameters of the ephase with compositions within the homogeneity range. (£=Pri-xGa2n+x)).


Fig.l shows the Ga-rich side of the Pr-Ga system derived from theresults of this investigation. Data from reference 1 and from reference 2were included for comparison. It can be seen that PrGa2 shows a ratherwide range of homogeneity which extends to about 78 at. % Ga. The sym-bol e was assigned to this phase. In addition a gallium-rich intermetal-lic compound was detected in this system having the composition PrGa6.The phase n was assigned to it. These two phases are now discussed.

The e phase

PrGa2 is the stoichiometric composition of the e phase which formscongruently at 1470°C(3) and has a hexagonal structure of A£B2 type.Its lattice parameters a-e listed in Table 1. In accordance with otherRGa2 compounds, where R is a light rare earth metal, the solubilityrange of Ga in PrGa2 is broad and extends to about 78 at. % Ga. Thereis also an increase in the lattice parameters with the addition of Ga.

101

An increase in the unit cell is expected according to the Pelleg -Zevin model(6) which is based on pairwise substitution by 6a atomsof the R constituent in the basal plane. Diffusion-couple resultsindicate the presence of a layer between Pr and PrGa6 as seen inFig. 2. In this layer a gradient in concentration of Ga was foundfrom the stoichiometric composition to about 79 at. % Ga as indica-ted by electron microprobe measurements.

The r| phase

PrGa6 forms peritectically at 466±6°C. The formation of PrGa6was established also by diffusion - couple experiments. In Fig. 2such a couple is seen between pure Ga and pure Pr. The predominantphase that forms at a relatively low temperature in a short period(360°C/2h) is PrGae. In agreement with similar RGa6 compounds(5-7,9,10] PrGa6 has a tetragonal structure of the PuGa6 type(ll). Itslattice parameters are presented in Table 1.

One of the authors (J. Pelleg) holds the Samuel Ayrton Chairin Metallurgy.

REFRENCES

1. Tandelli, Gazz. Chim. Ital., 19_, (1949) 70.

2. S.P. Yatsenko, A.A. Semiannikov, B.G. Semenov and K.A. Chuntonov,J. Less-Common Met. 64, (1979) 185.

3. S. Cirafici and Franceschi, J. Less-Common Met., 66, (1979) 137.

4. A.E. Dwight, J.W. Downey and R.A. Conner Jr. Acta Crystallogr.23_, (1967) 860.

5. G. Kimmel, D. Dayan, A. Grill and J. Pelleg, J. Less-Common Met.,75, (1980) 133.

6. J. Pelleg and L. Zevin, J. Less-Common Met., 77_, (1981) 197.

7. J. Pelleg and G. Kimmel, Materials Science and Eng., 52_ (1982) PI

8. R. Manory, J. Pelleg and A. Grill, J. Less-Common Met.,61_, (1978) 293.

9. J. Pelleg, G. Kimmel and D. Dayan, J. Less-Common Met.,81_, (1981) 33.

10. S.E. Hazsko, Trans. Metal. Soc. AIME 221_, (1961) 201.

11. F.H. Ellinger and W.H. Zachariazen, Acta Crystallogr. 19, (1965)281.

12. D. Dayan, U. Atzmony, and M.P. Dariel, J. Less-Common Met.,87, (1982) 87-98.

102

• -This work*-Ref. 2• -Ref. 3

1,0 50 60 70 80 90 100ATOMIC PER CENT GALLIUM

Fig. l . The ga.1 lium-rich side ot the Pr-UaSystem.

103

Table 1

Crystallographic data in the 50-100 a t . % range of the Pr-Ga System

PhaseCompo- Struc- constantssition ture

at.% Ga type a b

4.2817

4.272

4.3021

4.3182

4.3167

6.014

4

4

4

4

4

7

c

.2898

.298

.2924

.3022

.3113

.654

This work

10

This work

This work

This work

This work

PrGa_

RGa,

66.7

75.

78.7

85.7 PuGa,o

Fig. 2. Diffusion couple of praseodymium-galliumshowing the interface layer of the e phaseC?r x Ga ) . Heat treated at 360°C

for 2h. Pr is the grey and Ga is lightgrey layer, x540.

104

TEXTURE IN LOW ALLOYED URANIUM ALLOYS

* * **

J. Sariel , G. Kiminel , and J. Pel leg

* Nuclear Research Center Negev P.O.B. 9001 Beer-Sheva, Israel.** Ben-Gurion University Beer-Sheva, Israel.

INTRODUCTION

Orthorombic uranium is known to exhibit strong anisotropy and its con-sequence might be critical during irradiation growth. The degree of theanisotropy expressed by the dimensional changes in polycrystalline uranium,depends on the grain size and the preferred orientation that is present inthe material due to the prior processing. Thus, randomly oriented finegrained uranium might be almost free of irradiation growth on a macro-scopic scale, due to mutual canceling of this effect, while coarse graineduranium with a strong preferred orientation may change its dimensions ina very anisotropic manner. In practice it is very important to be able toevaluate quantitatively the degree of preferred orientation, and todecrease as far as posible, if not completely to eliminate this preferrredorientation.

In this work two uranium alloys, namely, adjusted uranium and uraniumchromium alloys were chosen to investigate anisotropy, by means of x-raydiffraction. Schulz method diffraction (1,3) and the inverse pole figuretechnique (2,3) were used for texture determination.

EXPERIMENTAL

The composition of the adjusted uranium is 0.02-0.05% Fe, 0.05-0.12% Aland 0.08-0.09% C. It was vacuum melted and cast into 3.6 cm diameter rods.The rods were 3 solution heat treated in a salt bath at 740°C, waterquenched and then vacuum annealed at 540°C. The uranium 0.1% Cr alloy wasprepared in the same way, 3 solution heat treated in a salt bath at 720°C,quenched to a second salt bath at 540°C for isotropic transformation andfinally vacuum annealed at 520

cC. The specimens were slices from the rod,

with the surface of interest perpendicular to the rod axis.

Cu Kα radiation was used for the regular diffraction, and a curved graphitecrystal monochromator was put in front of the proportional detector. Thearea illuminated by the x-ray beam was such that hundreds of grains werecovered,to achieve good statistical sampling. For the Schulz methoddiffraction a texture attachment was used in conjunction with the goniometer.The recorded intensities were sent in parallel to a TTY paper tape punchfor subsequent computerized data processing.

RESULTS

In the Schulz technique the diffracted intensities were corrected forbackground and defocusing, normalized and were devided into six groups.The groups were graded on a logarithmic scale. The relative intensities,P ((j),a), were presented on a stereographic projection, yielding the directpole figure. Here a is the angle of rotation of the specimen about an axis

105

normal to its surface, and <j> is the tilt angle about an axix located in thediffraction plane perpendicular to the goniometer axis. In figure 1 tworepresentative direct pole figures are shown, one of an as cast and thesecond of fully heat treated adjusted uranium. At the bottom of each figurethe relative intensity scale is presented. While the relative intensity3cale for the as cast material spreads up to 19.32, for the fully heattreated specimen, it reaches only 2.96. Similar pole figures were obtainedfor the U-0.1% Cr alloy, indicating a similar reduction in the intensityscale for the fully heat treated specimens.

In the inverse pole figure technique, pole densities Pi were calculatedfrom the peak intensities according to the equation:

I. / 1°.

P i = / ' *

Here n is the number of peaks. I. are the intensities from the specifiedplanes of the specimen and I? are the calculated intensities of a randomlyoriented polycrystalline material. These P^ values were used to constructthe inverse pole figures. Figure 2 shows a quadrant of the inverse polefigure for an as cast and a fully heat treated adjusted uranium. Here alsoa reduction was observed of the pole densities from a rang of 0.27-2.59in the as cast material to a rang of 0.58-1.83 in the heat treatedspecimen. Similar reduction was obtained for the uranium 0.1% Cr alloy.

DISCUSSION AND CONCLUSIONS

B .th, the direct and inverse pole figures characterize the manner ofpreferred orientation qualitatively rather than quantitatively. Furthermore,they are applicable in materials which were subjected to thermo-mechanicaltreatments, resulting in systematic and definite texture. However, ourresults indicate that specimens which were not subjected to such treatment,like as cast, or just heat treated specimens, show non reproducible textureeven if they came from the same batch. This observation can be attributedto the fact that as cast material, lacking any thermo-mechanical treatments,is supposed to be free of preferred orientation, and if there exists somedegree of preferred orientation, it is not systematic and not reproducible.Therefore a technique is suggested to overcome this difficulty. In essence,two parameters calculated form the results of the Schulz method aresuggested to characterize quantitatively the fine non-systematic texturein as cast and annealed materials. For a specimen ideally free of preferredorientation, the normalized relative intensities, P (<l>,cO, should all beequal to one, and thus, the standard deviation of these values is zero.For any other specimen the values of P would spread about the value of one,(although the mean value would still be one) and the standard deviationwould be greater than zero. The higher the maximum value of P (<ji,a) andthe greater the standard deviation are,the degree of preferred orientationis more pronounced. Thus the two parameters proposed for quantitativelycharacterization of preferred orientation are Pmax, the maximum value ofP O,°0 and SD, the standard deviation.

In order to confirm the validity of these two parameters, two extremecases were examined, a sintered silicon, which is expected to have minimalpreferred orientation, and a rolled uranium which has a pronounced texture.Indeed the results for P and SD were confirmed, resulting for the silicon

106

inimininmnniMMiiim.nnmH«i..^^.Mm^L- i

••,•„..'••Li.Miniiiiiiiiin.ii.w.,.•.,•••..„. •»••• :::::::

sssss:::: •KKS".:::ass:..

• •••••••if" >»KiW>> •

— "~----I«***lXI*""«II.M««K"ltM*Hf«M«Mtt"*"-—-***11"**'"** -"""?

wMjnnnnnnnnr mi i mm 11 ii"i*' ' *II tttu.*:V.,',..'.'.',ltMMMM D00000000011»IlIWIIHmi*«<IIlltllHI<...nlimiln.llI.»...,...i.»»«»t+».+»t*iI¥HII4.»» >t

ess xSSSSSXSS

I 1 1 I i •—*• i I mix I I •MΜ I f M H f

i i i ! " - * ; i ***** i I m w I i *M*I i i —•• i

I *" t I'..— 1 j •••»• i t inn 1 1 ww i I w-a !

! • j : " " j j "'•" j j!!!?? i J "*"• I " * * * ;

Fig. 1 : Representative direct pole figures of an as cast (left) and afully heat treated (right) adjusted uranium.

002 023 022 021 041 020 002 023 022 021 041 020

200

0.92 l.?3 1°83 0?81 0779 0.58

1.25

60

0 95

Fig. 2 : Representative inverse pole figures of an as cast (left) and afully heat treated (right) adjusted uranium.

107

in 1.32 and 0.09 and for rolled uranium 11.6 and 1.2 for PmaY and SDlIictA

respectively.P and SD for adjusted uranium observed in this work are 14.3±4.9 andl.?£0.2 respectively for as cast material, and 3.2±0.9 and 0.4±0.1 forthe fully heat treated specimen. These values are statistical averagesfrom a large number of specimens examined. Similar trend was observed inthe U-0.1% Cr alloy, but due to smaller number of specimens investigated,the statistical sampling is not adequate.Hie conclusions of this work are:1. The suggested parameters seem to be good quantitative indicators for

the level of preferred orientation even in the case of as castmaterials.

2. Cast materials have some degree of preferred orientation but this isnon-systematic and non-reproducible.

3. The preferred orientation can be reduced by appropriate heat treatments.

* One of the authors (J. Pelleg) holds the Samuel Ayrton chaii inmetallurgy.

REFERENCES1. L.G. Schulz, "A direct method of determining preferred orientation

of a flat reflection sample using a Geiger counter x-ray spectrometer",J. Appl. Phys. 20 , 1030-1033 (1949).

2. G.B. Harris, "Quantitative measurment of preferred orientation inrolled uranium bars", Phil. Mag. 43_ , 113-123 (1952).

3. J. Sariel, "Texture in low alloyed uranium alloys", M.Sc. thesis inengineering, Ben-Gurion University of the Negev, Beer-Sheva, July,1981.

108

MICROSTRUCTURE AND PROPERTIES OF TUNGSTEN-BASED HEAVY ALLOYS

D V Edmonds

Department of Metallurgy & Science of Materials,University of Oxford, Parks Road, Oxford 0X1 3PH, UK

ABSTRACT

The embrittlement of tungsten heavy alloys of typical nominal composition(wt%) W-5Ni-5Fe and W-7Ni-3Cu by the interfacial segregation of traceimpurity elements has been examined by Auger electron spectroscopy. Theoccurrence of precipitation at tungsten-matrix and tungsten-tungsteninterfaces, and in both the matrix and the tungsten phases, has beenexamined using high-resolution microanalytical transmission electronmicroscopy. The influence of certain processing variables and heat treat-ment is apparent, and the effect of segregation and precipitation onmechanical properties is discussed.

INTRODUCTION

Tungsten-based heavy alloys are manufactured by liquid-phase sinteringtechniques1•2, and after cooling from the sintering temperature the micro-structure consists essentially of a continuous network of approximatelyspheroidal tungsten grains embedded in a Ni-Fe(or Cu)-W matrix binder-phase. A useful combination of mechanical properties is exhibited, butthere is a marked variation in properties, particularly impact resistance,with composition and processing conditions. A better understanding of thechange in microstructure with heat treatment, and of the microstructuraldependence of mechanical properties, formed an objective of a researchprogramme of which this paper outlines the results obtained so far andrefers generally to alloys of typical nominal composition (wt%)90W-5Ni-5Feand 90W-7Ni-3Cu. Parts of the work are published in more detail else-where3"7 .

EXPERIMENTAL PROCEDURES

A detailed account of the experimental techniques employed is given else-where3 '5~8.

EXPERIMENTAL RESULTS AKJ DISCUSSION

Alloy Microstructure

The microstructural appearance of the alloys is illustrated by Fig.l. Itis conventional commercial practice to furnace cool the alloys from the

109

!0Qum

sintering temperature, and this

results in basically similar struc-

tures, although differences in specific

parameters can be identified by

measurement. Table I documents the

matrix volume fraction, average tung-

sten grain size, and contiguity of the

tungsten particle network as a func-

tion of selected treatments. (The

contiguity of the tungsten network,

C^, can be defined9 as the average

fraction of surface area that a given

tungsten particle shares with its

neighbours, and is given by

2N,

Fig.l Micrograph (optical) of as-

sintered furnace-cooled

W-5.2Ni-4.8Fe alloy. (Muddle)

WW

N. + 2N.(1)

WM ' WW

where [%\n| is the average number of

tungsten grain boundaries and Nyjyj is

the average number of tungsten-matrix interphase boundaries that are inter-

cepted per unit length of randomly positioned line).

Table I Measured microstructural parameters

W-5

As

2h

lOh

2h

5h

lOh

W-7

As

lOh5h

lOh

(FC

Specimen

•2Ni-4.8Fe

sintered, ]

@ 1200°C,

@ 1200°C,

@ 1350°C,

@ 1350°C,

@ 1350°C,

.2Ni-2.4Cu

sintered

@ 115O°C,

@ 1350°C,

@ 1350°C,

= furnace

FCWQWQWQWQWQ

WQWQ

WQ

Average

particle

21.5

23.1

23.7

tungsten Volume

size,

± 1.2

± 1.3

± 2.5

fraction

μm of matrix,

30.2

30.0

29.6

25.4

23.8

26.0

cooled, WQ = water quenched)

+ 1.

± 2.

± 3,

± 2,

± 2

± 1

%

,8

.0

.2

.1

.6

.6

Contiguity

tungsten

particles,

0.351 ± 0.

0.358

0.416 ± 0.

0.377

0.415

0.452 ± 0.

0.405 ± 0.

0.404

0.431

0.475 ± 0.

of

cw

023

027

026

021

023

The W-7.2Ni-2.4Cu alloy is representative of a typical commercial prepara-

tion sintered at the lower end of the permissible temperature range. This

has resulted in irregularly-shaped tungsten grains with a relatively high

level of contact between adjacent grains. By comparison, the W-5.2Ni-4.8Fe

alloy, prepared at a higher sintering temperature, contains approximately

spherical tungsten grains with a lower level of contiguity. The average

particle size is comparable in both alloys, but the volume fraction of

matrix is higher in the W-Ni-Fe alloy, and may be attributed to the greater

solubility of tungsten in the Ni-Fe matrix than in the Ni-Cu matrix. High-

temperature solution treatments do not appear to change significantly the

110

average tungsten particle size or the volume fraction of matrix, but dotend to increase the contiguity of the tungsten particle network in bothalloy systems. An increase in Cw is observed with increasing time at agiven temperature or with increase of solution treatment temperature for agiven time.

Figure 2 shows typical transmission electron micrographs of the two alloys,particularly of the tungsten-matrix and tungsten-tungsten boundaries.These boundaries exhibit structural features characteristic of those inother materials, and at least for specimens furnace cooled from the sinter-ing temperature or water quenched from a high solution treatment tempera-ture i 1200°C), there is no evidence of any major microstructural change.The impact resistance of the alloys in these two states can be markedlydifferent, however, and thereby warrants a more detailed examination.

Interfacial Segregation of Trace Elements

For both W-Ni-Fe and W-Ni-Cu alloys a post-sintering heat treatment pro-duced a significant improvement in toughness; for example, the impactenergy is approximately doubled (compared with an as-sintered specimen) bya solution treatment at 1350°C for 1 hour followed by a water quench. Thisbehaviour was reversible; specimens toughened by solution treating andquenching could be re-embrittled by solution treating and furnace cooling.

Figure 3 illustrates a fracture surface typical of both alloys in the as-sintered furnace cooled condition. The fracture is predominantly inter-facial in character and salient features are marked A, B and C on themicrograph; A corresponds to the tungsten side of a ti'ngsten-matrixinterface, B to the matrix side of the same type of interface, and C tothe impingement boundaries between tungsten particles.

Similar freshly exposed fracture surfaces were subjected to analysis byAuger electron spectroscopy using an instrument capable of measuringseparately the surface composition of individual fracture facets. Figure 4shows Auger spectra from the matrix side of the tungsten-matrix interface(i.e. area B in Fig.3), and of significance are. the strong phosphorus, andto a lesser extent sulphur concentrations segregated to this surface.These impurity elements derive from the elemental powders used in themanufacture of the heavy alloys, typical analyses of which are given inTable II, and are expected to lower the interface cohesion and thus

Fig.2 Micrographs (TEM) of W-5.2Ni-4.8Fe alloy," (a) tungsten-matrixinterphase boundary, (b) tungsten-tungsten grain boundary. (Muddle)

I l l

eruph (f=EMl of frac-'f as—sintered

-\: I'.'- .•..J!li-2.4Cu

Fig.4 Auger electron spectra

from fracture surface of as-

sintered furnace-cooled

W-5.2Ni-4.8Fe and W-7.2Ni-2.4Cu

alloys. (Lea and Muddle)

• • - i • I I - L . .

' i r f i . - t ctv

: /l I. . 1 ' ' ' .

'-•.::, . Lnt •_-; t s! ingly , on the corresponding tungsten side of•<:-:'.:*. Lri x interface (area A ) , only a small amount of phosphorus

;, ;iiiii similarly, very little phosphorus was evident on the••-,,•'-• \ .-p. interfaces (area C ) . This association of phosphorusK-kel-rr.ai,rix side of the fractured tungsten-matrix interface can.i->m.-•:is1"i-ated by rastering the incident electron beam across thei selecting the energy of an Auger electron peak respesentative-ui rrr --lenient to form an image. Figure 5 illustrates the simi—'.'t-n. r.h...» phosphorus and nickel images so formed.

Table II Analysis of elemental powders

Composition, wt.

i uv,ier Ca Fe O(ppm)

irbur.v L •

•irbui" v i I

-,! i 00

0

.010

.096-

.012

0.0040.0110.0260.002

0.0020.0030.0040.002

0.011 0.03

99.93

800

30

400

' . I . t

-? 'J .. I.

f rueinr-= ii f

lion i-CLd-r.cnt followed by quenching was found to reduce significantlyp'r.osohorus concentration at the interphase interfaces, but not thehur concentration, and also produced some significant changes in the•'ire mode, as shown by Fig.6. Although interface separation remains,najor component of failure (measured to be > 50pct of the fractureif"1] 'in improvement of interfacial cohesion is evident from the areas:.;*ilf- matrix failure (Di and an increase in the fraction of tungstenis railing by cleavage (E). Thus, it seems that phosphorus is mainlyaside for the measured reversible embrittlement by trace element

wl ly, Auger electron spectra from the facets resulting fromtungsten-tungsten boundary fracture (area C in Fig.3) showed segregationof nic!•-..•= I to levels as high as one monolayer in W-Ni-Cu alloys, but lessin ,v'-:!i-FV. In the W-Ni-Cu alloy segregation of copper to the tungsten-matrix ':-,'jundory during heat treatment was also detected, which would be

112

A-'

Fig.5 Augergraphs of the elements P and Ni on a fracture surface ofW-7.2Ni-2.4Cu alloy. (Lea and Muddle)

expected to lower boundary cohesion. It is known that segregation ofnickel will embrittle polycrystalline tungsten10 and so this and the possi-bility of copper segregation could be expected to contribute to the lowerimpact toughness levels generally found in the W-Ni-Cu alloy system.

Precipitation Reactions

Interface Precipitation

Figure 7 illustrates the fracture surface appearance of a commercially-supplied W-5Ni-5Fe material (of unspecified processing conditions) withparticularly low impact toughness. Fracture has occurred predominantlyby failure of the microstructural interfaces; the tungsten-tungsten bound-ary areas are smooth and featureless, but the tungsten-matrix interfaceareas show evidence of facetting and secondary cracking, suggesting theexistence of a brittle interfacial precipitate. In the present work itproved possible to reproduce this fracture surface appearance by firstgiving a solution treatment ( 1350°C) and quench, and then ageing in thetemperature range 750-850°C. However, there have been other reports of

Fig.6 Micrograph (SEM) offracture surface of W-7.2Ni-2.4Cualloy solution treated 2h at1350°C and water quenched.(Muddle)

Fig.7 Micrograph (SEM) offracture surface of as-sintered furnace-cooledW-5Ni-5Fe alloy. (Jones)

113

similar behaviour, also credited to precipitation, after furnace coolingfrom the sintering temperature11"13.

By deep-etching this fracture surface and exposing it to X-ray analysissome evidence of another phase was obtained, which was tentatively thoughtcould be an intermetallic compound of the form (Ni,Fe)W, isomorphous withthe orthcrhombic intermetallic compound NiW (whicn occurs at approximately75wt%W in the Ni-W binary system11*). Not all the X-ray lines expected forthis phase were detected, however, although a low success rate might beexpected from a rough fracture surface. According to the phase diagram,this phase might be expected to dissolve at temperatures 5 1000°C,and someimprovement in impact toughness values and fracture surface appearance wasobtained after solution treating and quenching from the range 1000-1200°C.Furthermore, Henig et al15 have claimed identification of a (Ni,Fe)Winterfacial precipitate in W-7Ni-3Fe alloys heat treated at < 1000°C.Consequently, there is some evidence to support the possible occurrence ofan interphase intermetallic compound in W-Ni-Fe alloys.

However, more detailed examination has revealed the existence of an alter-native (or additional) precipitation reaction. Firstly, high-resolutionAuger electron spectroscopy performed by Muddle1 6 on the precipitateexposed at the fracture surface shows a significant carbon peak. Theshape of the peak, and its apparent stability to argon-ion sputtering,suggests that the carbon is in the form of a carbide rather than a surfacedeposit. Secondly, thin foil electron microscopy has allowed a closerexamination of the precipitate structure (and morphology).

Figure 8 confirms the presence of a precipitate ^ 0.4ym thick at thstungsten-matrix interface. Similar precipitates have been induced inalloys simulating the composition of The matrix in the W-5Ni-5Fe heavyalloy (38.5Ni-40Fe-21.5W). Convergent beam electron diffraction and X-raypowder analysis of these precipitates have indicated a diamond cubicstructure (with lattice parameter, ao = 11A) and space group Fd3m(no.227)8>16, whilst electron microprobe analysis gave the W:Ni:Fe ratiosas 50.2:20.6:29.2 (at%). An experiment using a laser induced ion massanalyser also detected an additional small carbon peak in the spectrum,

Fig.8 Micrograph (high-voltageTEM) of W-5Ni-5Fe alloy solutiontreated lh at 1350°C, water-quenched and aged lOOh at 850°C.(Posthill).

Fig.9 Micrograph (TEM) ofW-7.5Ni-2.5Cu alloy solutiontreated lh at 1350sC, water-quenched and aged lOOh at 850°C.(Muddle).

114

thought to be real and estimated to result from a precipitate carbon con-tent of approximately 8at%." Both the ternary systems Ni-W-C and Fe-W-Ccontain n-oarbide phases of the forms Ni 6W eC" and ?e6W6C

1B»l9f respec-

tively, both with Fd3m structure and ao = 10.9A. Consequently, it is con-eluded that the embrittling tungsten-matrix interfacial precipitate inW-5Ni-5Fe alloys is an n-carbide of composition (Ni,Fe)6W6C, although the(simultaneous) presence of an intermetallic {Ni,Fe)W phase cannot be ruledout.

Although the W-7Ni-3Cu system has been subjected to much less metallo-graphic examination, the observation of precipitation at the tungsten-matrix interfaces has been made by transmission electron microscopy, asshown by Fig.9. The precipitates have been observed after ageing for lOOhat 850°C specimens previously given a high temperature solution treatmentfollowed by viuenching. The precipitates may occur singly, or as a groupof apparently equiaxed grains. Insufficient analytical metallography hasbeen carried out to identify this precipitate unambiguously at the presenttime. However, it is believed to have a similar embrittling effect tothat found in the W-Ni-Fe system.

Figure 10 shows evidence of precipitation at a tungsten-tungsten boundaryin the W-5Ni-5Fe alloy aged for lh at 1050°C. Convergent beam electrondiffraction identified the structrue as fee (ao = 3.60A), and thin filmmicroanalysis also demonstrated qualitatively high nickel and iron concen-trations in the precipitate. The precipitate thus appears to be similarto the matrix-phase of the W-5Ni-5Fe alloy which is also fee (ao = 3.60&).Evidence for an irrational orientation relationship with either of theadjacent tungsten grains, lying in the region between Kurdjumov-Sachs20

and Nishiyama-Wassermann21,22, has also been ob'tained8. It is concludedthat the precipitate forms to relieve the supersaturations cf nickel andiron in the tungsten grains resulting from liquid-phase sintering, whichhave been reported23 to be approximately O.lat%Ni and l.lat%Fe, respec-tively .

The raised flat areas representing cleavage along the tungsten-tungstengrain boundaries are readily evident on the fracture surface and normally

0-3pm

Fig.10 Micrograph (high-voltageTEM) of as-sintered furnace-cooled W-5Ni-5Fe alloy aged lhat 1050°C. (Posthill)

Fig.11 Micrograph (SEM) of as-sintered furnace-cooledW-3.5Ni-l.5Fe alloy aged lh at1050°C. (Hogwood)

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have a smooth featureless appearance. However, precipitation on theseboundaries is clearly apparent from the cleaved fracture surface, asillustrated in Fig.I?, which follows the work of Hogwood2", who has clearlydemonstrated this effect in W-Ni-Fe alloys heat treated for various timesat 1050°C. Other investigations25'26 have also interpreted similarfeatures on the otherwise smooth fra ;ture facet as probably resulting fromintergranular precipitation, although without experimental identificationof the precipitate.

The appearance of these precipitate features on the tungsten-tungstenfracture surface facets has also been associated with improvedductility25'27. The tungsten-tungsten boundaries are an intrinsicallyweak link in the microstructure; it has been shown that the impactproperties are generally reduced by an increase in contiguity of thestructure (essentially an increase in the tungsten-tungsten grain boundaryarea) and also that segregation of nickel to these boundaries takes place,which should also lower their cohesion. Consequently, any reduction inthe tungsten-tungsten grain boundary area, in this case by intergranularprecipitation of a 'ductile' precipitate with potentially stronger inter-faces, would be expected to result in increased alloy toughness.

Matrix Phase Precipitation (W-Ni-Fe alloy)

Ageing treatments in the approximate temperature range 750-850°C followinga solution treatment and quench result in lamellar and Widmanstatten formsof precipitation in the matrix phase of W-5Ni-5Fe (Fig.12). A 'dendritic'precipitate formed in the matrix of W-Ni-Fe alloys slowly cooled from thesintering temperature has been previously reported12, but not unambiguouslyidentified, whilst slow cooling was also shown to lead to a decrease inlattice parameter of the matrix phase and the appearance of a fine disper-sion of tungsten precipitates26.

During the present study both forms of precipitation were reproduced inalloys made up to simulate the matrix composition, from which they couldbe electrolytically extracted for analysis by X-ray powder diffraction.From the X-ra> results, and from electron diffraction in both the 'matrix'alloys and the' heavy alloy itself, both the lamellar and Widmanstattenprecipitates were identified a tungsten. It was also possible to

Fig,12 Micrographs (TEM) of W-5Ni-5Fe alloy solution treated lh at 1350°C,water quenched and aged 25h at 850°C; (a) lamellar precipitation and(b) Widmanstatten precipitation. (Muddle)

116

"i5

a

Fig.13 Microhardness of thetungsten phase in as-sinteredfurnace-cooled W-5Ni-5Fe alloyrolled 10.1%RA and aged for In.(Posthill)

•0.1 pm-

Fig.14 Micrograph (weak beam TEM)of the tungsten phase in as-sintered furnace-cooled W—5Ni—5Fealloy rolled 10.1%RA and aged forlh at 600°C. Arrows point todefect-dislocation interactions.(Posthill)

associate formation of the lamellar morphology with migration of a high-angle grain boundary, suggesting a typical Type 1 discontinuous precipita-tion reaction from supersaturated matrix, e.g. Ysupers. + Yequl. + W. Anindependent observation of discontinuous precipitation of tungsten in thematrix phase of W-7Ni-3Fe heat-treated in the range 800-1000°C has alsobeen reported recently15.

The present investigation did not attempt to examine the influence ofmatrix-phase decomposition on mechanical properties. Conflicting reportsexist in the literature15>2B'z9, although it should be noted that the pre-viously discussed precipitation reaction at the tungsten-matrix interfacecould occur concurrently, and result in simultaneous embrittlement of thealloy.

Tungsten Phase Precipitation (W-Ni-Fe alloy)

Ageing in the temperature rang,3 500-700°C for lOOh leads to an increase inthe yield and ultimate tensile strength of the W-5Ni-5Fe heavy alloy.Microhardness measurements «n the tungsten particles reflect this increasein the tensile properties of'-the alloy, whilst the hardness of the matrixphase remains unaltered. Prior deformation of the alloy accelerates thisage hardening reaction as «3hown in Fig.13 (as well as giving a work harden-ing increment). Examination of aged specimens by transmission electronmicroscopy, using the weak-beam technique to enhance diffraction contrast,identified a fine dispersion of elastic strain fields •** 20A in diameter(e.g. Fig.14). Analysis of the strain fields suggested that they would becaused by the formation of platelet precipitates on {100} tungsten planes,resulting in a tensile strain field in <100> tungsten directions. Althoughit has not yet proved possible to obtain analytical data from such a smallarea, it ±s proposed that the precipitates form to relieve the supersatura-tion of nickel and iron in the x-ungsten; it has already been shown thatthis can occur at much higher ageing temperatures by precipitation at thetungsten-tungsten grain boundaries.

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CONCLUSIONS

The mechanical properties of tungsten heavy alloys are very dependent onthe interfaces present in the microstructure. Segregation of phosphorusimpurity to the tungsten-matrix interfaces results in embrittlement, butthis can be reversed by a high-temperature solution treatment followed bya quench. Sulphur segregation to the same boundaries, with an expectedloss in cohesion, was also observed. Segregation of nickel to tungsten-tungsten boundaries and the possibility of copper segregation to tungsten-matrix boundaries was similarly detected, and is expected to contributetowards generally lower toughness levels. The relative proportions oftungsten-matrix and tungsten-tungsten boundaries are also shown to varyaccording to processing and heat treatment conditions and should influencealloy properties. The formation of an n-carbide (Ni,Fe)eW6C and possiblyan intermetallic compound (Ni,Fe)W, are shown to occur at the tungsten-matrix interfaces in W-5Ni-5Fe, depending on processing conditions and heattreatment, and can severely embrittle the alloy. Precipitation of an uni-dentified phase at the same interfaces in W-7Ni-3Cu alloys has also beenfound. Precipitation can also be induced at the tungsten-tungsten grainboundaries, but this phase is similar to the heavy alloy matrix binder-phase, and consequently is thougtt could enhance cohesion across theseotherwise relatively weak boundaries. Matrix-phase decomposition was alsoshown to be possible, and resulted in the discontinuous precipitation oftungsten with a lamellar morphology, or in tungsten with an intragranularWidmanstatten morphology. Low-temperature ageing resulted in intragranularprecipitation in the tungsten phase, and this age-hardening response couldbe accelerated by prior deformation.

ACKNOWLEDGEMENTS

I am grateful to Professor Sir Peter Hirsch FRS andProfessor J W Christian FRS for the provision of laboratory facilities. Iacknowledge and thank my colleagues B C Muddle, J B Posthill, P N Jonesand M C Hogwood, both for stimulating discussion and permission to usesome of their published and unpublished results. This research has beencarried out with the support of the Procurement Executive, Ministry ofDefence.

REFERENCES

1. R H Krock, Metals for the Space Age, Proc. Plansee Seminar, 1964,p.257.

2. D J Jones and P Munnery, Powder Metall., 10, 1967, 156.3. D V Edmonds and P N Jones, Met.Trans., lOA, 1979, 289.4. B C Muddle and D V Edmonds, Residuals, Additives and Materials

Properties, Royal Society (London), 1980, p.129.5. B C Muddle and D V Edmonds, Metal Sci., 17, 1983, 209.6. C Lea, B C Muddle and D V Edmonds, Met.Trans., 14A, 1983, 667.7. J B Posthill and D V Edmonds, Intl.Conf. Phase Transformations in

Solids, Crete, 1983.8. J B Posthill, D Pnil Dissertation, University of Oxford, 1983.9. J Gurland, Trans AIME, 212, 1958, 452.10. M P Seah, Acta Met., 28, 1980, 955.11. R V Minakova et al, Second Intl. Powder Metall.Conf. 2, 1966, 91.12. RVMinakova et al, Poro.Metall., 65, 1968, 73.

118

13. Alfennappann, Report NO.R70/71, Chalmers University of Technology,Gothenburg, 1971.

14. J M Walsh and M J Donachie, Met.Trans., 4, 1973, 2854.15. E T Henig, H Hofmann and G Petzow, Proc. 10th Plansee Seminar, 2,

1981, 335.16. B C Muddle, Monash University, Melbourne, unpublished research.17. M L Fiedler and H H Stadelmaier, Z Metallk., 66, 1975, 402.18. Ya S Umanskii and N T Chebotarev, Izv.Akad.Nauk. SSSR, 5, 1951, 24.19. J Leciejewicz, J Less-Common Metals, 7, 1964, 318.20. G Kurdjumov and G Sachs, Z.Physik, 64, 1930, 325.21. Z Nishiyama, Sci.Rep.Tohoku Univ., 23, 1934, 637.22. G Wassermann, Arch.Eisenhuttenw., 6, 1933, 347.23. L Ekbom, Scand.J.Metall., 5, 1976, 179.24. M C Hogwood, R.A.R.D.E., Sevenoaks, unpublished research.25. D G Brandon, E Ariel and J Barta, Proc. 5th Intl.Symp,Electron

Microscopy and Strength of Materials, ed. G Thomas, 1972, p.849.26. P A Verkhovodov et al, Poro.Metall., 196, 1979, 8.27. R Gero and D Chaiat, Materials Engineering Conf., Technion, Haifa,

Israel, 1981, p.46.28. S S Kiparisov et al, Nauchn.Tr.Vses. Naucho-Issled.Proektn.Inst.Tugo,

Met.Tverd.Splanov., 16, 1976, 280.29. L AKonyukhova, Yu A Eiduk and L S Vodop' yanova, Tsvetny Metally.,

No.10, 1974, 57.

119

INVESTIGATION OF THE CREEP FAILURE MECHANISM IN THEMO - 5% W ALLOY

E. Freund, D. Agronov and A. Rosen

Department of Material Engineering Technion I.T.T. Haifa

INTRODUCTION

Previous, investigations have established the fact that creep failure formost metals and alloys is a result of grain boundary cavitation (1,2,3).In an earlier paper (4) by one of the authors it was reported that nearthe fractured surface of a Molybdenum alloy specimen which failed in creepat high temperatures, almost every grain boundary which was normal to thedirection of the applied stress contained a crack. Moreover it was shownin the same paper that thermo-mechanical treatments which effect the sizeand shape of the grains, also influnce creep ductility. For example aslightly deformed specimen by swaging contains large almost equi-axedgrains and therefore a high fraction of the boundaries are normal to theapplied stress, while a heavily swaged specimen is composed of very longand thin grains where most of the grain boundaries are parallel to thedirection of the applied stress. The first specimen failed in creep aftera few percent of elongation, while the latter specimen crept to approxi-mately 50 percent strain before fracture.

The aim of the recent investigation was to understand the kinetics ofcrack nucleation and growth. Since the investigation is not yet completed,we are going to report here only the work done on a slightly swagedspecimen.

EXPERIMENTAL

All the experiments were carried out on a Mo-5%W alloy which was manu-factured by powder compacting, sintering and finally swaging. The reduc-tion by swaging was 18% and the resulting grain diameter was 77 micron,with an aspect ratio of 3,13. Creep experiments were performed in aspecially constructed creep apparatus, in dry Hydrogen, under an appliedstress of 50 MPa and at a temperature of 1200 deg. centigrade. Threecreep experiments were carried out: (1) up to the end of the primaryregion, (2) up to the end of the steady state region, (3) up to fracture.The specimens were removed from the creep apparatus and i ere prepared formicroscopic examination.

Small and sharp notches were machined within the grip section of thespecimens. It was found that when the specimens were broken at -10 deg.centigrade at the location of the notches a large percentage of the frac-ture surface is intercrystalline. From one specimen approximately 10fractured surfaces were created. The broken specimens were examined ina scanning-transmission electron microscope (STEM) using the scanning mode.It was found that even as-received specimens contain very small micro-voids having various shapes and locations, there is a large variation in

120

the early stages of creep and then increases in the tertial region. Therecan be possible reasons for the increasing void density: 1) new cavitiesare created and 2) all the cavities grow, now even the very small oneswhich were not seen before are large enough to be detected.

The area of the average cavity multiplied by the density gives the areafraction of cavities. This value is shown in Fig.6 as a function of creeptime.

DISCUSSION

As mentioned in the introduction, the fact that creep failure is causedby grain boundary cavitation during creep is well known, however thekinetics of cavitation is still not completely understood. Even fromthe few experimental data reported here the following trends can besummarized:

1) Microvoids exist in the as-received material.2) The microvoids are seen mostly in grain boundaries.3) During the first and second stages of creep neither the density

nor the average size of the cavities change significantly.4) The last stage of creep is accompanied and probably caused by

very rapid cavity growth.

It would be difficult to draw final conclusions from partial results andtherefore the above listed four points of conclusion should be treatedonly as indications.

The investigation is now extended to the following directions:

The role of temperature and stress on cavity growth and the effect ofswaging on the phenomenon. Unfortunately, the study require? timeconsuming and rather expensive use of scanning electron microscopy andtherefore we are seeking the critical experiment which will give us thesolution to the problem.

ACKNOWLEDGEMENT

All specimens were manufactured, prepared and heat treated by MetallwerkPlansee, GmbH, Reutte, Tirol. This work is sponsored by Metallwerk Plansee,GmbH and their support is highly appreciated.

REFERENCES

1) Hull, D., Rimmer, D.E., The growth of grain boundary voids under stress,Philosophical Mag. 1959, 4,673.

2) Muller, D., Langdon, T., Independent and sequential cavity growthmechanism. Scripta Metallurgica, 1980, 14,pp J43-148

3) Svenson, L.E., Dunlop, G.H. Growth of intergranular creep cavities.International Metals Review, 1981, 2, pp 109-131.

4) Bendersky, L., Rosen, A. The effect of Thermo-mechanica1 Processingon the creep ductility of Mo-5%W alloy. Material Science and Engi-neering, Vol. 62, no. 2, pp 211-216.

5) Konig, G., Blum, W. Comparison between cell structure produced Inaluminium by cycling and monotonic creep. Acta Metallurgica,1980, 28, pp 519-537.

121

Fig.l. As received specimen. Fig.2. Fractured specimen.

OS 10 1.5 20

OIAHETER OF CAVITIES I pm I

Fig.3. Cumulative size distribution Fig.4. Average void diameterof diameter of cavities. versus time of creep.

S - I S O -

% - 'no

20 10 60 30 0 0 IZO

TIME OF CREEP I HOURS)

20 40 60 80 "CO 120

TIME OF CREEP (hsursl

Fig,5. Mean number of cavitiesversus time of creep.

Fig.6. Area fraction ofcavities versus timeof creep.

122

ISOCHRONOUS CREEP OF COPPER-BERYLLiUM-NICKEL ALLOY(Cu-O.^Be-2-ONi) SOLUTION TREATED AND AGED

Nahum Ni rRafael - A.D.A. Haifa

ABSTRACT

Short-time, constant-load uniaxial creep tests were performed ona high intensity electromagnet coil candidate [1] alloy Cu-0.4Be-2.0Niin the aged [2] 1750HT condition". The test temperatures varied fromroom temperature to 250°C. The stresses used at each temperature were inthe 0.2% yield strength range. These tests were designed to simulateperiodic short-time loading conditions of the coil turns.Isochronousstrength (stress-strain) curves were generated from these creep testresults for time scales of 0.5sec to 8.0sec and compared to the tensile(standard) stress-strain curves of the alloy at these temperatures.It was found that at each temperature the stresses in all stress-straincurves (standard one as well as isochronous) were within <S% of eachother, for a given strain. It was also found that by normalizing thestresses to the 0.2% yield stress at the testing temperature, all curves,independent of temperature and batch, were within 5% of each other (instress, for a given strain). As a result, it is proposed here to use, inpractical design, the lower boundary of the range of the normalizedcurves. This would be a conservative, unique stress-strain curve forthe 175OHT Cu-0.ljBe-2.ONi alloy, independent of small variations of yieldstrength from one batch to another and of temeprature in the rangementioned above.

EXPERIMENT AND ANALYSIS

Test Procedure

Short-time, cyclic creep tests (not reaching steady-state) werecarried-out using MTS tensile test machine and furnace. Most data wasobtained at 150°C. However, a few trend-check tests were also caried outat room temperature, 200cC, and 250°C. Tests were load controlled,monitoring strains with an extensometer and loads with a load cell.Stress levels were reached in 0.5sec, kept constant for lOsec oftransient creep, and unloaded in 0.5sec. Uniaxial specimens were usedafter being heat reated. Each loading cycle was repeated 10 time (untilhaving an apparently repeating curve), with 2-3min between cycles.

Typical creep curves at 150°C are shown in Fig. 1. Isochronousstrength curves were generated from the creep data for creep times of0.5 to 8.0sec, simulating the pulsed creep of the magnet coil duringloading. Simulating coil working conditions, 150°C was chosen to be

solutionize and quench, k0% cold work and *tS2°C age.

123

the major testing temperature, at which several stress levels were applied.At the other temperatures (room temperature, 200°, and 250DC) only twostress levels were tested to check the trend due to change of temperature.

Results

There was little difference between the isochronous strength curvesin the range of 0.5sec to 8.0sec. This can be seen from the two curves,solid line and dashed line, on Fig. 2. These curves were generated fromthe first cycle of each test and are the boundaries of the time scalementioned. It was also found that the same type of curves, whengenerated after a large number of creep cycles, resulted in an evennarrower difference; and all match, within a reasonable accuracy, thelower isochronous curve (of the 0.8sec creep time of the first cycle).

It can also be seen, from Fig. 2, that the scatter band of theisochronous curves is not sensitive to temperature change in the relevantrange. Also, the higher the temperature, the lower the isochronousstrength curve. It is to be noted that both the yield stress (ay) andthe elastic modulus (E) reduce as the temperature increases. Moreover,in our case, the amount of their change with temperature and differentbatches is roughly the same [3] [E(T)/CTY(7)=16Q±10]. Therefore,normalizing the stress with <?y(T) shrinks down even more the scatter bandof isochronous results due to differences in temperature. Practically itgathers all results, independent of temperature and isochronous time, intoone unique strength curve for the material tested here.

Standard tensile (stress-strain) curves (carried out at a loadingrate of 65Opsi per second which is two orders of magnitude slower thanthe loading rates of the short-time creep tests) were also added to thesenormalized-stress versus strain results. These curves were obtained atInesco Inc. in the past, at room temperature, 150°C, and 200°C, andrepresent two batches of Cu-0.4Be-2.0Ni alloy, 175OHT condition: #044and #056. This first batch is the one the isochronous specimens in thiswork were taken from. It had a room temperature yield strength of113.7-Il4ksi. The second batch had a higher room temperature yieldstrength of 119—119.4ksi, representing a realistic variation of strengthfrom one batch to another. The stresses in all cases were normalizedto the measured 0.2% yield strength of each batch at each particulartemperature. The results taken as a whole led to the construction of aunique strength curve for this material as shown in Fig. 3.

Analysis

The unique normalized stress versus strain behavior of theCu-0.4Be-2.0Ni alloy at 1750HT condition was also plotted on a log-logscale using the minimum vahiu of the scatterband of the curve in Fig. 3-These values, together with minimum values of the lower elastic range ofour measured stress versus strain for the alloy, generated a morecomplete, yet conservative curve from0.\% total strain to strains ofabout 8% (and more!)- This curve is shown in Fig. 4. It is to be notedthat the proportional limit in all the measured stress-strain curvesvaried from 0.72av to 0.76ciy. In our conservative curve (Fig. 4) it wastaken to be 0.75tfy for best curve fit.

The temperature and batch independent normalized stress versusstrain curve, shown in Fig. 4, is divided into three parts:

124

1. Linear elastic range-up to 0.75oy(T) (which is the normalizedproportional limit).

2. Transition range from the proportional limit up to about o/ay(T)=l(which is also of the order of about \% total strain).

3. Uniform plastic deformation range, starting at the end of thetransition range (at about 1% total strain).

The uniform plastic deformation range, can be described by thefollowing equation:

—°—=KEn, (1)

ay(T)

where a/oy(T) is the applied stress normalized by the 0.?% yield stress,which is a function of temperature. The total strain at the particularstress level is shown by e. The strain hardening exponent n is found tobe 0.038 and the strength coefficient K is found to be 1.208 for Fig. *».

In the transition range, the curve deviates from Eq. (l). Thisstress deviation is plotted on a log-log scale, against the total strain,and shows a straight line with a slope of -k.M as shown in Fig. 5-

That transition deviation is described as sol lows:

(2)

where A[a/ay(T)] is the deviation of normalized stress from Eq. (l), e hasthe same definition as in Eq. (1) and the constants are measured fromFig. 5 to be m=-k.\2 and A=7.5*10"11.

In reality, the whole non-elastic range of the curve (starting atthe end of the proportional limit) can be described by subtracting Eq. (2)from Eq. (1), to obtain one equation for that range, as follows:

°y(T)

The linear elastic range described here as

°y(T) a "

where E(j)/ay(T) is here taken to be 150, which is the minimum measuredvalue in the range of 150 to 170 measured for the different batches andat different temperatures.

125

CONCLUSIONS

The uniaxial isochronous creep strength of the Cu-O.'iBe-Z.ONi alloy,1750HT condition, was measured by simulating high intensity electromagneticcoil cyclic loading conditions, with hold times of up to 8sec, andtemperatures of up to 250°C. There was almost no effect of the creep onthe strength performance of the alloy that usually is given by a standardtensile stress-strain curve.

By normalizing the applied stresses to the Q.2% yield stress at thegiven temperature, a unique curve was found and mathematically expressedthat describes the complete stress-strain relationship of the alloy.That curve is independent of both temperature and batch of material withinthe working temperature range and specification of material. Its use isrecommended in calculating both short-time creep as well as regular stress-strain behavior since it was generated also from available stress-straindata of that alloy at the 1750HT condition.

ACKNOWLEDGMENT

The auther gratefully acknowledges the assistance of Mr. R. Akin aswell as the discussions with Mr. S.N. Rosenwasser, Mr. R.D. Stevenson andMiss J.E. McGregor. The permission to publish these results granted bythe Inesco Inc. management is also acknowledged.

REFERENCES

TM[1] The Riggatorn '" Tokamak Design, Inesco Inc. 11077 N- Torrey PinesRd. La Jolla, Ca 92037-

[2l Heat Treatment Routin, Developed by Inesco Inc. (1981) .

[31 Inesco Inc. Metals Data File, (I98I).

IJL

2-3 MINUTES.

BETWEEN CYCLES

0.004STRAIN

10 (Sec) - 113.3(K5l)

Fig. I: Representative Cyclic Creep Raw Data.

126

120

110

100

90

SO

70

CO

50

40

30

20

10

0

TIM KUXC WAS

NOT COvEftFD BV

SHORT* Tine CR££P

FOR ISOCHRONOUS

S T O E M T H CURVES

• % a - -

150 (y tt5 (si.) _

ISOCcl 8.0 (»c)

40.6.1C

SOCHROKQUS 03

- 22 (c-JtR.T.)

150 (rt

200 (C1)

250 (rf

(sec)

S

0

e®

B.0 (Stc)

a•aa

i 4 i ft IS fc 14 15 16 l> 18 19 20

Fig. 2: Isochronous Strength Curves.

1.05

100

-.55

i

/T

B•.80

-7J

REGULAR STRES -STRAIN

BATCH *

ROOM TCMP

150 (y

200 fclROOM TEMP.

0 2 % VIELO STRESS

ottom

A

0

113J-IM

SATf: —

0S601

I19-1IS<

SOCHRONOUS 6.5 (5. .1

ROOT TEMP. 9

KOfclO

200(1)0

250ft) S

8.0 (Sic)

a0

as

( * ) All 5p«iIm«M en from l»t(i<^ 0440)73

STRAIN

Fig. 3- Normalized Stress vs. Strain Data from Both RegularTensile Tests and Isochronous Creep.

127

Fig. 4: Normalized Stress vs. Strain Characteristics ofCu-0.4Be-2.0Ni Alloy, 175OHT Condition, Independent ofTemperature and Batch.

03-

QD3-

002

3 « 5 7 10

STRAIN

Fig. 5: The Transition Range Characteristics Shown by theNormalized Stress Deviation from the Plastic RangeEquation Shown in Fig. k.

128

THE HOT TEARING OF LEAD ALLOYS

P. Ari-Gur* and F. Weinberg

Department of Metallurgical Engineering,

University of British-Columbia, Vancouver B.C., Canada

The minimum temperature for which hot tearing occurs, is relatedto the ductile-brittle transition temperature, TQB- Measurements of TQBhave been made in as cast binary Pb alloys containing Sb,Bi,Sn,Ma andCa. The difference between TDB and the solidus temperature was found tobe greatest for the PbSb alloys and least for the PbNa and PbCa alloys.Homogenization markedly changed TQB in the PbSb alloys, indicating thatboth non-equilibrium and equilibrium solute segregation at the grainboundaries contributed to hot tearing. For a certain maximum stressnear TnBr the PbSb and PbCa alloys were much less suseptible to hottearing than PbSb.

INTRODUCTION

In one of the containment procedures being considered for thedisposal of irradiated CANDU fuel, the fuel bundles would be sealed in acorrosion resistant metal container filled with lead1. It is importantthat the lead cast around the fuel bundles will be free of voids orcracks.

During the solidification process most impurities and solutes areconcentrated at the grain boundaries, where the final liquid freezes,lowering the solidification temperature of this liquid. If a tensilestrain develops across the boundary with liquid present, the boundarywill open, producing a hot tear. At present the concentration of soluteat the grain boundary cannot be established quantitatively in a givensystem, nor can the final solidification temperature of the liquid atthe grain boundary2. In some system it is possible to calculate thelocal strains due \o thermal contraction during solidification usingheat transfer models3.

The present investigation was undertaken to measure the temperatureat which the ductile-brittle transition occurs i,t lead and lead alloys.The ductile-brittle transition temperature is a direct indication of thetemperature at which the final liquid solidifies and therefore thelocal temperaturt- above which hot tearing can occur if the material isstrained. Because of shortage in space, only a brief description isgiven, full details are given elsewhere k.

PROCEDURE

The materials investigated were Pb with binary alloy additions ofSb,Sn,Bi,Ca,and Na. Test samples of the alloys listed in Table I wereproduced, in shapes as shown in Fig. 1.

•• Now with Armament Development Authority, Haifa, ISRAEL.

129

The mechanical tests were performed with a table type Instronmachine. During the tests, specimens were kept hot in a well stirredoil bath.

After the failure the cross sectional area at the failure wasmeasured. The appearance of the fractured surface was examined with SEM.

RESULTS

The variation of the reduction in crossectional area with testtemperature for the Pb alloys is presented in Fig. 2. Values of JQQ fromFig. 2 are listed in Table I, as well as the solidus temeprature Tsolidus and the difference AT between TQB and T^jj^g. The values of ATare large for the PbSb alloys, negligible for the PbNa and PbCa alloys.Increased solute concentration results in an increase in AT, with theexception of PbNa alloy.

For the Pb \% Sb alloy, TQB is at the eutectic temperature.Homogenization reduces AT.

Alloy additives significantly increase the strength and decreasethe creep rate of pure lead. cruTS ^ o r ductile samples (close to T D B ) ,are listed in Table I. Na additions have no effect on crm-_, whereasalloying with 1% Sb gives the largest increase from 1.17 to 3.45 MPa.

CONCLUSIONS

Solute additions to lead decrease the ductile-brittle transitiontemperature TQB near the solidus. The temperature interval &T betweenthe solidus and T[JB, during which hot tearing can occur, increases withincreasing solute content. The interval varies appreciably for thedifferent alloys examined. Homogenizing reduces AT indicating that hottearing is associated with both equilibrium and non-equilibrium solutesegragation at the grain boundaries.

The maximum stress in tension of a given lead alloy is not directlyrelated to AT. Accordingly a given maximum stress level can be attainedwith either little probability for hot tearing (PbNa and PbCa) or witha high probability (PbSb).

REFERENCES

[1] D.J. Cameron, J.L. Crosthwaite and K. Nuttall; Can. Met. Quarterly,1983 - in press.

[2] F. Weinberg; Progress in Materials Science 1980, Vol. 30,PP. 295-328.

[31 A. Grill, J.K. Brimacombe and F. Weinberg; Ironmaking andSteelmaking, 1976, No. 1.

[4] P. Ari-Gur and F. Weinberg; Can. Met. Quarterly, 1983 -in press.

TABLE,: DATA FOR LEAD ALLOYS.

wt* 0.5

MaximumStress MPa

1.85

0

SnNa

M

! 2 0.25 0.750.02 0.03

LEAD

15 "06 „,

1 55 1.17

homogenized sample-

131

:••'•••'•:„*!'/;.<",•'/.••/

r

LO

Grip

-Pin

• Sample

Fig. 1: Gripping system for tensile tests.

1

A>

Fig. 2: The change in percent reduction in area with test temperaturefor the Pb alloys shown (weight % concentration).

132

325 310

-o-o-o

PbO 73N3

315 320 320Temperature °C

325 310 315 320

Fig. 2 (b)

80

a<

C 6 0

- PbO 02 Ca PbO 03 Ca

I 1330 320 325 330

Tempera tu re C

PbO5SbH

31U 315

Fig. 2 (c)

133

THE EMBRITTLEMENT OF STEELS BY LOW MELTING POINT METALS

Dr. Norman N. BreyerIllinois Institute of Technology

In the late 1930's it was found that lead in small quantities, about1/4% by weight, added during the teeming or ingots would improve themachinability of steel. It was determined, after extensive testing,that the mechanical properties of low carbon unalloyed steels wereunaffected by the lead addition1. Improved production rates usingautomatic and semiautomatic screw machines resulted in lower unit partcosts. The introduction of lead to alloy steels had been initiated inthe post World War II period with the machining experience showingsimilar cost improvements. Unfortunately, the leaded alloy steelsoccasionally failed, both during processing and during subsequentservice^>3. In some cases, the unexpected failures were catastrophicand occurred after extensive and expensive machining operations toproduce the part. Figure l i s a complex qear from which the teeth onone of the helical iears spalled during the heat treatment. The sametype of violent failure can also occur in service under similar stress-temperature conditions. Several service failures which have come tolight recently include jet engine failures and extensive cracking ofdies in forging production operations^.

Extensive research has shown that the reason for such failures lies inthe action of lead as a "liquid" metal embrittler in high strenathsteels, a phenomenon which has been found to occur in many other metal-metal systems. The embrittlement, which can start, in some cases,110°C (200°F) below the melting point, will exhibit a trough with aminimum just above or at the melting point has been called metal inducedembrittlement (MIE), with that portion occurring up to the melting pointlabeled solid metal induced embrittlement, SMIE^. At temperatures abovethe melting point the liquid metal induced embrittlement (LMIE) is amaximum and eventually will decrease with a return to ductility, theso-called recovery temperature. (See Figure 2) Extensive testing hasrevealed that the steel can fail in a brittle manner if threeconditions coexist:

a) The presence of lead.b) Tensile loading.c) Temperatures between 200°C (400°F) and 480°C (900°F).

The lead need not be present internally in the steel to cause theembrittlement; externally applied lead has also been found to causebrittle failures. Many other low melting metals, including Cd, Zn, Snand In, will embrittle steel below the melting point; all of thosetested revealing a rapid decrease in RA to a minimum at the meltingpoint^. Additions of tin, antimony or zinc as dilute second elementsin the lead and bismuth or antimony in the tin markedly increase theextent of embrittlement. As an example, the alloying with tin of theexternally applied lead results in increasing severity of embrittlement.(See Figure 3) There was not a consistent trend in the SMIE or onsetregion. On the other hand, a systematic change in the brittle-to-ductile recovery temperature was found to accompany changes in thecomposition, with increasing tin alloyed with the lead progressivelyincreasing the transition temperature.

134

Although the embrittlement is a kinetic phenomenon with the rate ofarrival of the low melting metal atoms to the crack tip controllingthe extent of loss in ductility, the effect of changing the testingstrain rate is not consistent with a strict rate of arrival effect.Figure 4 shows the results of testing AISI 4145 leaded steel specimensover the temperature range at which an embrittlement trough was found.Using three orders of strain rate change it was found that the samegeneral shape of the RA trough was produced with some inconsistentresults. As would be expected, the shorter test times accompanyinghigh strain rates shifted the onset (SMIE) region to highertemperatures, consistent with shorter exposure times. What wasunexpected was the finding that high strain rates shifted the recoverytemperature to higher values, a revelation inconsistent with shorterexposure times. (The complicating effects of strain rate on theproperties of the base metal, steel, should play a part in theeventual answer.) The same atoms which cause embrittlement had beenfound to be responsible for other embrittlement phenomena includingtempered martensite embrittlement and temper embrittlement3.

Cold working the steel surprisingly tends to alleviate sensitivity tothe metal induced embrittlement (MIE)8. (See Figure 5) This is truein spite of the fact that, cold work increases the strength level ofthe steel, an effect which normally increases the sensitivity toembrittlement. The fracture characteristics reflect the increasedresistance to fracture.

REFERENCES

1. "Properties and Machinability of a Leaded Steel", T.J. Dolan andB.R. Price, Metals and Alloys, (January 1940) p. 20.

2. "The Effect on Lead on the Mechanical Properties of 4145 Steel",S. Mostovoy and N.N. Breyer, Trans. ASM, 61 (1968), pp. 219-232.

3. "Some Effects of Certain Trace Elements on the Properties of HighStrength Steels", N. Breyer, Proceedings of the 31st ElectricFurnace Conference, (1973), pp. 183-189.

4. "Lead Induced Brittle Failures of High Strength Steels", N. Breyerand P. Gordon, Proceeding of the Third International Conference onthe Strength of Metals and Alloys, (August 1"73), Cambridge,England.

5. "Solid Metal-Induced Embrittlement of Steel", J.C. Lynn, W.R. Warkeand P. Gordon, Material Science and Engineering, 18 (1975),pp. 51-62.

6. "Liquid Metal Embrittlement of 4145 Steel of Lead-Tin and Lead-Antimony A l loys" , N.N. Breyer and K.L. Johnson, Journal of Testingand Evaluation, 2, No. 6, 1974, pp. 471-477.

7. "Environmental Sensi t iv i ty of Structural Metals: Some DynamicAspects of Liquid Metal Embrittlement", K.L. Johnson, N.N. Breyerand J.W. Dal ly, Proceedings of Conference; Environmental Degrada-t ion of Engineering Materials, Virginia Polytechnic Ins t i t u te ,1977, pp. 91-103.

8. "Effect of Cold Work on Liquid Metal Embrittlement by Lead Alloyson 4145 Stee l " , M, Watkins and K.J. Johnson and N.N. Breyer, IVInteramerican Conference on Materials Technology, June 29-July 4,1975, Caracas, Venezuela.

135

Fig. 3.

70 -

40

50

s 4 0

o1 3 0

3 20

10

Reduction of area, RA,for 200-ksi (1379-M Pa)4145 steel surfacewetted with variousPb-Sn alloys as afunction of testtemperatures.

O H.T.

A 10% C.W.

O 20% C.W.

O 30% C.W.

O 50% C.W.

O 200 400 600 800 1000

TEST TEMPERATURE, °F

Fig. 5. Ducti l i ty Properties ofInternally Leaded 4145Steel Processed to 200 ksiNominal UTS by Heat Treat-ment Alone and by HeatTreatment Plus 10, 20,30, and 50% Reductionsby Die Drawing.

400 500 600 700

Temperoture, °F

136

Fig. 1. Complex Leaded Alloy SteelHelical Gear with SpalledTeeth Resulting from HeatTreatment.

Fig. 2. SchematicRepresentation ofDuctility Trough

oG

Unembrittled».

\ Ductility\ trough

/ Recovery. - J — — * temperature

Test Temperature

120

o: 80 -

uiu£40

TEST TEMPERATURE, "C100 200 300 400 500

£

' ' KEY€ IN/IN/WIN

.0025~ .025

.232.46

1 1

I

M/M/SEC.0000417.000417.00383.041

1 f 1

' • '

i _

t

200 400 600 800TEST TEMPERATURE, *F

Fig. 4. Effect of Strain Rateon Ductility-TemperatureProfiles for Lead Steel.

1000

137

THE EFFECT OF SUBSTRUCTURE ON CREEP PROPERTIESOF THE T Z M ALLOY

D. Agronov, E. Freund and A. RosenDepartment of Materials Engineering

Technion, Israel Institute of Technology, Haifa

INTRODUCTION

Previous investigation with Mo-5%W alloy has revealed that the creepstrength is strongly dependent on the thermo-mechanical history of thealloy (1,2). It was found that swagging at high temperatures introducesa subgrain structure and the higher the percentage of reduction by swag-ging the greater is the average dislocation density in the subgrain bound-aries and the higher is the creep strength of the alloy.

Mo-5%W is a solid solution alloy and therefore the only possible hardeningmechanism is the one which stems from dislocation rearrangements. Thestudy reported here involves an entirely different moly alloy, TZM, whichis strengthened by dispersion (3,4,5). It is well known that TZM is muchstronger than the moly-tungsten alloy and its creep resistance is muchhigher (6). However, little is known about the effect of high temperaturepre-deformation, such as swagging on its creep resistance. The purpose ofthis investigation was to find out whether the creep strength of thisalloy is also affected by thermo-mechanical treatment and to compare theresults with that of Mo-5%W obtained during the previous investigation.

EXPERIMENTAL

The specimens were made from powder compacted, sintered and swaged TZM,with the following chemical composition in weight percentage: 0.5% Ti,0.07% Zr and 0.01-0.05% C. Five series of specimens were prepared byMetallwerk Plansee. The difference between the series was in the finalreduction by swagging after the last recrystallisation treatment. Thefinal reductions were 18.6, 36.0, 50.2 and 75.0 percent. The fifth serieswas received in the recrystallised condition. The complete productionschedule of the specimens as well as the various grain sizes in thelongitudinal and transverse direction are given in detail elsewhere (7).It is to be noted that the grain size of TZM is considerably smpller thanthe grain size of Mo-5%W.

All creep experiments were carried out in a high temperature creep appa-ratus and the procedure of testing are explained in detail in (8). Testswere performed at the temperature range of 1000-1300°C and in the stressrange of 70 - 450 MPa. Specimens for TEM were prepared from the as-received as well as from the crept specimens. Thin films were obtainedfrom both transverse and longitudinal directions and were examined in aJeoi lOOCX scanning transmission electron microscope.

RESULTS

Fig.1,2 and 3 exhibit the variation of the steady state creep rate (SSCR)with stress for various temperatures for the recrystallised, for the least

138

(18%) and for the most severely swaged conditions (75%), respectively.Similar behavior was observed for the other batches of specimens, i.e.swagged to 36% and 50% reduction. The main features of these diagrams aresummarized below. There are enough data points at 1300°C to indicatethat they can not be connected by a single straight line. Instead, twostraight lines can well represent the results. The line which connectsthe high-stress results is very steep (high n-va.lue), while the linewhich connects the lower stress results has a much smaller slope. Thistendency is very well known and was published by several investigators(9,10,11). It is difficult to determine the exact value of the slope forthe lower stresses since the number of experiments is small in this range,The variation of the n-values with temperature or the amount of swaggingis not systematic and therefore we assumed that the slope n is independ-ent of these factors. The average slope n = 13.6 which is very high. Onthe other hand, the n value of dispersion hardened alloys is known to bevery high (12,13). The practical implementation of the high n value isthat the stress range of the measurable creep rates is narrow. Thispoint will be discussed later in detail.

The recrystallised alloy has the lowest creep resistance at everytemperature. The creep strength of the swaged alloy increases withincreasing amount of swagsing. At low stresses however, where the creeprate is small the degree of reduction by swagging has almost no effect onSSCR. Similar behavior was observed for the Mo-5%W alloy (14).

The subgrain size for all four types of specimens in the as-received con-dition is listed below. The values were obtained by measuring a largenumber of samples.

amount of reductionby swagging (%)

average subgraindiameter (ym)

18.6

0.74

36

0.88

50.2

1.03

75

0.76

The subgrain size 6Q is practically the same for all the specimens andits average value is 6O = 0.85 pm. The subgrain size of Mo-5%W wasalso independent of the amount of swagging, however &o of Mo-5%W was2.1 p about three times larger than 6O of TZM.

The Disorientation between the subgrains was measured by the Kikuchiline shift method, which was developed for the Mo-5%W alloy (8). Fig.4represents the cumulative distribution function of the subgrain mis-orientation, 7(0)), versus the subgrain Disorientation angle, a). It isto be noted that the distribution shown in Fig.4 is very similar to thatobtained for the Mo-5%W alloy,(8). For both alloys the percentage ofsmall misorientation angles decreases with increasing reduction byswagging and the average misorientation angle increases. The distribut-ion curves for the 211 and 212 TZM alloys (50 and 75% reduction) arepractically identical. The distribution curves of the 111, 112 and122 Mo-5%W alloy (50,75 and 85% xeduction) were also similar. This factindicates that swagging ov-z 30% reduction in both alloys could not bevery effective in improving the creep strength.

139

We have also started to measure the subgrain size and the distributionof misorientation angles after creep, however at the time of preparationof this report we have only partial results.

DISCUSSION

Since the study is not completed, we do not want to draw final conclu-sions. Nevertheless, the results which were obtained until now indicatethat the creep mechanism in the dispersion hardened TZM and in thesolution hardened Mo-5% alloy is probably the same. Accordingly theidea of effective subgrain size can be applied also for TZM. A detailedexplanation of the theory is published elsewhere (8). One of the impor-tant factors, according to the theory, is the strengthening effect ofthe subgrain boundaries and the importance of the subgrain size.Undoubtedly, the very high creep strength of swaggedTZM compared to thatof the swaged Mo-5%W alloy is due to its finer subgrain size. We calcu-lated the factor of strengthening only due to the finer subgrain size inTZM and obtained a two order of magnitude increase in creep rate. Onthe other hand experimental results show a four order of magnitudeincrease in creep rate (1200°C, 140 MPa) ( 7 ) . The discrepancy betweenthe calculated and measured values is due to the contribution ofdispersion or precipitation. We believe that when the investigation iscompleter1 it will be possible to evaluate this contribution quantita-tively.

Another important observation of the recent study is the very high stressdependence of the SSCR in the high stress region. This phenomenon istypical for dxspersion hardened alloys and has been reported in theliterature (12,13). For TZM the average value, independent of tempera-ture and swagging conditions, was n = 13.6, while in Mo-5%W alloy n=8.5.As mentioned before, due to this very high slope the stress range of themeasurable creep rate is small. For example, at 1300°C under the stressof 160 MPa, the 212 (50.2% reduction) specimen creeps at a rate of 10~^hr-1, while under the stress of 270 MPa the creep rate is 10"1 hr~l.A 40% reduction of the load results in a three order of magnitude dropof the creep rate. A further reduction of the load causes creep atextremely slow rates, in any case below the sensitivity of our measure-ments .

The investigation is now conducted in the following directions:(i) the effect of creep deformation on the distribution of misorientationangles (a>), (ii) quantitative measurements of the free dislocation lengthin subgrain boundaries (A) and the relationship between oj and X andfinally (iii) the interaction between dislocations and dispersoids orprecipitates during creep.

ACKNOWLEDGEMENTS

All specimens were manufactured, prepared and heat treated by MetallwerkPlansee, GmbH, Reutte, Tirol. This work is sponsored by MetallwerkPlansee, GmbH, and their support is highly appreciated.

140

REFERENCES

1. Bendersky, L., Komem, Y., and Rosen, A., High Temperatures - HighPressures, 13, 511-520 (1981).

2. Bendersky, L., Rosen, A., and Mukherjee, A.K., Scr. Met., 16,467-470 (1982).

3. Chang, W.H., Trans, of the ASM, 218, 254-256 (1960).

4. Chang, W.H., Trans, of the ASM, 56_, 107-124 (1963).

5. Wilcox, B.A., and Gilbert, A. Acta Met., l^, 601-606 (1967).

6. Eck, R., Planseeberichte fur pulvermetallurgie, 27^ 53-74 (1979).

7. Agronov, D., Bendersky, L., and Rosen, A. To be

published in International Journal of Refractory & Hard Metals.

8. Bendersky, L., Dr. of Science Thesis, Haifa (1982).

9. Weertman, J., Trans, of the ASM, 6^, 681-694 (1968).

10. Sherby, O.D., Burke, P.M. Progress in Materials Science, Vol. 13,340-350, (1967), Pergamon Press, Oxford.

11. Mukherjee, A.K., Bird, J.E., Dorn, J.E., Trans, of the ASM, 62^155-178 (1969).

12. Lagneborg, R., Bergman, B., Metal Science, January, 20-28 (1976).

13. Gibeling, J.C., Nix, W.D. Materials Science and Engineering, 45_,123-135 (1980).

14. Bendersl.y, L., Rosen, A., and Mukherjee, A.K. Stress and Micro-structure Dependence of the Creep Resistance of Mo-5%W Alloy, inStrength of Metals and Alloys., ed. by R.C. Gifkins (ICSMA 6,Melbourne), 585-600 (1982).

141

Fig.l. Steady state creep rateversus stress for TZM inrecrystallised condition.

Fig.2. Steady state creep rateversus stress for TZM in swaggedcondition.

Fig.3. Steady state creep rateversus stress for TZM in swaggedcondition.

Fig.4. The cumulative distributionfunction F(o)) versus disorientationangle u) for TZM in as receivedcondition.

142

FRACTURE TOUGHNESS EVALUATION OFBRITTLE MATERIALS USING INDENTATION METHOD

Z. Nisenholz

A.D.A. HAIFA

A simple and economic technique for the evaluation of fracturetoughness (Kjc) for brittle materials has been applied to several brittlematerials.

The technique consisted of a two step procedure:(i) Introduction of a radial crack pattern into the surface by means

of a Vickers indenter and(ii) measurement of the crack length under a microscope.

The materials tested by this method were:Alumina based ceramics, ZnS, WC-Co, and Zirconia-Ytria. The results werein good agreement with those obtained by other methods.

1. INTRODUCTION

During indentation with a sharp indenter, two orthogonal radialcracks evolve within the elastic/plastic field under the indenter. Thefield contributes 2 components to the driving force on the crack system:an elastic (reversible) and a residual (irreversible) component [1-6].

At the indentation surface the elastic stress is compresive whilethe residual stress is tensile and therefore the radial cracks grow totheir final length due to the residual driving force as the ir.denter isunloaded. This force is suitably characterized by the residual stressintensity factor (3):

Kr=xrP/C3/2 (I)

r (2)

where

P - Indentation Load

c - crack length

A - material independent constant

I|J - the characteristic half angle of indenter (7A° in Vickers indenter)

E - Young's Modulus

H - Vickers hardness

143

3- RESULTS

Table 1 summerises the results for the various materials, underdifferent test conditions (load and environment).

Table 1: results under different test conditions.

Material

AD-85(A.R.)

ii

ZnS(H.P.)

II

II

II

AD-85(H.T.)

II

II

Environment

ai r

oil

air

oil

air

oil

ai r

air

air

air

Indentation Load(N)

50

50

2

2

5

5

10

50

70

100

1 ' l(MPa-m*)

3.0+0.11

3.2+0.40

0.56±0.06

0.62+0.02

0.64±0.06

0.67+0.02

0.67+0.10

4.19±0.32

4.22+0.22

4.30+0.17

Average indentation toughness results for the materials which havebeen studied in this work, are listed in Table 2, with results obtainedfor these materials using other toughness measurements.

Table 2:

Material

AD-85(A.R.)

AD-85(H.T.)

Al20a-10^Ti02

ZnS (H.P.)

ZnS(C.V.D.)

WC-6£Co

ZrO2-4£Y2O3

Fracture Toughness

E(GPa)

220

220

300

95

102

686

250

H(GPa) KTr

9

11

17

1

1

17

12

i1

5

0

1

8

9

5

11

Results—^_—_—____——~—(indentation)

(MPa-m*)

3.0

4.2

3.46

0.6

1.04

12.3

6.0

K

2

1

lc

3

.0

12.

5.6

(Other Method)

(MPa-m*)

(D.C.B.)1

-

-

-

(D.T.)2

0 (D.T.)

(C.N.)3

144

At post indentation equilibrium conditions: C=C0 and K =K and weobtain the basic equation for evaluating material toughness:

2 / / 2 (3)

The average calibration factor according to Lawn et al [3] andAnstis et al [5] for Viclcers indenter is:

A(cotip) 2/3=Kc(H/E)4Co/2/P=0.015

The fracture toughness value can be calculated directly, knowingthe indentation load (P) and the crack length (c):

K=0.0l5(E/H)Vc3/2

2. EXPERIMENTAL

The fracture toughness of 5 different brittle materials wasdetermined using the post indentation method.

1. Al203-(AD-85)

2. Ai2o3- iorno2

3- ZnS

k. WC-63XO

5- 4

All specimens were shaped to a form with flat parallel surfaces.One of the surfaces was polished with diamond paste. The minimum sizesof specimen surface area and thickness were 7mm2 and 2mm respectively.5 measurements or more were made on each specimen. Effect of atmosphereon crack growth was examined by comparison of results from specimens withand without protection. A drop of immersion oil was placed on theindentation site of the following specimen: AD-85 as recieved (A.R.),heat treated AD-85 (H.T.) and Hot Pressed (H.P.) ZnS.

Part of the AD-85 specimens were studied after the following heattreatment: rapid cooling from T»50oC, stress relief at 800°C andreheating to 1150°C. This treatment was found to improve significantlytheir toughness.

Several different indentation loads were applied on thesespecimens.

145

1. Double Cantilever Beam.2. Double Torsion.3. Chevron Notch.

4. DISCUSSION

The oil immersed specimens produced slightly higher toughnessresults than those obtained from specimens which were exposed to air.This difference is insignificant considering the higher standarddeviation of the results.

The indentation toughness values seem to be independent ofindentation load. This load independence was already confirmed by Anstiset al [5] for other materials.

The similarity between the indentation toughness results and theresults obtained by other toughness measurements as listed in table 2,leads to the conclusion that this relatively simple fracture toughnessmeasurement method can and should replace the other methods for K.measuring of brittle materials. c

5. REFERENCES

[1] A.G. Evans, T.R. Wilshaw, Acta Metall, 24(10), 939[2] M.W. Swain, J. Mater Sci, H_(12), 2345 1X976).[3l B.R. Lawn et al, J. Amer. Cer. Soc., 63(9-10), 574 (I980).[4] B.R. Lawn et al, J. Aust. Cer. Soc. j|Jl) 4, (I980).[51 G.R. Anstis et al, J. Amer. Cer. Soc. 64(9), 533 (1981).[6] D.B. Marshall, B.R. Lawn, J. Mater. Sci. 14(8), 2001 (1979),

146

FAILURE OF WELDED INCONEL-600 PIPE IN THE COOLING SYSTEMS OF A NUCLEAR

REACTOR

G. Kohn, B. Herrmann, A. Stern, E. Rabinovitz, and S. Addess

Nuclear Research Centre - Negev, P.O.Box 9001 Beer-Sheva, ISRAEL

ABSTRACT

Serious leaks were detected in the inlet and outlet pipes of the heatexchanger in the primary cooling loop of the nuclear reactor at the NRCN.Non-destructive tests were conducted which included: ultrasonic testing,tests with dye penetrants and radiography. The flawed part was replacedand mechanical tests were performed on it. The crack areas of theInconel 600 tube were examined using optical and scanning electronmicroscopy. Chemical analysis of both cracked and intact tubes werecarried out. It is concluded that stress-corrosion cracking was the maincause of failure, while minor evidence of fatigue was encountered as well.Measures for the pre\'ention of similar failure in the future are suggested.

INTRODUCTION

Routine inspection of the cooling system of the nuclear reactor atNRCN detected serious leaks of heavy water both in the inlet and outletpipes of the heat exchangers in the primary cooling loop. A schematicdrawing of the defective area is shown in Figure 1.

The upper part of the outlet pipe which contained the cracks wassectioned out and replaced by a new pipe and flange. The aim of thispaper is to describe the tests carried out in order to locate thecracks, determine their cause and suggest measures to avoid theirfuture occurrence.

EXPERIMENTAL

Upon dismantling the heat exchanger it was submitted to. a series ofnon destructive tests (NDT) which included ultrasonic testing, dyepenetrants and radiography. Two cracks were detected on the outlet pipe:one was a circumferential crack parallel to weld No. 1 and at a distanceof about 15 mm from it. The length of the crack was about 310 mm on theinside and about 5 mm on the outside of the pipe. A second crack about20 mm long was detected near weld No. 2.

Further destructive tests were conducted on the part of the pipewhich was sectioned out and replaced by a new pipe. Chemical analysisof the pipe material has been carried out both on samples near the crack

147

and away from it. Mechanical testing was conducted both on specimens fromareas near the cracked region and away from it. The latter ones weretaken both from parts of the pipe containing a weld and parts withoutany weld. Tensile specimens were machined according to ASTM E-8 (1) andthe tests were performed on an Instron machine.

Curved bending specimens were machined from areas containing thelongitudinal weld and from unwelded areas. Vickers micro-hardness testshave been carried out at a load of 10 gm on a part containing the longi-tudinal weld. The microstructure of the pipe material has been studiedusing both conventional optical microscopy and scanning electronmicroscopy (SEM). Fracture surfaces of the tensile specimens werecompared to the surface of the cracked area and the different modes offracture were established.

RESULTS

The results of the chemical analysis of the pipe material establishedthat the composition of the steel was in accordance with regularspecifications of Inconel 600 (Table 1) .

Table No. 1: Chemical Composition of Inconel 600

Sample Concentration in w/oNi Cr Fe Mn Si Cu

Specification: >72.0 14-17 6-10 <1 <0.5

At crack: 74.6 14.15 5.9 0.18 0.2

Away from crack: 76.5 14.8 6.3 0.20 0.2

0.5 <0.08

0.015 0.08

0.009 0.08

Both tensile and hardness tests proved that the properties of the pipeaway from the circular weld were as expected of Inconel 600 (Table 2),and the bending tests indicated that welding did not have an adverseeffect on the pipe material.

Table No. 2: Tensile Properties of Inconel 600 Specimens

Specimen

Unwelded

Unwelded

Welded

Welded

Yield Strength0.2% (kg/mm 2)

28.5

25.4

46.5

34.3

Tensile Strength(kg/mm 2)

65.3

61.1

66.1

63.6

Elongation

(%)

33.5

34.7

28.7

33.0

Typical values from tht literature are yield strength of 25.3 kg/mm 2and Tensile Strength of 63.3 kg/mm 2.

148

Optical micrograph of the longitudinal deffect-free weld is shownin Figure 2 while Figure 3 shows a series of micrographs taken from thecracked area of the pipe. From the first picture it is clear that theweld area is surrounded by a small heat affected zone (HAZ) beyond whichno grain growth can be observed. On the other hand, from the secondpicture it is evident that while welding did not affect much the grainsize on the upper part of the weld (the heavy flange side), it had astrong effect on the lower part, containing the crack in which largegrains are clearly seen. The intergranular nature of the crack is alsoevident in this picture.

figure 4 shows a fracture surface of a deffect-free sample typicalboth to welded and unwelded areas. The mode of fracture which is dimplerupture is typical of ducticle material with transgranular tearing. Frac-ture surfaces of the cracked pipe showed a completely different mode offracture.

Figure 5 shows a SEM micrograph taken from the crack area. It isclearly seen that the mode of fracture is intergranular with little orno signs of tearing. Al rich layers (as proved by EDAX measurement),can be seen on the intergranular surfaces, while Cr rich particlescontaining up to 3 w/o sulfur are also evident on the grain boundaries.Some evidence of fatigue crack propagation was evident as seen in Figure 6.

DISCUSSION

The results of the tests carried out on the pipe material indicatedthat away from the circular weld the properties of the Inconel were asexpected from a sound and ductile material. The micrographs presentedshow that the crack propagated in an intergranular mode from the insideof the pipe towards the outside. Intergranular fracture is the typicalmode of failure of Inconel steels when submitted to stress corrosion(2-4). Stresses on the pipe in this cooling system originated probablyfrom the weld and from external loads such as the weight of the pipeitself.

The large difference in grain size between the two sides of thecircular weld are probably due to the different thickness of the materialon both sides of the weld which lead to different heat extraction capacity.The larger grains had Al rich layers on their boundaries which originatedprobably from alumina which is known to be dissolved in the water. Ithas been previously established that alumina can absorb large quantitiesof chlorides and increase their local concentration (5). This, plus thefact that the Cr rich particles found at the crack area were also richwith sulfur may indicate the presence of a corrosive environment leadingto crack propagation.

CONCLUSIONS

It has been concluded that stress corrosion cracking was the main-cause of failure. Since Inconel 600 is not very susceptible to stresscorrosion cracking and since the time to failure was very long it wasrecommended that the external tension loads on the pipe system be made assmall as possible with maximum elimination of mechanical vibrations anddistortion. Heat treating complete parts to eliminate internal stresseswas not practical due to the large size of the components and was

149

therefore not recommended. It was also recommended to carry outperiodical NDT to uncover cracks before they become critical.

REFERENCES

ANSI/ASTM E8-77a, in 1977 Annual Book of ASTM Standards, part 10,American Society for Testing and Materials, Philadelphia, Pa,, 1977,p. 154.S. H. Bush and R. L. Dillon, in Steess Corrosion Cracking andHydrogen Embrittlement of Iron Base Alloys, R. W. Staehle,J. Hochmann, R. D. McCright, and J. E. Slater, eds., NACE, Houston(1977), p. 61.J. R. Cels, Corrosion, 34, 198-209 (1978)J. Blanchet, H. Coriou, L. Grail, C. Mahieu, C. Otter, andG. Turluer, p. 1149 in Ref. 2.W. E. Berry, Corrosion in Nuclear Applications, John Wiley, New-York,1971, p . 184.

flange

No. 1

weld No. 2

Fig. 1 Inconel-600 Outlet Pipe

Fig. 2 Micrograph Containing the Longitudinal Weld

150

Cb) Ca)

(c)

n Cβ)

Fig. 3 A Series of Micrographs from the Cracked Area O60)

151

Fig. 4 Fracture Surface (x625)

Ca)

(xl35) (xlllO)

Fig. 5 Crack area (a] Intergranular fracture (b) Solid layers

Fig. 6 Crack Area Showing Fatigue Marks (xl670)

152

EFFECTS OF METALLURGICAL VARIABLES ON HYDROGENEM3RITTLEMENT TN TYPES 3 l 6 , 321 and

STAINLESS STEELS

P . Rozenak and D. E l l e z e r

Department of Materials EngineeringBen-Gurion University of the Mepev,

Beer Sheva, Israel

SUMMARY

Hydrogen embrittlement of 3l6, 321 and 3^7 types austeniticstainless steels has been studied "by charging thin tensilespecimens with hydrogen through cathodic polarization.through-out this study we have compared solution annealed sampleshaving various prior f.usteniiip /rra i • -sii.c with Rfinples giventhe additional sensitization treatment. The results shov thatrefinriu 'rains i~-r roves the reni •--tance to hydrogen crackingregardless of the failure node. The sensitized specimenswere predominantly intergranular, while the annealed speci-mens show massive regions of microvoid coalescence producingductile rupture. 3^7 type stainless steel is much moresusceptible to hydrogen embrittlement than 321 type steel,and 316 type is the most resistant to hydrogen embrittlement.The practical implication of the experimental conclusionsare discussed.

1. INTRODUCTION

The occurrence of hydrogen embrittlement in austenitic stain-less steels is substantial with ductility losses, and bythe appearance of nonductile fracture surfaces [1-6]. Thedeleterious effect on mechanical properties caused by hydro-gen charging has been found to depend strongly on metallurgi-cal factors [7-11].

The object of the study was to evaluate the effect of grainsize and heat treatment on the hydrogen susceptibility of3l6, 321 and 3^7 types austenitic stainless steels.

2. EXPERIMENTAL PROCEDURE

Commercial austenitic stainless steels of types 3l6, 321 and3^7 having the compositions shown in Table I were used forthe present study. Various austenitic grain-sizes, wereobtained following various times at austenizing temperatures.The grains size as measured by ASTM E-112 method were 7, 9and 11 ASTM in 3l6 stainless steel, 8, 9 and 13 ASTM in 321stainless steel and 7, 9 and 13 ACTM in 3^7 stainless steel.

153

Table 1

Chemical composition of AISI type 3l6, 321 and 3^7 austeniticstainless steels

Element

CrHiMnS iCMoT iNb

Amount ( wtstainlessAISItype 316

17.9812.091.710.8790.052.08

-

.%) of elementsteels

AIEItype 321

18.0710.111.321.0330.072

O.Ul-

in the following

AISItype 3^7

17.909 . 0 01.511.0250.061|

__

0 .6 l

Some samples of each group were given a further senitizationheat treatment at 65O C for 2.X hr. The samples were tensiletested at room temperature at an extension rate of 0.005 cmmin" while undergoing cathodic polarization. The hydrogencharging cell contained IN H?S0, solution with 0.25 grl ofNaAsO . A platinum counter electrode and a current densityof 50 mA/cm were used. For comparison purposes, specimenswere tensile tested in air at room temperature. After failurethe fracture surfaces were examined with a scanning electronmicroscope (SEM). Microstructure from the various startingconditions were characterized using transmission electronmicroscopy (TEM).


Sensitization for 2k hr. at 650 C produced discontinuousgrain boundary carbides in 3l6 stainless steel (Fig. la)which varied in size from 100 2 to 1(00 2. Diffraction patternsfrom these particles were readily indexed as M__C,-. Exanina-tion of the microstructures of sensitized both 321 and 3^7types revealed grain boundary carbides (M _C/-) and a densedistribution of (MC) carbides which precipitated on matrixdislocations (Fig. lb, c).

The results given in Fig. 2 and Table II demonstrate theinfluence of grains size and heat treatment on the mechanicalproperties of 316, 321 and 3^7 types austenitic stainlesssteels under cathodic polarization. A significant featureof the results is that decreasing prior-austenite grain-sizeincreases the mechanical properties and the resistance of3l6, 321 and 3^7 steels to hydrogen embrittlement. Thedependence of the tensile properties on the grains size is

154

shown for annealed and sensitized 316, 321 and 3!'T typessteels in Fig. 2. V it-Id s«.r.. r (fnt of samples with grain siz-.A5TM 9 were decreased ai i.ut 15f , as compared to "! ii e grainedsize (ASTM 11, 13) in sensitized 31o, 321 and 3^7 ;./ne steels,but were unaffected whether the material wa;; ur.neai ed orsensitized in various grain-size groups, ii'.h •••_•; i. ••-•+, >o Mi-tensile prootrties, ujdvogaa h .. •- '<•"•.•;"/ IU-A-KV •! --ffect on theultimate tensile strength. UTS or fine grained size (AGTM 11)decreased 35^ as coirmared to coarse grained (ACTM T) 3l6type sensitized specimens and h2% reduction of elongation.As can be seen from Fig. 2 and Table Ii, the sensitized coarsegrained 3^7 type steel wMch is the most susceot ible steel,resulted in 7855 reduction of elongation and 31* reduction inultimate strength while the sensitized coarse grained 3?1type steel resulted in 75? reduction of elongation and ?8%reduction of ultimate strength. Most resistant were sensitizedcoarse grained 316 type specimens that resulted in ^k% reduc-tion of elongation and 2?% reduction of ultimate strength.However, the total reduction of elongation at fracture ascompared to those of uncharged specimens is reduced about60% in fine grained sairnles whether the material was sensiti-zed or not.

Following testing, the fracture surface of each specimen wasexamined to determine the modes of failure. The fractures ofthe specimens tested while cathodically charged show consi-derable differences between heat treatments and prior auste-nite grain-size range. Ductile dimpled rupture wit}: raicro-void coalescence was the main feature mode in sensitizedfine grained (ASTM 11) 3l6 type steel (Fig. 3(a)) withnarrow inter granular zones of about 5 pra on both fracturesurface sides. A completely brittle fracture was observedin sensitized coarse grained (ASTM 7) 3l6 type (Fig. 3(b),(c)). The initial hydrogen fracture is mostly intergranularwith only small transgranular cleavage-like areas .

U. CONCLUSIONS

( l) Refined grain-size improves the resistance to hydrogencracking regardless of the failure node.

(2) Examination of the mi crostructures of beth 321 and 3!t7types using TEM revealed discontinuous grain boundary carbides(Mp_C^) and a dense distribution of ('1C; carbides with preci-pitates on matrix dislocation. It is suggested that the impro-ved ductility in sensitized 3l6, 321 and 3^7 steels stronglydepends on grain size and carbide morphology.

(3) Examination of the fracture surfaces of coarse grainedspecimens tested while cathodically charged shot.- considerabledifferences between the annealed and the sensitized specimens.3l6 type sensitized specimens were nrodominantly intergranu-lar, while the annealed specimens show massive region ofmicrovoid coalpscencr T-,ror1uc i r.p, dunt""le runturn.

155

(I t ) 3I17 t y p e s t a i n l e s s s t e e l i s much more s u s c e p t i b l e t ohydrogen e m b r i t t l e m e n t t h a n 321 t y p e s t e e l , and 3 l 6 t y p e i st h e most r e s i s t a n t t o hydrogen e m b r i t t l e m e n t .

REFERENCES

1. M.L. H o l z w o r t h , C o r r o s i o n 26 ( 1 9 6 9 ) 107.2 . M.B . Whiteman and A.R. T r o i a n o , C o r r o s i o n 21 ( 1 9 6 5 ) 53-3 . M.R. Lou than , J r . , Hydrogen i n M e t a l s , I .M. B e r n s t e i n

and A.W. Thompson, e d s . ( A . S . M . , Meta l s P a r k , Ohio , 197*0

5 3 .k. A.W. Thompson, Hydrogen in M e t a l s , I .M. B e r n s t e i n and

A.W. Thompson, e d s . ( A . S . M . , M e t a l s P a r k , O h i o , 197*0 9 1 -5. H. Hanninen and T. H a k k a r a i n e n , C o r r o s i o n 36 ( l 9 8 0 ) hi.6 . D. E l i e z e r , D.G. C h a k r a p a n i , C. A l t s t e t t e r and E.N. Pugh,

Met. T r a n s . lOA (1979) 935-7 . C.L. B r i a n t , Hydrogen E f f e c t s i n M e t a l s , I . M . B e r n s t e i n

and A.W. Thompson, e d s . (TMS-AIME, P e n n s y l v a n i a , 1950)527-

8 . H. Hannir .en, ^. Hakkarp.ipen ; n d P . " e n o n e n , HydrogenE f f e c t s i n M e t a l s , Met. Soc . of AIME, I98O, p . 575 .

9. I .M. B e r n s t e i n and A.'-/ . Thompson, e l s . v ''MS-AIME,P e n n s y l v a n i a , 1980) S75.

1 0 . C.L. B r i a n t and A.M. R i t t e r , S c r i p t a Met. 13 (1979) 177 .1 1 . C.L. B r i a n t , Met . T r a n s . PA ( 1 9 7 8 ) 7 3 1 .

TA3.LE II

Percentage reduction in tensi le properties o** charged speci-mens compared with those of uncharged specimens for AISI type3l6, 321 and 3*»7 stainless s teels

Type ofstainlesssteel

AISIAISIAISIAISIAISIAISIAISIAISIAISIAISIAISIAISIAISIAISIAISIAISIAISIAISI

3l631632132134731*7316316321321

31*73l*73163163213213't7

Grain

ASTMASTMASTMASTMASTMA f-THASTMASTMASTMASTMASTMASTMASTMASTMASTMASTMASTMASTM

size

111113131313Q

99999778877

Heattreatment-

AnnealedSensitizedAnnealedSensitizedAnnealedSensitizedAnnealedSensi tizedAnnealedSensitizedAnnealedSensitizedAnnealedSensitizedAnnealedHens'tizedAnnealedSens it i zed

Reduction(#) in the<• ensile propertiesYield Ultimatestrength tensile

k20

505030

10

n1500n•-)

c0

strength11q

21161618O

15152021152

2?2U2S2331

following

Elongation

5 5526857676359

65576355k65hC>3755878

156

Fig. 1. TEM micrographs showingcarbides in the sensitized AISItype 316, 321 and 3^7 stainlesssteels: (a) grain boundary carbidesin the sensitized AISI type 316stainless steel; (l>) precipitationof carbides within the grains, ondislocations and at grain boundariesin the sensitized AISI type 321 stain-less steel; (c) precipitation of car-bides in the sensitized AISI type 3^7stainless steel. (c)

ANNEALED - -

Alii 318« — AlSt 347 A 5 T M

Aisi m

1

Jib Asm lap

"fiP i

lgr"

SENSITIZED

* ^ ^ ^ * ~ - • " — •

— AISI 31BAISI MT

4 — .1 ^ 9 »STi• ASTH T

V113 ASTH 11

g

(a) DISPLACEMENT (nun) (b) DISPLACEMENT imml

Fig. 2. Engineering stress vs. displacement curves illustrat-ing a comparison between (a) annealed and (b) sensitized AISItype 316 (grain sizes, ASTM 7,9 and 11) (-.-), 321 (grainsizes, ASTM 8,9 and 13) (—) and 3U7 (graineizes, ASTM 7,9 and13) (- - -) stainless steels tensile tested while undergoingcathoaic charging.

157

Fig. 3. Micrographs of fracture surface of AISI type 3l6stainless steel, grain sizes(a) ASTM 11 and (b), (c) ASTM 7S

tensile tested while undergoing cathodic charging after asensitization heat treatment.

158

MARTENSITIC TRANSFORMATION IN 3O4L and 316L TYPESSTAINLESS STEELS CATHODICALLY HYDROGEN CHARGED

E. Minkovitz and D. E l i e z e r

Department of Mater ia l s Engineering,Ben - Gurion Univers i ty of the Negev,

Beer-Sheva, I s r a e l

ABSTRACT

This paper reports a TEM study on the role of phase transitions at thecrack tip in 304L and 316L Types Stainless Steels cathodically hydrogencharged in the absence of any eternally applied forces. The possiblerole of o-'and e martensite phases in the fracture mechanism is discussed.

1. INTRODUCTION

Hydrogen induced martensite phase transformation in austenitic stainlesssteels have been extensively studied with respect to the relative sta-bility of the austenite p(f.c.c) phase, which in fact, is character-ized by its tendency to transform to martensite on cooling or duringplastic deformation below the critical M_ temperature. The effect ofhydrogen on the γ-phase stability is that hydrogen decreases the y-phasestability and muy induce transformation of the y-phase to a" and emartensite (1-4) . It has been shown that most steels that form or'--martensite are quite susceptible to hydrogen cracking, however, thisissue has become controversial due to the fact that embrittlement tookplace when no martensite has been formed (3). This paper, based partiallyon recent work (5,6,9), demonstrates TEM study on the formation andpropogation of microcracks caused by hydrogen charging, in connectionwith (a';e ) martensite phases. The possible role of a'and e martensitephases in the fracture mechanism is discussed.

2. EXPERIMENTAL PROCEDURE

The stuuies were carried out on 304L and 316L types austenitic stain-less steels. The steels were of commercial grade, and were receivedin the form of sheets 0.2 mm thick. All of the samples used in theseexperiments were first solution annealed for 1 h at 1100°C and thenwater-quenched. Specimens suitable for electron microscopy were thenprepared by electrolytic polishing at 65 V in a Tenupol polishingcell using 30 cm perchloric acid, 300 cm methanol and 520 cm br.canolsolution at -18 C. In the electrolytic polishing process, an attackof the hole edge is normally unavoidable. Thus, in order to avrid un-certainties about the origin of the crack and the possibility of obser-ving cracks which were formed by the thinning process, only specimenshaving "perfect" hole edges (Fig. la) were selected to be suitable forthe hydrogen charging process. After a TEM examination, where the holeedge was checked carefully to ensure that no deformation structure had

159

been induced during the thinning process the specimens were cathodically

hydrogen charged and then were investigated again by TEM. The hydrogen

charging was performed in a charging cell, at room temperature in^the

absence of an external force in a IN H-SO. solution with 0.25 gl" of

NaAsO- added as hydrogen recombination poison. A platinum counter

electrode and a current density of 0.5 A cm were used. The charging

time was 15 min. TEM analysis was carried out in a JEOL-200B electron

microscope operating at 150 kV.


A TEM micrograph in the "Mesh-Image" magnification showing the TEM

specimen's hole after the thinning process is represented in Fig. la.

Forming cathodic charging to the same specimen in the absence of any

externally applied stresses for a few minutes, revealed mainly inter-

granular cracks and some transgranular cracks at the hole edge of the

TEM specimen. Further charging and re-examination by TEM analysis re-

vealed crack propagation. Fig. lb represents cracks at the hole.edge

which were formed following the charging conditions of 0,5 Acm for 2 h.

Dark field electron micrograph of a transgranular crack in 304L,

following charging conditions of 0.5 Acm" for 15 min. is shown in Fig.

2a and 2b. The Selected Area Diffraction Pattern (SADP) at the crack

tip and along the crack surfaces is represented in Fig. 2c. These

findings have revealed Debye rings reflections from a fine grained bcc

ft'-martensite and a fee )'-matrix. The bcc phase and the surrounding

a'ly interface showed a completely fine grained polycrystalline

structure, while the SADP taken at about 0.5 ^m from the a'/y interface

showed single crystal fee y-phase, indicating that the formation of the

polycrystalline structure for the p-phase was highly localized to that

area. Further charging to the same specimen and re-examination by TEM

studies have revealed crack propogation through the martensite phase in

front of the crack tip. The bright field electron micrograph of a

transgranular crack and the microstructure around the crack tip in 316L

Type following the charging conditions of 0.5 Acm" for 15 min. is

shown in Fig. 3a. SADP taken at the crack tip is represented in Fig. 4a

and its schematic diagrams are given in Figs. 4b and 4c. These findings

have revealed single crystal spot patterns of (fee) )>-austenite and

(hep)t-martensite (Fig. 4b) where the v/e orientation relationships are

(lll)v II (0002)

e and [01i]

y || [1120]

f (6) . The SAD at the crack tip also

showed Debye rings characteristic of a very fine-grained polycrystalline

structure, which were indexed as (bcc) a' reflections (Fig. 4c). Fig. 3b

represents a dark field electron micrograph of the e-martensite plates,

taken from a (OllO) reflection spot of e-martensite. The location of

this phase is in the front of the crack tip. A dark dield image (Fig. 3c)

using the (211) a' ring intensities showed that the a'-phase is located

ahead of the crack tip within the c-phase. Further charging to the same

specimen revealed crack propogation through the mixed area of a' and e

-martensite phases, however, mainly through the (hep) e-martensite (5).

No evidence of the appearance of a'-martensite phase after hydrogen

charging was found within the graiis; in the thicker region of the

specimen as previously reported (6). It has been shown that the resist-

ance to hydrogen embrittlement is improved by increasing the μ-phase

stability (7,8). Transformation of the )>-phase to the a'(bcc) phase in

front of the crack tip has been shown in Type 304, fractured in 10 Pa

160

of H gas (7) or when Type 304L was cathodically hydrogen charged in the

absence of any externally applied stresses (9), where the fracture occurs

through the o'-phase. However, recently an HVEM observation of crack

tips of Type 304 sheet specimens which were severly embrittled by hydro-

gen pre-charging and thr.n fractured in air have shown that the crack

propagated mainly along the e-martensite and partly in the region

having a mixed structure of a'and e-martensite phases (10). In the

case of Type 316L, the significant observation of the present experi-

ments is that the y to e -phase transition occurs in front of the crack

tip, where the crack propagation occurs mainly through the e-phase.

Furthermore, it has been reported (11) that Type 310 steel was embritt-

led by hydrogen and that the crack proceeded along the interface

between austenite and the hydrogen induced e-phase. However, crack

propogation in 304L Type, occurs through the a'-phase in front of the

crack tip.

4. CONCLUSIONS

This paper reports a TEM study on the formation and propogation of

microcracks caused by hydrogen charging in connection with (a';e)

martensite phases. The presence of α-bcc martensite in front of the

crack tip of both steels support the view that embrittlement of unstable

steel can have the form of an auto-catalitic process. Crack initiation

occurs through triaxial stresses caused by hydrogen penetration and

evidently followed by the formation of a1bcc martensite when hydrogen

egresses the specimen. The existence of the bcc phase in front of the

crack tip escalates hydrogen entry and crack propogation which in turn

aid martensite formation. However, in the case of 316L Type, it is

now recognized that the presence of the e-phase may be quite important

in the fracture mechanisms. Thus, it is suggested that increasing the

f-phase stability increases the relative importance of the e-phase in

the fracture mechanisms.

5. REFERENCES

1. J.M. Rigsbee and R.B. Benson, J. Mater. Sci. 12, 406 (1977).

2. D. Eliezer, D.G. Chakrapani, C.J. Alstetter and E.N. Pugh, Met.

Trans. A, lOA, 935 (1979).

3. C.L. Briant, "Hydrogen Effects in Metals" I.M. Bernstein and A.W.

Thompson, eds., The Metallurgical Society of AIME, New York (1981),

527.

4. N. Narita, C.J. Altstetter and H.K. Birnbaum, Met. Trans. A., 13A,

1355 (1982).

5. E. Minkovitz and D. Eliezer, Scripta Met. 16, 981 (1982).

6. E, Minkovitz, M. Talianker and D. Eliezer, J. Mater. Sci. 16, 3506

(1981) .

7. N. Narita and H.K. Birnbaum, Scripta Met. 14, 1355 (1980).

8. E. Minkovitz and D. Eliezer, J. Mater. Sci. 17, 3165 (1982).

9. E. Minkovitz and D. Eliezer, J. Mater. Sci. Letters 1, 192 (1982).

10. T. Nakayama and M. Takano, Corrosion-NACE 38, 1 (1982).

11. A. Inoue, Y. Hosoya and T. Masumoto, Trans. Iron Steel Inst. Japan19, 170 (1979).

161

Fig. 1. TEM micrographs in the "Mesh-Image" magnification of the TEMspecimen.(a) After the thinning process.(b) Cracks which were formed by cathodic charging at the

hole edge of the TEM specimen.

Fig. 2. Dark field electron micrographs and Selected Area DiffractionPattern (SADP) of 304L Stainless Steel, which was cathodicallyhydrogen charged.(a) Dark field, taken in (200) a' (bcc) ring. Note that the

martensite phase is located along the crack surfaces.(b) Dark field, taken in (200") a1 (bcc) ring. Note that the

martensite phase is located in front of the crack tip.(c) SADP taken along the crack surfaces and the crack tip,

which was indexed as y(fcc) matrix and a'(bcc) rjartensite.

162

Fig. 3. TEM micrographs showing a transgranular crack and themicrostructure around the crack tip of Type 316L whichwas cathodically hydrogen charged.(a) Bright field image.(b) Dark field image. Note that the e-martensite phase is

located in front of the crack t i p .(c) Dark field image. Note that the a'-martensite phase is

located ahead of the crack t ip within the e-plates.

200

042 01109

10111101

242 200

O [012]z.A. - v

[ ] 2 .A . - hcpf

110

211022310123

bcc

Fig. 4 (a] Selected area diffraction patterns of Type 316L taksnat the crack t ip shown in Fig. 3.

(b) Schematic diagram of electron diffraction pattern ofp-austenite and e-martensite.

(c) Schematic diagram of the Debye rings of a'-martensite.

163

MECHANICAL PROPERTIES DEGRADATION OF HYDROGENATED AUSTENITIC SS

I. Gilad, Y. Katz and H. Mathias

Nuclear Research Center-Negev, POB 9001 Beer-Sheva, Israel.

INTRODUCTIONAustenitic stainless steels (ASS) are often preferred as structural mate-rials, due to their beneficial combination of improved mechanical proper-ties and aggressive environment resistance. Clearly, metallurgical effectsmight be significant in ASS related to phase stability, segregation andsensitization processes. This turns to be particularly important by consi-dering forming or thermo-mechanical processes. Plastic deformation belowthe MJ temperature causes martensitic phase transformation following thereaction: y •* E' + a' (y, e1 and a' are the FCC austenitic phase, HCPand BCC martensitic phases, respectively). The influence of y decompositionon ASS properties are not well established. As proposed by Jewett andothers (1)» ductility might be selected as a possible criterion to evaluatethe susceptibility of ASS to hydrogen effects. On the other hand, ductilityvariation with temperature, in air, are non-regular with minimum valuesabove 900 K (2). In spite of this ductility reduction, necking occurs duringuniaxial tensile tests resulting in a ductile fracture mode.

ASS uniaxial tensile tests during or after hydrogenation show higher yieldstrength, lower ultimate strength and ductility loss associated with alter-native brittle fracture modes (3,4). This behaviour is mainly observed inAISI 304L ASS hydrogenated by high fugacity methods. For example, cathodiccharging provides hydrogen concentration up to 34 \ at a depth of about1 urn (5).

The present paper is aimed to track hydrogen effects due to gas chargingmethods which result in relatively low hydrogen concentration. More of amapping concept is adopted, intended to spot possible sensitivity regionsalong the temperature scale. For comparative purposes, effects with andwithout hydrogen were investigated by means of mechanical testing, X-Raydiffraction technique and fractographical study.

EXPERIMENTAL PROCEDUREStandard tensile specimens with the axis parallel to the rolling direction,were machined from commercial 304L and 316L ASS plates, with a diameterof 7.14 mm and a gauge length of 50 mm. Tensile tests were performed inan inert atmosphere at stable temperatures from 77 K up to 1050 K, utiliz-ing a nominal strain rate of 0.06 min"1. Precharging of the specimens wasperformed in dry hydrogen at 800 atm and 300 K for one month, or 500 K forone day. At these conditions, the theoretical hydrogen concentration nearthe surface is 0.56 a-s, while the average concentration of 0.085 a-i wasactually measured.

164

In order to perform X-ray diffraction of the mechanically tested specimens,

samples of the fractured specimens were cut along their axis by me^.s of

spark errosion and then mechanically and electrolytically polished.

Reflection X-Ray profiles were obtained from the uniformly strained region,

utilizing Mo Kα radiation in conjunction with a graphite monochromator.

SEM was applied for the examination of the necked region and of fracture

surfaces of the strained specimens.


The temperature dependencies of the mechanical properties for the AISI

304L, before and after hydrogenation, are illustrated in Figs.l and 2.

As shown in Fig. 1, the strength of the tested 304L drops monotonically,

but with varying rates at different temperature ranges. It can be seen

that hydrogenation caused no significant changes of the strength, except

a slight increase of the yield strength between 200 K and 300 K. Referring

to Fig. 2, the total strain to fracture, e-j-jis given as the sum of the

uniform strain, eu, and the strain at the neck, e

n. Figure 2 demonstrates

that the ductility of unhydrogenated 304L is strongly temperature dependent.

Values of eu reach a maximum of 65% at 270 K, and a minimum value of 15%

near 1000 K.

Figure 2 clearly indicates that the same trends of ductility changes have

been obtained for hydrogenated 504L, with the exception that a loss of

ductility occurred between 200 K and 300 K. Notice also that at 77 K, a

higher value of EU has been obtained after hydrogenation. It is well

observed that for hydrogenated and unhydrogenated materials, the reduction

of area (RA) curves are very similar to the corresponding en curves but

not to E curves. This is not surprising, taking into account that both

RA and eu curves correspond to the plastic deformation at the necked region.

Testing of several specimens at 3C0 K and at a strain rate of 0.006 min

showed a better ductility, but no change in the strength, whether the

specimens were hydrogenated or not.

The mechanical behaviour of the 316L ASS was similar to that of the 304L,

but with less pronounced hydrogen effects.

Valume fractions, Xa, of strain induced mariensitic BCC a' phase, were

evaluated from X-ray diffraction profiles. Curves of Xa vs. test temperature

are shown in Fig. 3, for hydrogenated and unhydrogenated 304L and 316L

ASS. In order to allow a comparison between Xa and e

u, the corresponding

eu vs. T curves are included. There are indications that the presence of

hydrogen influences the formation of the strain induced a1 phase, especially

for the 316L. In spite of this, approximately the same M^ temperature,

namely 300 K, was obtained for the present testing conditions. It should

be mentioned that in case of as-recieved 304L and 316 ASS, Caskey (4)

obtained different M^ values, as indicated by arrows in Fig. 3.

Fractography after strain to failure revealed a ductile fracture mode

characterized by dimples. The size of the dimples increased with the

temperature, but was usually smaller in the outer circumferential region

as compared to the central region. Hydrogenation resulted in a change of

the fracture mode at the outer region, only. Instead of dimples (Fig. 4a),

facets formation and transgranular cracking occurred (Fig. 4b). Such a

brittle fracture mode was obtained by straining at temperatures between

240 K and 300 K. The complex state of stress which developes during strain-

ing at the outer periphery of the reand tensile specimens, plays certainly

165

an important role in the occurrance of the brittle fracture mode.

Generally, it is well recognized that hydrogen may concentrate at pref-erential sites during charging and redistribute in the course of plasticdeformation. Local hydrogen concentration after cathodic charging may inducephase transformations and cracking at the surface, even without an externalapplied stress (5). After gas charging the hydrogen concentration is supposedto be low. But during straining at specific temperatures, the hydrogenconcentration at discontinuities, could exceed a critical value needed tocause damage. As shown, 300 K was the highest temperature at which hydrogenaffected the mechanical properties, and corresponds to the measured M^.No hydrogen degradation has been observed below 200 K, but straining at thistemperature range resulted in extensive formation of a' martensite.

In the curent study, mechanical properties degradation due to hydrogenationhas been observed under the following conditions: 1) the presence of biaxialstate of stress, 2) low austenitic stability which enabled induced marten-sitic transformation to occur and finally 3) the existence of sufficiantthermal drive at which the permeability is high in order to reach thecritical hydrogen concentration at preferential sites. In fact, these con-dition really existed in the relative strongly hydrogen affected rangebetween 200 K to 300 K.

Referring back to the mentioned proposal of considering ductility loss ascriterion for hydrogen susceptibility, an additional remark should be men-tioned briefly. Based on the current findings, small differences betweenductility loss values are not sufficiant for a proper evaluation regardingthe extend of hydrogen damage in the broad sense of hydrogen degradationpotential of 304L and 316L. Consequently, it is proposed to use hydrogeninduced ductility loss only as a warning sign of possible hydrogen influ-ences which often seems to be masked in metastable austenitic stainlesssteels.

ACKNOWLEDGEMENT

The authors wish to express their appreciation to Mr. M. Kupiec and Mr. M.Aberman, for experimental assistence.

REFERENCES

1. R.P. Jewett, R.J. Walter, W.T. Chandler and R.P. Frohmberg, Rocketdyne,Division of North American Rockwell, Canoga Park Calif. 91304 (NASACR-2163), March 1973, p. 67.

2. V.K. Sikka, R.W. Swindeman and C.R. Brinkman, Proc. ICF4, (Waterloo,Canada, June 19-24, 1977) D.M.R. Taplin (Ed.), University of WaterlooPress, 1977, vol. 2, pp. 561-567.

3. M.R. Louthan, G.R. Caskey, J.A. Donovan and D.E. Rawl, Mater Sci. andEng., 1972, vol.10, pp. 357-368.

4. G.R. Caskey, Proc. Symp. on Environmental Degradation of EngineeringMaterials, (Blacksburg, Va., Sept. 21-23, 1981) M.R. Louthan and others(Eds.), Blacksburg Va. Polytechnic Ins., 1981, pp. 283-301.

5. H. Mathias, Y. Katz and S. Nadiv, Proc. Int. Symp. on Metal-HydrogenSystems, (Miami-Beach, Florida, April 13-15, 1981) T.N. Veziroglu (Ed.),Pergamon Press, Oxford, 1081, pp. 225-249.

166

r

-

\

a

-

-

304 L

0

5i

\

—i

unhydr

hydr

T

ogcnated

ogenated

13-.

1

—1

-

-

O i

• •

-a

i

•

-

i

3041

•

/

*• •II

1

\

1

i n

1

unhydroênated °

hydrogenated .

i

-

. . .

•

y

Fiy. 1 Stranght vs. temperature. Fig. 2 Ductility vs. temperature.

^ " •jnhydroRenated

hydrogena ted e__

Fig. 3 Volume fracture of a1 anduniform strain vs. temperature.

Fig. 4 Fracture mode at peripherialregion of 304L tensile specimen.

167

TENSILE FLOW OF AUSTENITIC STAINLESS STEELAFTER THERMAL AGING IN HYDROGEN ATMOSPHERE

Y. Rosenthal1^*, M. Mark - Mark p,w i t c h 1 ,A. Stern and D. E l i e z e r

1 - Nuclear Research Centre-Negev, P.O.B. 9001, Beer-Sheva, I s r a e l .

* - In p a r t i a l fu l f i lment of the requirements for a M.Sc. degree .

2 - Ben Gurion Univers i ty of the Negev, Dept. of Mater ia ls Engineering.

INTRODUCTION

Blanchard and Troiano [ 1 ] , Whiteman and Troiano [ 2 ] , Holzworth [ 3 ] ,Kolts [4] E l ieze r e t . a l . [ 5 ] , Hanninen and Hakkarainen [ 6 ] , Nar i ta e t .a l . [7] and others (8,9) have shown t h a t considerable hydrogen embr i t t l e -ment (HE) or hydrogen-induced degradat ion of p rope r t i e s (HIDP) may beproduced in any a u s t e n i t i c s t a n i l e s s s t e e l (ASS), however, s t a b l e withrespec t to deformation mate rns i t e , provided two condi t ions were met:(1) supply of envi ronmenta lHat high inpu t f u g a c i t i e s ; (2) bu i ld -up oflarge concentrations of internal H across a significant part of thespecimen's cross-section, rather than in some shallow sub-surface layer.These results, while essentially correct, apparently led to the s t i l lpopular view that high charging fugacities of internal H contents are aprerequisite to considerable H effects in AGS. Thus, relatively l i t t leinterest v.as left, the purpose of this study was to investigate theeffects of low contents of internal H, of the order of a few hundreds ofappm, on the following1 properties of selected austenitic stainless steel(ASS) in their thermally-aged (sensitized) condition: 1) plastic flow,strainhardening, ductility and fracture characteristics; 2) micro-structural stability, i .e . extent of strain-induced martensite formationand possible changes of the Md temperature. Two ASS were chosen: 304Lunstable at RT and 316L stable at RT in the hydrogen free condition.

EXPERIMENTAL

The two commercial-grade ASS were received as 0.2 mm th ick sheets inthe br ight -annealed cond i t ion . H_ gas precharging of t e n s i l e specimenswas conducted a t 600°C in vacuum-tight ASS r e t o r t s , equipped with thermo-couple ports and i n l e t / o u t l e t vacuum v a l v e s . The H-pressure in the repor t swas kept constant a t 0.5 MPa-gage for an exposure dura t ion of 170 h r .Control specimens were heated in an argon atmospheis. All t e n s i l e ,t e s t i n g was conducted in a i r a t RT. a t a s t r a i n r a t e of O.5«'1O" 5" . Theload-extension readings were fed to a PDP-11 minicomputer programmed toreduce such data to s t r e s s - s t r a i n and s t r a in -ha rden ing parameters , andcomplete flow curves t o o . Hollomon's empi r ica l power equat ion wasemployed as a b e s t - f i t function:

a = a • £o p

168

where a- true flow s t r e s s ; a = strength coefficient; £ - true p l a s t i cs t r a in ; n - s t ra in hardening°exponent. With accurate, high-resolutionload-extension readings, the Hollomon equation yields for some ASS log-logflow curves comprising three l inear s tages , characterized by different nand a values.


Tables 1 and 2 summarize the mean values of t ens i le properties ofthe 304L and 316L ASS, respectively, thermally aged at 600°C. Theproperties of hydrogenated specimens are compared to those of controlspecimens (Ar) in terms of a "hydrogen-induced change" parameter,AH% 100(11,,-". )/H, . The additional parameters serve to emphasize theidea that hydrogen-induced or ass is ted fracture is but the final eventof a continuous, probably complex s t r a in h is tory . When t h i s idea isoverlooked and the "conventional" parameters solely are invoked usc r i t e r i a of H effects , an incomplete or even misleading picture mightemerge.

The duc t i l i ty loss of the 316L (-18.8%) i s roughly hal f that of the304L (-34.8%), the more stable s tee l would appear to be considerablysuperior to the metastable one, at l eas t in the given conditions ofhydrogenation and t e s t i n g . On the other hand, judging by the loss inuniform elongation, Ae , the two s tee l s appear to be embrittled to nearlythe same degree: -30.#% for 316L vs. -36.6% for 304L. Thus, the super-io r i t y of the 316L over the 304L in the annealed condition drops consider-ably in the sensi t ized condition. The somewhat larger localized (necking)s t ra in of the 316L, e~ - e = 8% vs. approximately 3% for 304L, seems tobe associated with a different fracture mechanism, as discussed la te r .

Further, hydrogenation apparently had l i t t l e effect on the 0.2%yield s t ress and UTS of ei ther s t e e l , i f these parameters are examinedout of context of the respective flow curves. The very small differencesin 0.2% Y.S. and UTS between hydrogenated and control specimens couldreadily have been dismissed as experimental errors and data sca t t e r . Theimplicit and misleading conclusion would have been that hydrogenation inour experiments had l i t t l e effect on the ' 'overall" strength of the 304Land 316L ASS. The picture i s s igni f icant ly al tered, however, i f weexamine the whole complex s t ra in h is tory by the events of gross yieldingand p las t i c i n s t a b i l i t y (onset of necking). To th is purpose we propose anew cri ter ion of hydrogen effects , already implicit in the definition ofthe "hydrogen-induced property change", AH in Tables 1 and 2: the"hydrogen-sensitivity of flow s t r e s s " , which is simply a plot of thehydrogen-induced change in engineering flow s t r e s s ,

SH " SArAs = —£ 100 vs. the log of engineering p l a s t i c s t r a in e ., %°Ar P1

(Fig. 1).

The plot of As - log e j in Fig. 1 is drawn up to e = 40% andthus does not include AUTS:P the UTS of hydrogenated andpcontrol (Ar)specimens are not proper flow s t resses as (1) they occur a t differents t ra ins , and (b) they mark the t rans i t ion from a uniaxial s ta te of s t r e s sto a t r i v i a l one. The flow s t ress of both 304L and 316L ASS shows

169

significant H effects over a large part of the uniform s t ra in range. Forthe 316L the effect is one of general "hydrogen hardening" from macro-yielding up to e = 40% (and beyond, see Table 2). The 304L exhibitsslight "hydrogenpsoftening" over the entire gross yielding range, up toabout 0.8% strain; the effect then reverts to one of hydrogen hardening,quite parallel to that of 316L. Beyond e , = 40%, however, As dropsstejply towards a negative value of AUTS giving rise again to slightsoftening.

Thus, ther- ...-barging to low contents of internal H did notaffect the overall strength of both s tee l s , and the largest effectsoccurred at flow stresses ( s ^ - s.o5.) which are not among the mechanicalparameters usually reported in the l i t e ra tu re .

Any H v3ffects on the strength of a metal ultimately must originatein some interaction between the dissolved hydrogen and the mechanismswhich govern the plas t ic deformation of the metal. Some qualitativeinformation relevant to such interactions may be obtained from strain-hardening displays, e.g. comparative Hollomon plots (Fig. 1 and 2).Further, Table 3 displays the values of the strain-hardening exponent nin the Hollomon equation and the true p las t ic strain £_ marking theonset of stages 2 and 3 in three-stage Hollomon plots .N

CONCLUSIONS

1. A low content of internal hydrogen, of the order of 300 appm, wasshown to induce significant degradation of tensile properties in thermallyaged (sensitized) 304L and 316L type austenit ic stainless s t ee l s :ducti l i ty losses, hydrogen hardening (flow stress increase), decreasedstrain-hardening capacity, and changes of fracture mode (304)L ormorphology (316L) .

2. Hydrogen hardening and the changes in strain-hardening capacitymainly occurred over an intermediate range of s train; the yield stressand UTS of both steels were l i t t l e affected by hygenation and thus shownto be doubtful or even misleading c r i t e r i a of hydrogen effects whenevaluated out of the context of overall deformation behaviour. Uniformelongation was shown to be an important criterion of hydrogen effects,besides the elongation to fracture.

3. Thermal pxecharging to low and uniform contents of internalhydrogen, with no surface damage, and a "scan" of the ent i re p las t icrange appear to be experimental techniques particularly useful in thestudy of intr insic hydrogen effects in austenitic stainless s tee ls .

ACKNOWLEDGEMENTS

The authors thankfully appreciate the competent assistance of Mssr.R. Frenkel and A. Magen (NRCN).

170

REFERENCES

1. P.A. Blarichard and A.R. Troiano, Mem. S c i . Rev. Met., 57 (1960) 409.

2. M.B. Whiteman and A.R. Troiano, Corrosion, 21 (1965) 53.

3. M.L. Holzworth, Corrosion, 25 (1969) 107.

4. J . Kol ts , in S t ress Corrosion - New Approaches, STP 610, ASTM,

Phi lade lphia , Pa. 1976, p . 366.

5. D. E l i eze r , D.G. Chakrapani, C.J. A l t s t e t t e r and E.N. Pugh, Meta l l .

Trans. lOA (1979) 935.

6. H. Hanninen and T. Hakkarainen, Meta l l . Trans. lOA (1979) 1196.

7. N. Nar i t a , C.J. A l t s t e t t e r and H.K. Birnbaura, Metal l . T r a n s . , 13A

(1982) 1355.

8. E. Kinkovitz and D. E l i e z e r , "Grain Size and Heat Treatment Effects in

Hydrogen Assisted Cracking of Aus ten i t ic S ta in le s s S t e e l s " , J . of

Mater ia ls Science, 17 (1982) 3165.

9. P. Rozenak and D. E l i e z e r , "Effect of Metal lurgical Variables on

Hydrogen Embrittlement in Type 316, 321 and 347 S ta in l e s s S t e e l " ,

Mater ia ls Science and Engineering, 61 (1983) 31.

171

(I.

71 .

II .

II .

tt .

II .

I II I II S II'

Fig. 1: Hollomon Flow curves of 316L ASS thermally aged at 600 C.

i ii' i 10* 5 It

Fig. 2: Hollomon flow curves of 304L ASS thermally aged at 600°C.

i.-ible 1 Effects of Hydrogen on Tensile Properties of Thermal ]y Aged 304 L ASS

Property\

\

Kngineering stress S, MPa

at Engineering strain e ,

Fract. Uniform TotalUTS Stress elong. elong.

MPa MPaExposure 0.05 0.1 0.2 0.5 O.b 1 5.5 10 20 SO 40

2 6 6 . 5 2 1 > 9 . 3 2 7 3 . 2 2 S 4 . S 2 S S . 5 3 1 2 . 0 3 4 4 . 5 5 " 6 . 4 4 2 0 . 0 4 S 0 . 5 5 S 8 . S 6 1 7 . 5 6 9 5 . 2 6 9 6 . 6 6 3 9 . 5 4 S . 2 3 5 1 . 1 7

il,,'h0i'OC

2.8 2.4 2.1 2.S 2.5 4.6 t>.3 8.4 16.2 24.6 1.5 25.4 59.8 55.9 2.24 6.48

2 7 4 . 4 2 7 o . 5 2 8 0 . 4 2 9 0 . 3 2 9 3 . 0 5 0 5 . 4 5 3 1 . 5 5 2 5 . 6 5 9 0 . 3 4 5 6 . S 5 4 7 . 1 6 0 7 . 1 6 4 7 . 5 7 - 1 7 . 5 6 9 6 . 1 7 6 . 0 5 7 S . 5 O

A r / 6 0 0 C

S.S 9.0 9.0 11.5 12.b 12.0 14.8 15.4 17.2 15.0 13.2 9.5 11.7 9.1 14.5 0.80 1.50

Hydrogeninducedpropertychange

-2.9 -2.5 -2.6 -1.9 -1.5 5.9 5.S 7. (> 5.1 7.6 10.0 7.4 -S.I -56.6 -54. S

ii.2 O.2 (1.2 0.2 0.1 0.2 0.4 0.5 0 . " 0.5 0.5 i>. 3 0.4 2.1 5.1

Standard deviation irrelevant, as UTS and Sf are not flow stresses.

Table 2 Effects of Hydrogen on Tensile Properties of Thermally Aged 516 L ASS

Fract. Uniform Total

Propi-rty Engineering stress S, MPa UTS stress elong. elong.\

at Engineering strain e , °\ >' MPa MPa

LVOburc Li. 05 0.1 0 .2 0 .5 0 .6 1 2 5.5 5 10 20 50 40

Ar/(i0i'°C

32 S.!- 329.5 531.0 357.0 345.5 559 .3 392.8 426.2 475 .0 539.0 642.7 twd.O 695 .0 690.7 656.6 43.01 51.24

4.4 4 . 5 6.2 6.7 6.4 7.5 2 .2 2.1 10.8 4 .2 7.6 8 .3 1.7 12.2 7.K 4 .55 0.4 7

525.S 326 .5 326.8 522.3 556.6 345.6 56S.S 400 .5 450.5 492 .2 578.S 626.7 656 .0 6S4.7 657.0 62.21 65.14

5.5 5.9 5 .6 5..S 3 .1 A.I 5 .9 7.2 7.9 5 .3 3.0 5.4 4 . 5 20.0 11.3 3.76 4 .02

inJiiccd

i'!\>l'CTty

i-li.mgoi> 9 1.0 1.3 4.6 5.0 4.6 l>.5 6.5 10.3 9.5 11.0 ".9 5.9 0.9 -0.1 -50.9 -JS.S

an 6 _ _ _ _ _ _ - _ - - - - - - - -

0.(i 0.U 0.(1 il.l H.l 0.1 0.: 0.2 (1.4 0.2 0.2 (1.2 0.1 * * 5.1 2.4

* Standard deviation irrelevant, as UTS and S, are not flow stresses.

174

Tables: Strain-hardening Parameters of Thermally-aged

304L and 316L ASS-

\ Material,Parameter Vondition

n

304L

316L

nH • nAr ,n *00nAr

e

304L

3161

304L

316L

ZAr

H-IAr

H92

Ar

H

Ar

Stage

0

0

0

0

07

.07

.05

.05

-2

-2

1

+ 0

± 0

± 0

+ 0

.9

.2

/

/

t

.01

.02

.01

.02

Stage 2

0.

0.

0.

0.

4.

1.

5.

1.

17+0

24+0

14 ± 0

22 ± 0

-31.

-35.

5-10"3

7-10"2

4-10"3

6-10"2

.01

.02

.01

.05

3

5

Stage 3

0

0

0

0

6

9

5

9

.47 + 0

.50 ± 0

.38 ± 0

.43 ± 0

-6.6

-11.8

.3-10"2

.6-10"2

.2-1O"2

.o.io"2

.01

.04

.01

.01

* computed from three-digit values of n, in order to resolve the

minute yet real H-induced decrease of n over Stage 1.

175

CORROSION BEHAVIOR OF Al-Hu ALLOY THIN FILMS IN MICROELECTRONICS

J . Zahavi and M. Rotel

I s r ae l I n s t i t u t e of Metals, Technion, Haifa, I s r a e l .

H.C.W. Huang and P.A. Totta

General Technology Division, IBM Corporation Z/48A

Hopewell Junct ion, NY 12533, U.S.A.

INTRODUCTION

Uuminum and aluminum-copper film a l loy have been used ex tens ive ly asconductor l ines for t he production of VLSI c i r c u i t s . I t i s known tha t\1-Cu bulk a l loys a re suscep t ib le t o general and loca l i zed corrosionCI,2) as was also observed with Al-Cu t h i n film al loys by severa linves t iga to r s ( 3 , 4 , 5 ) .

Recently (6 ,7 ) , e lectrochemical p o l a r i z a t i o n techniques have been usedin assessing resistance to localized corrosion processes in bulk (6)and thin film materials (7). However, no substantial or systematicresearch work has been reported on the behavior of Al-Cu thin filmalloys and their susceptibility to localized pitting corrosion.

The present work aimed at studying the relationship between Al-Cu filmelectrochemical polarization behavior and film susceptibility tolocalized pitting corrosion, as well as various factors such as heattreatment, copper concentration affecting film corrosion behavior andcorrison resistance.

EXPERIMENTAL

Al-Cu al loy th in fi lm specimens were E-beam evaporated on oxidizeds i l i c o n subs t ra tes with Cu concentrat ion varying from 0% t o 8% wt%.Film thickness was 900 nm. Specimens were annealed for 1 h r a t 400°Cin i n e r t atmosphere of flowing argon gas and cooled in a i r stream.

Po la r i za t ion was conducted in a s t i r r e d ba th of deionized water con-t a i n i n g 10~3M of Cl- ions at 25°C and a t a sweep r a t e of 0.2 mv/sec,while ni t rogen gas was purged continuously in to the s o l u t i o n . Amicroprocessor-based corrosion measurement system (8) model 350A,toge the r with IR compensator module 365, manufactured by EGGG, NJ, USA,was used in conducting the po l a r i z a t i on t e s t s .

Cyclic polarization scans started at init ial potential, which was about

176

250 mV below the given corrosion potential, E c o r r. The scan potentialincreased toward the noble direction while ';he corresponding currentwas recorded. After reaching the corrosion potential, the scan poten-tial continued to increase in the noble direction reaching the Fladepotential, Ef, and thereafter the breakdown potential, E^, as indicatedby a rapid increase in current. Thereafter when the scan potentialreached a predetermined vertex potential and vertex current, a reversescan was started and continued toward more active potentials untilrepassivation occurred, characterized by the protection potential, E_.The hysteresis loop was completed when final predetermined potentialwas reached (Fig. 1).

Log Current Density

Fig. 1: Typical Cyclic Polarization Curve.

Samples before and after polarization were examined by optical micro-scopy, scanning electron microscopy (SEM) and Auger electron spectro-scopy (AES). AES survey and depth profile were taken with a PHIScanning Auger Microprobe, model 590A.


Typical potentiodynamically cyclic polarization curves of as depositedand heat-treated Al-Cu thin film specimens are shown in Figs. 2A and2B, respectively. The presence of hysteresis correlated very wellwith the mode of corrosion observed on film surface after the polari-zation tests. Localized pitting corrosion processes occurred onpolarized specimen that displayed a well-defined hysteresis as shownin Fig. 2.

177

0.4

0.0 -

-0.4

-0.8

LLJ

q -1-2en

£ 0.4o

0.0

-0.4

-0.8

-1.2

FILM ALLOY : AL - 4%Cu ; 9000^TREATMENT: AS DEPOSITEDSOLUTION : D.I. WATER+ 10"3MCI" "

FILM ALLOY : AI - 4%Cu ; 9000A"TREATMENT: HEATING, 400°C v 1 hrSOLUTION : D.I. WATER +10"3MCL~

10° 10' 1OZ 105

nA /cm2104 105 1Oe

Fig. 2: Typical c'yclic potentiodynamic polarization curves for Al-4. %Cuthin film alloys. A. As deposited film. B. Heat-treatedfilm.

A correlation between the corrosion resistance expressed by thedimensionless ra t io (TL^-E^/EyEf) to the various amounts of copperconcentrations in as deposited and heat-treated films is given inFig. 3. Each data point represented at least two measurements.

Increase of copper concentration up to 4% in film specimen resultedin sharp decrease of film apparent corrosion resistance, namely thedimension!ess r a t io , E_-Ef/Eb-Ef, for as deposited and heated filmspecimen (Fig. 3). However, further increase of copper concentrationup to 8% did not significantly affect the rat io values. I t should benoted that Al-Cu film containing less than 1% copper concentration

178

exhibited very high corrosion resistance.

Heat treated film specimen exhibited higher values of the ratioE -Ef/E, -Ef compared to non-treated or as deposited films (Fig. 3),although the nature of the correlation given in Fig. 3 did not changewith specimen heat treatment. Higher values of this ratio indicatedhigher corrosion resistance.

1.0 -

ae -

o.e -

ui 0.2 -

1 1

FILM ALLOY :TREATMENT :

SOLUTION :

_

I i

l

Al-Cu: MOOX

A AS DEPOSITEDo HEATING -tOO'CD.I. WATER + 1O"1

1

, IhrMCl"

i

_

-

• o "

12 4COPPER CONCENTRATION

6[Wt%]

Fig. 3: Dependence of Al-Cu thin film corrison resistance on filmcopper concentration.

Figs. 4 and 5 show typical views of localized corroded flower typezones in heat treated films. It was noted that the "flower", or the de-iendritic type corroded sites resulted from lateral corrosionprocesses. Copper rich second phase particles (point 1, Fig. 5B)were associated directly with localized corrosion processes as shownin Fig. 5B. These particles acted as efficient cathodic sites incomparison to the Al-Cu matrix (point 3 in Fig. SB], and resultedin localized corrosion process attack at the grain boundaries whichled to the removal of complete grains (point 2, Fig. 5B).

Fig. 4: Typical view of corroded zone of polarized heat treatedAl-8%Cu thin film alloy (900 mm in thickness). Opticalmicroscopy (x200).

179

Fig. 5: Typical SEM observation of localized corroded zones ofpolarized heat treated Al-l-2%Cu thin film alloy (900 ironin thickness). A. General view of "dendritic" or "flower"-like localized corroded zones (xlOOO). B. High magnifica-tion of localized corroded zone shown in (A), Point 1 -copper-rich particle; point 2 - missing grain; point 3 ••grain matrix (x5000).

ocalized pits in the as deposited film are shown in Fig. 6. Exposuref the silicon substrate and the absence of aluminum copper film at

180

the localized pits were confirmed through X-ray mapping for silicon(bright spots in Fig. 6B) and aluminum (bright spots in Fig. 6C) .High magnification view of aluminum film-silicon substrate at thebottom of the pit (Fig. 6D) indicated that the corrosion processesresulted in gradual film dissolution up to the silicon substrate.

Fig. 6: Typical SEM observation of corroded zones of polarized asdeposited Al-4%Cu thin film (2000 nm in thickness).A. Typical view of rounded pits reaching the silicon subsrate.B. X-ray image of Si(Ka) showing the presence of silicon at

the bottom of the pits (bright spots in the picture).C. X-ray image of Al(Ka) showing the absence of aluminum at

the bottom of the pits (dark areas are very low concen-tration of bright spots).

D. X-ray image of Al(Ka) of the area shown in (E) . T.rightspots indicated the presence of Al.

E. Typical view of the interface between corroded film and tliesilicon substrate at the bottom of the pits.

Typical AES depth profiles of pitted zones in as deposited Al-Cufilms are given in Fig, 7. These depth profiles showed that the Cu-to-Al signal intensity ratio in pit's walls (Fig, 7) was about ten timeshigher than that for film background while pit's center is mostlysilicon .

181

<LJQ.

1C

UJ d

a.

PIT CENTER

AL

PIT SLOPE

Si

AL

4 6 8 10 12 14SPUTTERING TIME (min)

16 18

Fig. 7: Typical AES depth profile analysis of as deposited Al-4%Cuin the pitted area shown in Fig. 6. Sputtering rate was2.5 nm/min.A. AES depth profile taken x>om the bottom of the pit,

point 1 in Fig. 6.B. AES depth profile taken from the pit's walls, point 2 in

Fig. 6.C. AES depth profile taken from non-corroded polarized film

surface, point 4 in Fig. 6.

Mode of corrosion processes. Aluminum alloy usually exhibited highcorrosion resistance due to formation of passive stable ronconductivealuminum-oxide layers. However, when discontinuous conductive oxidefilms were formed, either due to embedment or fast dissolution ofintermetallic particles (10,11), or to breakdown events within theoxide coating (12), severe local corrosion processes took place (2).Aluminum-copper bulk alloys containing reactive copper-rich inter-metal lies usually form nonuniform oxide film with non-regular struc-ture (10,ll)which resulted in initiation and propagation of localizedcorrosion process.

182

Strehblow et al. (5) indicated the importance Oi oxide coating incorrosion protection and corrosion behavior of Al-Cu thin film alloys.Coating stability and protective properties were dependent on copperconcentration in Al-Cu thin film alloy. Heat treated Al-Cu thinfilms obtained in this study with copper concentration up to 8% con-tained copper-rich intermetallics (up to 1 \im in size - Fig. 5),which significantly affected the uniform and continuous formation ofoxide coating.

Copper concentration significantly affected the formation of copper-rich intermetallics in heat-treated Al-Cu thin film. With increase offilm copper concentration the number of copper-rich particles wasincreased and therefore affected film corrosion behavior and resistanceto localized corrosion. Heat-treated Al-Cu thin film containingcopper concentrations of 0.5 to 1% did not produce copper-richparticles which immediately affected the corrosion resistance of thefilm, as shown in Fig. 3.

Copper-rich particles served as very efficient cathodic sites comparedto anodic copper depletion zones at film grain boundaries or filmgrains (2,3,9). Anodic dissolution of Al (Al •> Al + 3 + 3e~), therefore,took place preferentially from grain boundaries while cathodicreactions consuming electrons such as 2H2O + 02 + 4e~ ->- 40H tookplace on copper-rich particles resulting in localized intergranularcorrosion followed by removal of complete grains (Fig. 5) known asthe missing grain phenomenon (3). The presence and remainder ofcopper-rich particles within the localized corroded areas (Fig. 5B)indicated that these particles actually served as cathodic sitesmaintaining their initial shape and size.

As deposited Al-Cu thin film alloys with copper concentration up to 8%neither showed a distinc: '. re grain structure nor the presence ofcopper-rich intermetallies before and after polarization. However,severe localized corrosion processes characterized by single isolatedrounded pits (Fig. 6) were observed,

AES examination of corroded and non-corroded zones (Fig. 7) indicatedthe presence of localized copper concentration particularly withinpitted zones as well as the presence of about 40 nm oxide coatings innon-corroded zones. Aluminum oxide coatings covered the Al-Cu filmsurface. However, oxide thickness was reduced significantly at Al-Cuthin film zones containing local high concentrations of copper.Furthermore, these were probably copper-rich coatings with conductiveproperties (5,13,14) acting as local sites for cathodic reactions.The aluminum oxide coating formed adjacent to the copper-rich oxidecoating consisted of flaws (15) and/or breakdown events (12). Thepresence of chloride ions in the D.I. water used prohibited the repairor the healing processes of the oxide coating (12) which resulted inlocalized corrosion processes. The initiation and propagation oflocalized pitting corrosion process took place while the adjacent Al-Cuthin film copper-rich zones served as efficient cathodes with surfacearea much larger compared to the copper depleted anodic zones.Anodic dissolution of the aluminum ions resulted in pit propagationreaching the silicon wafer substrates (Fig. 6).

183

Conclusions. The contribution of this study was to evaluate thesusceptibility of Al-Cu thin film alloys to localized corrosion pro-cesses making use of electrochemical potentiodynamic cyclic polariza-tion techniques. A dimensionless ratio between perfect passivation(E -Er) and total passivation range (E^-Ef) was defined to indicatethl susceptibility to localized corrosion. The higher the ratio, thehigher the resistance to localized pitting corrosion. Heat treatedAl-Cu thin film exhibited higher resistance to localized corrosioncompared to as deposited Al-Cu thin films. The susceptibility tolocalized corrosion increased with increase of copper concentrationsboth in heat treated and as deposited Al-Cu thin film alloys. Heattreated Al-Cu thin films exhibited shallow branched dendritic or flower-like localized corroded areas which were directly associated with thepresence of copper-rich intermetallics preferentially at grainboundaries.

In as deposited Al-Cu thin film alloys, well defined isolated roundedpits were observed. These were associated with localized copperconcentrations, but not with copper-rich particles or grain boundaries.

REFERENCES

1. N.D. Tomashow, "Theory of Corrosion and Protection of Metals",Macmillan, New York (1966).

2. J.R. Galvele, S.M. de DeMichelli, I.L. Muller, R.B. de Wexlerand I.L. Alanis, "Critical Potentials for Localized Corrosion ofAluminum Alloys", in "Localized Corrosion", NACE, EditorsR.W. Stachle, B.F. Brown, J. Kruger and A. Agrawal, Volumepublished in 1974.

3. P. Totta, "The Missing Aluminum Problem" IBM Technical ReportTR 22.1447 (1972).

4. W.T. Lee, J.M. Eldridge and G.C. Schwartz, J. Appl. Phys. 5_2_,4, pp. 2994-2999 (1981).

5. II.H. Strehblow and C.J. Doherty, J. Electrochem. Soc. Vol. 125,No. 1, pp. 30-33 (1978).

6. R. Baboian and G.S. Haynes, "Cyclic Polarization Measurements -Experimental Procedure and Evaluation of Test Data", Electro-chemical Corrosion Testing ASTM STP 727, F. Mansfeld andU. Bertocci, Eds., American Society for Testing and Materials,pp. 274-282 (1981).

7. J. Zahavi, M. Rotel, H.C.W. Huang and P.A. Totti?, "CorrosionBehavior of Thin Films in Microelectronics". First AnnualReport, No. 5095-32, September 1983.

8. W.M. Peterson and H. Siegerman, "A Microprocessor-Based CorrosionMeasurement System", in Electrochemical Corrosion Testing, ASTMSTP 727, F. Mansfeld and U. Bertocci, Editors, American Societyfor Testing and Materials, pp. 390-406 (1981).

184

9. K.R. van Horn, "Aluminum", Vol. I, p. 49, 115, 2091, AmericanSociety of Metals (1967).

10. J. Zahavi, A Zangvil and M. Metzger, J. Electrochem. Soc.125, 268 (1974).

11. J. Zahavi, H. Kerbel and 0. Korotkina, J. Electrochem. Soc. 129,7 (1982).

12. J. Zahavi and M. Metzger, J. Electrochem. Soc. 121_, 268 (1974).

13. G.C. Wood and A.J. Brock, Trans. Inst. Met. Finish. 44, 189 (1966),

14. J. Cote, E.E. Howlett and H.J. Lamb, Plating, 57_, 484 (1970).

15. J.A. Richardson and G.C. Wood, Corrosion Science, 10, 313 (1970).

185

QUANTITATIVE NONDESTRUCTIVE EVALUATIONUSING ULTRASONIC WAVES

Laszlo Adler

Department of Welding EngineeringThe Ohio State University

Columbus, Ohio 43210

INTRODUCTION

Recent developments of materials characterization by ultrasonic wavesis summarized by presenting a systematic approach to discontinuity analy-sis. Figure 1 presents a systematic approach to ultrasonic evaluation ofmaterial structures which may be a weld or a joint, for example:

• Locate all flaws• Characterize each flaw (determine size, shape, orientation andcomposition)

• Characterize the material (determine the elastic properties, grainsize, surface roughness, etc)

• Evaluate the seriousness of the flaws' presence (using fracturemechanics techniques).

In order to determine the presence (or absence) of weld discontinui-ties, an ultrasonic image is produced. Defects will be delineated asareas of increased ultrasonic-echo return. For flaws large compared withthe beam dimensions, the ultrasonic image shows the extent of the defect.Flaws on the order of and smaller than the ultrasonic wavelength scatterthe incident sound beam. Some of this energy returns to the transducerand appears on the image as a weakly scattering region in the material.Regions such as these are flagged for later investigation.

Next, the characteristics of the weld metal are measured. The grainstructure and roughness of the ultrasonic beam entrance surface will havean effect on all subsequent measurements, and so must be determined. Inaddition, the mechanical properties of the material surrounding any flawswill profoundly affect the strength of the weld. Also, regions of theweld containing dense porosity or clouds of inclusions will lower weldstrength. Determinations of the concentration and size distribution ofvoids or inclusions should permit calculations of estimated weld mechani-cal properties.

One now returns to the suspect regions in the welded structure,namely areas of increased echo return in the ultrasonic image. A broad-band ultrasonic wave is directed toward the suspect region and the back-scattered signal is then processed to obtain a magnitude spectrum (ampli-tude versus frequency)„ If the spectrum shows deep and periodic modula-tion, the suspect region contains a planar discontinuity (eug. crack). If,however, the spectrum is relatively smooth, the flaw is volumetric (e.g.,pore). Characterization of the defect requires further signal processing.,Discrimination of planar from volumetric discontinuities permits theappropriate processing algorithm to be selected.

Ultrasonic spectroscopy (using cepstral analysis) is utilized for

186

determining the size of crack-like flaws. Upon processing, the distance(along the ultrasonic beam path) is calculated which separates the nearand far edges of the flaw. Interrogation of the defect from a number ofangles may be used to obtain the flaw's size and shape-

Volumetric defects are characterized by "Born inversion" processing..This algorithm returns both a line-of-sight estimate of flaw radius andthe cross-sectional area of the flaw (along the ultrasonic beam path)."Observation" of the defect at a number of "look angles" discloses itsshape and dimensions.

Once all flaws are characterized, data tabulating• Coordinates of all flaws• Shape and size of all defects• Material elastic properties (Young's and shear moduli, Poisson'sratio)

• RMS surface roughness• Concentration and size distribution for areas containing clouds,of porosity, or inclusions

can be output for evaluation using fracture mechanics techniques.

Weld Discontinuity Analysis System

ImagingMapping of Large Flaws and Location of Suspect Regions

Material CharacterizationDetermine Mircostructure, Porosity and Surface Roughness

Investigate Suspect RegionsUse Broadband Backscattering Technique

Planar Discontinuity Volumetric Discontinuity

Ultrasonic Spectroscopy Born Inversion!0>ka>l 2.5>ka>0.5

Determine the Distance Determine the FlawSeparating the Edges of the Flaw Size and Shape

Figure 1. Systematic Procedure for Ultrasonic Nondestructive Evaluation.

187

DESCRIPTION OF SYSTEM ELEMENTS

IMAGINGThe purpose of ultrasonic imaging is to locate regions in the weld

which contain flaws. A focused ultrasonic beam is directed toward theweld. Changes in material elastic properties or density will cause reflec-tion and scattering of the incident beam. A portion of this echo energyis intercepted by the receiving transducer. The receiver output isrecorded with the spatial coordinates of the scattering region (Figure 2).If the entire volume of the weld is scanned one builds up a 3-dimensionalmap of echo amplitude versus position.

Figure 2. Assignment of Coordinates to a Weld.

Display ModesDisplay of echo information is normally presented one plane at a

time. The image is termed a C-scan (echo amplitude at constant depth) ifthe image plans is parallel to the sample surface (x-y plane). It is aB-scan (display brightness indicates echo amplitude) if any other plane(e.g. x-z, y-z or oblique) is presented. It is also possible to displaythe echo strength as a brightness-modulated three-dimensional isometricpresentation (Figure 3). Echo amplitude is shown as display height as afunction of two orthogonal spatial coordinates. Because of the ease inimplementing this isometric display format, it was chosen for all our B-and C-scans.

183

Figure 3. B-Scan Ultrasonic Image of a Fatigue Crack (a). The Rayleighsurface waves used (b) to investigate the crack were generatedby the wedge method.

Image Optimization

The scattering amplitude measured for a particular element in theweld is dependent on

* the ultrasonic wave mode (longitudinal, shear, surface wave),• the frequency of the incident wave,the incident angle and

- the spatial distribution of the incident beam.

189

Each of these factors may be adjusted such that flaw detectibility isoptimized,, Analytical as well as experimental work is needed in thisarea.,

Image Processing

It is possible to accentuate the presence of flaws in an ultrasonicimage by postprocessing. A method, gradient processing, for enhancingedges in an image is introduced here. Let the echo amplitude as a func-tion of spatialposition be represented by f(x,z) [B-scan display mode],,Then the gradient is defined as the vector [1]

G[f(x,z)] = 6f5x (1)

&f_

5Z

The gradient vector points in the direction of the maximum rate of changeof f(x,z) and its magnitude gives the maximum rate of change in f(x,z) perunit distance in the direction of G . For manipulation of the digitalimage, the magnitude of the gradient is approximated by

G[f(x,z)] = |f(x,z) - |f(x,z) - f(x,z+l) (2)

The gradient of the ultrasonic image is displayed if the gradient isabove a threshold value. If the gradient is below threshold (presumablythe case for pixels representing unflawed regions in the weld) the imagepixel is sat to zero. Flaws in the processed image stand out more clearly(Figure 4b) than in the original ultrasonic image (Figure 4a). Researchis needed for adaptive threshold setting, other methods of image pro-cessing and automated flaw recognition schemes.

190

Ultrasonic B-Scan Images of a Specimen containing MultipleDefects. Image before (a) and after (b) gradient processing.

MATERIALS CHARACTERIZATION

The mechanical properties of the material in which a flaw is embeddedmay be as important as defect size in determining strength. Materialcharacteristics which are of importance are listed in Table 1. It ispossible to nondestructively determine many of these properties fromultrasonic velocity and attenuation measurements. Both analytical andexperimental studies have been made to determine how frequency-dependentvelocity and attenuation measurements may be used to infer the concentra-tion and size of pores occurring in dense clouds.

Inhomogeneities (such as pores and inclusions) weaken the structuralcomponents in which they occur. It is important to nondestructivelydetermine the size distribution and concentration of pores or inclusions.A multiple scattering theory was developed for treating wave propagationthrough inhomogeneous material. Matrix and second phase elastic proper-ties inclusion concentration and size distribution are used as input..Dispersion and frequency-dependent attenuation are calculated usingtheory.

191

Table 1

MATERIAL PROPERTIES

Tensile ModulusShear ModulusTensile StrengthShear StrengthBond StrengthHardnessSurface FinishImpace StrengthFracture ToughnessAnisotropyMicrostructureGrain SizePorosity, Void ConcentrationPhase CompositionHardening DepthResidual StressHeat Treatment ProfileFatigue Damage

Ultrasonic Wave Propagation in Cast Iron-Graphite Composite

The same multiple scattering treatment applied to porous media mayalso be used to analyze the problem of an ultrasonic wave traversing amaterial containing clouds of solid inclusions. Cast iron containingcompact flake and nodular graphite, and also specimens of gray iron isstudied. Because the graphite in nodular iron appears mostly as spheres,it was thought the multiple scattering theory could be used with onlyminor changes to include the elastic properties of the graphite nodules.

The properties of the cast iron matrix were estimated from ultrasonicvelocity measurements on a sample of 1045 steel. This type of steel waschosen because it has approximately the same relative percentages offerrite and pearlit-2. Table 2 lists pertinent material properties for thematrix (iron) and inclusions (graphite nodules). Elastic moduli werecalculated from the velocities.

Table 2. Properties of the Constituents in Nodular Cast Iron.

Velocities (km/sec)Material Longitudinal Transverse Density (g/cr3)

iron (1045 steel) 5.8 3.1 7.2

graphite 3.2 1.8 2.2

Values for hte matrix material, iron, and the inclusions, graphitespheres were input to the multiple scattering theory to determine thephase velocity and attenuation as a function of frequency. Narrowband andbroadband ultrasonic experiments were performed on carefully preparedspecimens of cast iron. Figure 5 plots theoretically-determined attenu-ation and experimental measurements. The shape of the curves are

192

identical; however, there is an offset. We ascribe this to incorrectassumptions of the wave speed in the graphite and to the fact we assumeda uniform size of graphite nodules, when there is actually a distributionof sizes.

Surface Roughness

The condition of the surface through which ultrasonic beam enters thematerial can have a profound effect on the frequency content of the pulse,the angle of the refracted energy and on the spatial distribution of thebeam in the solid. Both imaging and defect characterization techniqueswill be affected by the surface finish on the weld under evaluation,, Muchof the unreliability of contact ultrasonic testing arises from multipatheffects and trapped air due to roughness on the sample surface..

We have begun studies aimed at nondestructively inferring parametersdescribing surface roughness (rms roughness and correlation length forrandomly-rough surfaces and periodicity and peak to valley height forperiodic surfaces). Surface roughness may be measured ultrasonically fromthe angular or frequency dependence of wave scattering.

Ultrasonic measurement of rms roughness (h) may be carried out in thefollowing manner. An ultrasonic transducer, operated in the pulse-echomode, is used to record the intensity <I> of backscattered ultrasoniccompressional waves,. The experiment is repeated for a smooth-surfacedsample of the same material (giving I ).

If low frequencies are used, the backscattered intensity is approx-imated as

where k is the wavenumber of the ultrasound in the liquid be.th. Theroughness, h, may be estimated from

-17.37k2/ (4)

dhere AdB is the relative amplitude of the waves backscattered from therough surface compared to that of the smooth surface. The advantages ofthe ultrasonic measurement of surface roughness are

• it is nondestructive,° determinations may be made rapidly• whereas mechanical profilometers measure roughness only along theline of stylus traverse, the ultrasonic method averages over theentire insonifiecl area (be it small or large),,

195

E

eCD•a

12

1.0

0.8

a>'5

S 0.6U

0.4

0.2

0.

« — Theoretical

« - Experimental

.06 .08 .10 .14 .18

Kpa

.22

Figure 5. Attenuation of Ultrasonic Waves in Nodular Cast Iron as aFunction of Frequency (o - multiple scattering theory andx - experiment).

DEFECT CHARACTERIZATION

Ultrasonic images produced in the first stage of our evaluationdisclosed the extent of large flaws and identified weakly scatteringregions as areas potentially containing smaller defects. For a fracturemechanics evaluation, the size and shape of these small discontinuitiesis required (in addition to their location, as determined from the ultra-sonic image). Flaw characteristics may be determined from ultrasonicscattering data; however, the defect must first be classified as volume-tric (3-dimensional, e.g., pore or inclusion) or planar (2-dimensional,e.g. crack-like).

The classification of defect type is carried out by identifyingdiscriminatory features in the backscattering frequency spectrum. Theprocedure below is followed.

1) A broadband ultrasonic pulse is directed toward the suspect

194

region in the weld.2) The backscattered signal from the region is sampled digitized and

stored in computer memory.3) A system normalization signal is acquired by sending the incident

pulse through a unflawed region in the material toward a perfectreflector (polished solid-air surface),,

4) Perturbing effects of the data acquisition system and materialintervening between the ultrasonic transducer and the defect areremoved by deconvolving the defect signal by the signal from thereference reflector.

5) The magnitude spectrum is displayed.• The defect is classified as planar if the spectrum has deep,periodic modulation (explanation given later).

• The defect is classified as a volumetric flaw if the spectrumis relatively smooth.

6) If there is some ambiguity in classifying the defect type (5),then the suspect region is interrogated from a number of angles.• A planar discontinuity will return large backscattering signalsfor orientations in which the incident ultrasonic beam isdirected normal to the plane of the flaw,,

• Amplitude of signals returned from a volumetric discontinuitydepend less strongly on angle.

7) If the discontinuity is planar, an inversion scheme termedultrasonic spectroscopy is followed.

8) If a volumetric flaw is to be characterized, the Born inversionalgorithm is utilized.

The Born inversion procedure is summarized here.

Born Inversion

As an ultrasonic wave strikes a volumetric flaw (Figure 6), thecross-sectional area encountered increases as the wave propagates. Whenthe wave is just incident on the defect (time, t-j) the area encountered isminimal. As the wave moves onward (time t^, then t^, . . „) the are en-countered increases [k^, Ag, . . .)„ The cross-sectional area of the flawreaches a maximum when the wavefront reaches the center of the flaw. Thearea encountered decreases thereafter. The Born inversion algorithmreturns the area function A. versus t^. From this, a line-of-sightestimate of flaw radius may be calculated.

19S

t, M J W

Plane

Wave

Area

A R E A

R a d i u s = 1'2 ( w a v e s p e e d < I,

T i me

Figure 6. Plane Wave Encountering A Scatterer.a) Position of wavefront at times t..b) Area of the scatterer encountered at times ti andc) Area function.

Ultrasonic Spectroscopy

Ultrasonic spectroscopy applied to defect characterization waspioneered by Adler, et« al. [1,2,3].. The fundamental ideas underlying thetechnique are summarized here (details may be found in the book by Fittingand Adler, [4]).

TRANSDUCER

REFLECTOR

Figure 7. Ultrasonic Waves Scattered from the Edges of a Planar Reflector.

An ultrasonic wave directed toward a planar discontinuity will bein part specularly reflected from it, but also scattering will occur fromthe near and far edges (flashpoints, [5]) of the flaw (Figure 7).

196

If the echo is transformed to the frequency domain (via the Fast FourierTransform, FFT) the time spacing of the signals may be determined.

Consider a single signal y(t) which has a spectrum Y(2irf). The mag-nitude spectrum of two such signals, separated in time by 2tQ, has beenshown [6] to be |2 cos2irf tQ| |Y(2Trf) |. That is the spectrum is modulated,

and the period is determined by the time separation (Figure 8). Thespacing (Af) of the frequency minima may be used to determine the timeseparation (At) of the ultrasonic echoes:

Af = l/2t0 = I/At.(5)

If the wave speed is known, then the separation of the near and far flawedges may be calculated and the flaw dimensions determined from

d =(6)2 Afsine

The orientation of the crack (e) may be determined from a number ofangular measurements.

EXAMPLES OF ULTRASONIC NONDESTRUCTIVE WELD EVALUATION

As a test of our systematic method for evaluating welded structures(Figure 1 and previous descriptions) we began with two welds containingwell characterized defects.

Weld sample number 2 was carefully prepared to contain only planardiscontinuities. Three artificial flaws, nominally 1/16", 1/8" and 1/4"in diameter were incorporated into the weld, along its center line(Figure 8 ) . The plane of the defects was parallel to the sample surface..Both the top and bottom surfaces of the specimen were ground flat.

Figure 8. Weld Sample #2 Containing Planar Discontinuities 1/16",i/8" and 1/4" in Diameter.

197

Imaging

A 5 MHz focused transducer (approximately 6" focal distance, 1"diameter) was scanned in a water bath above the weld. The water pathlength was adjusted to focus the ultrasonic beam approximately at the depththe discontinuities were thought to occur. Spacing between transducerpositions (along the weld) was 5 mm.

• the ultrasonic waveform returning from the weld is sampled, digi-tized and displayed (Figure 8 ) ,

• the signal is full-wave rectified (Figure 8 ),• an estimate of the envelope of the waveform is determined(Figure 8 ) ,

• the number of positions along the weld where the transducer willbe located is input (Figure a ),

' wave speed for the mode of ultrasound used in the imaging is input(Figure 9 ),

" the depth increment at which the echo amplitude is to be sampledis input (Figure 9 ),

' a file name identifying the image data to be stored is given(Figure 9 ) and

• the depth (time) range to be used for image display is defined(using the graphic terminal's cursor) (indicated by the verticallines in Figure IO).

Figure 9. B-Scan Ultrasonic Image of a Weld Containing Three PlanarDiscontinuities (1/16", 1/8" and 1/4" in diameter).

198

dB

-40

-60

-80

—

ii'r\ ^7'7 \ AI V \

1

\\jAr

!

//L /\ /

V

\

\'\

i

—

Ni—'"""\ \

\

110 15

Frequency (MHz)20

Figure io- Backscattering Spectra from Suspect Regions in Weld #2..

Ultrasonic Spectroscopy. Broadband ultrasonic pulses from the 15 MHzunfocused transducer were coupled to the weld through a 6" water path.Backscattering signals were acquired for each flaw at 0-, 10- and20-degree refracted angles in the solid. The spectra computed from thesesignals was plotted versus frequency (Figure ). The absence of deepnodulation at O-degrees indicates the plane of each flaw is approximatelyparallel to the surface of the weld. The average spacing of minima inthe spectra was determined (Table 3) and equation (6) used to calculatethe flaw diameter.

Table 3. Average Spacing of Minima in the Backscattering Spectra fromFlaws in Weld #2..

11A

B

C

10-degree Measurement 20-degree Measurement

9

4.7

2.55

MHz

MHz

MHz

4

2.3

1.13

MHz

MHz

MHz

199

Nondestructive Evaluation Summary For Weld #2

Longitudinal Wave Speed:

Shear Wave Speed:

Young's Modulus:

Shear Modulus:

Poisson's Ratio:

Suspect Regions:

Flaw Type:

Flaw Or ientat ion:

Flaw Dimension:

(assume c i r c u l a r l y

symmetric)

5939 m/sec

3267 m/sec

2.566

8.219

0.283

IS.A

B

C

IS.A

B

C

ID

A

B

C

ID

A

B

C

X 101 0

X.

2.0 cm

7.5 cm

12.5 cm

Type

planar

planar

planar

1 10 7.5

0 7.5

0 7.5

Angle with Respect to Surface

0°

0°

0°

10-degree Measurement 20-degree Measurement

1.8 mm 2.06 mm

3.5 mm 3.6 mm

6.36 mm 7.3 mm

Compare these ultrasonically determined results with the actual position anddimensions below

IDA

B

C

X

1.8

7 cm

12.05

cm

cm

Y

0

0

0

7

8

8

Z

. 5 cm

cm

cm

Diameter

1.6

3.2

6.35

(mm)

200

REFERENCES

[1] Ho U Whaley and L. Adler, Flaw Characterization by UltrasonicFrequency Analysis," Mat. Eva!., 2j3 (8), (1971).

[2] H. L. Whaley and L. Adler, Model for the Determination of the Sizeand Orientation of Reflectors from Ultrasonic Frequency Analysis,"J. Acoust. Soc. Am., 48 (1), (1970).

[3] H. L. Whaley and L. Adler, "A New Technique for Ultrasonic FlawDetermination by Spectral Analysis," Technical Memo, Oak RidgeNational Laboratory, ORNL-TM-3056, (1970).

[4] D. W. Fitting and L. Adler, Ultrasonic Spectral Analysis for Non-destructive Evaluation, Plenum Press, New York (1981)„

[5] J. D. Achenbach, L. Adler, D. K. Lewis and H. McMaken, "Diffractionof Ultrasonic Waves by Penny-Shaped Cracks in Metals: Theory andExperiment," J. Acoust. Soc, Am., 66 (6), (1979).

[6] W. A. Simpson, "A Fourier Model for Ultrasonic Frequency Analysis,"Mat. Eva!., 34 0 2 ) , (1976).

201

ON THE HOMOGENIZATION PROBLEM IN SINTERED ALLOYS

L. LevinDepartment of Materials Engineering

Technion, Haifa

A. SternNuclear Research Center, Beer-Sheva

andDepartment of Materials Engineering

Ben-Gurion Univ. of the Negev, Beer-Sheva

ABSTRACT

The present study contains a detailed analysis of the sintering processat temperatures at which one of the components is in a liquid state.Experimental results were obtained for the Fe-Cu system containing 5 wt.%copper. Completion of the sintering process is accompanied by achieve-ment of homogeneity in composition. The different stages of homogeniza-tion were observed by methods of optical microscopy, hot—stage metallo-graphy, scanning electron microscopy plus EDAX and microprobe analysis.At small amounts of copper the process consists of the following stages:dissolution of iron in the liquid copper, covering of iron granule sur-faces by liquid solution; penetration of Cu into the Fe matrix; migrationof the copper atoms towards a uniform distribution. At higher amountsthe process is complicated by penetration of the liquid phase at thegranular interfaces and along grain boundaries. It was shown that theduration of each stage can be calculated when the amount of copper, andmode of penetration are known - by solution of the diffusion equation.The model is applicable to any system with limited solubility of thecomponents.

INTRODUCTION

During liquid phase sintering in the iron copper system a series ofprocesses take place: dissolution of iron in the liquid copper, spreadingof the liquid copper over certain surfaces in the material, penetration ofthe copper into the iron matrix, rearrangement of the iron particles andchanges in the total porosity of the alloy^"^ in the present studyattention was directed mainly at the diffusion processes that bring to ahomogeneous composition of the alloy.

In an earlier work a model for diffusion penetration in the system wasdeveloped. The model was based on the assumption that the copper atomspenetrate the iron matrix from the free surfaces on melting of the copper.The assumption referred to the cast of small amounts of copper in thealloy - 2%, and its validity increases with decrease of Cu content.

In the present study an alloy with 5% Cu was chosen. The behaviour ofthe material during the sintering process was investigated by opticalmetallography, including vacuum hot stage microscopy, and by the SEM andEDAX techniques.

202

EXPERIMENTAL

Table 1 contains data on the powders used in this investigation.

Table 1. Powder Characteristics

Material

ElectrolyticIron

ElectrolyticIron

ElectrolyticCopper

Size

125-150Mm

45-63 ym

45-63 ]im

Amount

66.5%

28.5%

5%

ImpuritiesTotal

240 ppm

240 ppm

415 ppm

Source

BDH LaboratoryChemicals

BDH LaboratoryChemicals

Baker Laboratory

The average grain size of the iron particles was about 20 ym.

After blending for one hour the powder mixture was placed in a neoprenesleeve and isostatically pressed at 10° Pa. The specimens were heatedin vacuum for 10 min at 750°C to achieve mechanical strength necessaryfor further handling.

The sintering process was carried out at 1100°C in two variants:

(1) For short times at temperature, the specimen was heated in a Reichartvacuum hot stage with the sample temperature determined by a thermocoupleembedded directly in the sample. The hot stage microscope was used inan attempt to observe the sintering process in situ.

(2) For long periods up to 24 hours at temperature, the specimens weresintered in a vacuum resistance furnace.

In order to perform most of the process at the chosen temperature, thelast stage of heating was rapid: from 750°G to 1100°C in 4 min. In thefirst variant, cooling was effected by a stream of helium, in the second.-by water quenching. Specimens sintered in the vacuum resistance furnacewere examined by the above-mentioned techniques. The change in thecopper distribution after sintering times of 0, 3, 10, 30 min and after2, 6, 12, 24 hours was checked.

RESULTS

The different stages of the copper melting as obtained with the aid ofthe hot-stage microscope are presented in Fig.l : la - a copper granulejust before -melting; lb - onset of the melting process; lc - a pore atthe spot where the granule had been located; Id - paths of liquid copper

203

penetration in the iron matrix. All the stages shown in the figure areaccomplished in the interval between 1083°C and 1100°C.

Tha ••..'c.roscopic structure obtained in the vacuum resistance furnace after10 mi-i of sintering is seen ic 7ig.2. The dark spots represent pores.A Scanning Electron Micrograph after the same time is presented in Fig.3.On the topography of the structure, the profiles of Fe and Cu obtained byEDAX (Scanning along the central line) are superimposed. The same profilesare shown in Fig.5 at lower magnificatian. The valleys in both profilesindicate pores and do not belong to the "real concentration profile"developed as a result of diffusional processes. The spacing of the "realmaxima" of the Cu profile is not uniform but in most cases follows the Fegrain size (V^se of copper penetration along the grain boundaries). Thedistribution of copper is much more uniform after 6h sintering see Fig.6.An almost uniform distribution of copper is achieved after 12h of sinter-iuij at 1100°C - Fig.4.

ESTIMATION OF TIMES NECESSARY FOR. DIFFERENT STAGES OF H0M0GENIZATI0N

For diffusion calculations we begin with, the time count when 1100"C isachieved and the copper is already melted.

The copper melt can dissolve 4 at.% of iron. The dissolution process ispractically instantaneous as shown in an earlier work^. The FeCCu) meltspreads along the boundaries of the Fe particles of different types(granule surfaces, grain boundaries). The time necessary for the melt toarrive at "starting positions" for the diffusion into the Fe matrix dependson the type of position and is of the order of several minutes^.

An estimate of the thickness of the liquid layer at the starting positionis essential for further calculation; Assuming that most of the melt wasspread along the grain boundaries, the thickness was found as h = 0.5 ym.During penetration of the liquid copper into the iron matrix, the copper,concentration in iron can be estimated^ as

C = CQerfc (x/2/5t) (1)

The total volume of copper penetrating per unit surface of an iron particlecan be expressed as

dx = h<^ (2)

where c_ = 0.96 is the concentration of copper in the liquid, t. theduration of this stage of the reaction, characterized by Co = 0.08 Cmax.solubility of Copper in iron at the sintering temperature) , D is the ,diffusion coefficient of Cu in Fe. According to Smithells and Brandes

D = 1.8 x 10"10 cn^sec"1 at 1100"C.

Taking into account that

/ erfc zdz = 0.5642 , 5

204

the time of penetration can be calculated from eq. (2) :

t = 1.07 x 103 sec - 18 min.

The next stage of the process is migration of the copper atoms towards auniform distribution. The relationship between the effective distancethe solute atoms have to transverse to achieve homogeneity, and theduration of this process can be presented as

KDt, C3)

where K depends on the assumptions concerning the shape of the particlein which the diffusion takes place and on the degree of homogeneity^.

Assuming the simple case

I = 2Dt, C4)

we find the correlation presented in Table 2.

Table 2: Effective distance vs. homogenization time

Effective distancefor Homogenization

5.3 ym

7 ym

13 ym

32 ym

45 ym

55 ym

63 ym

Homogenizationtime

10 min

18 min

1 h

6 h

12 h

18 h

24 h

Taking into account that the center of mass of the copper in the iron authe moment the liquid disappears is removed from the surface by about3' ym , we have for the case when penetration proceeds from a granulesurface I - 55 liia; if penetration proceeds from agrain boundary, £ = 7 ymIn the last case the diffusion time is "VL8 rain; in the first case ^18 h.This is the time necessary for full homogenization. The result is inaccordance with experimental data.

205

ACKNOWLEDGEMENTS

The authors would like to thank F. Simca and D. Shmariahu for technicalassistance and helpful discussions.

REFERENCES

1. L. Levin, A. Stern, S. Dirnfeld, Z. Matallkde 21 (1980), 621.

2. W.A. Kaysser, W.J. Huppmann, G. Petzow, Powder Metallurgy, Q (1980),86.

3. W.J. Huppmann et al., Zeitschrift fur Metallkde, 2P_ (1979), pp. 707,792.

4. C.J. Smithells and E.A. Brandes, Metall. Ref. Book, 5th ed.Butterworths, London (1978), 874.

5. H.S. Carslaw and J.C. Jaeger, Conduction of Heat in Solids, ClarendonPress, Oxford (1967), 59.

6. L.E. Larson and B. Karlson, Materials Science and Engineering, 2£(1975), 155.

206

Fig.l: Dissolution of copper as seen under hot-stage microscope:

(a) copper particle just before melting in iron environment,

(b) onset of melting, (c) pore created as result of Cu

melting, (d) first stages of penetration into Fe matrix.

207

Fig.2: Microstructure of the

alloy after 10 min of

sintering at 1100°C.

Fig.3j SEM + EDAX picture after

10 min of sintering at

1100°C; frott top to

bottom: Fe profile,

scanning line, Cu profilc-

Cu background.

Fig.4: SEM + EDAX picture after 12 hr

sintering at 1100"C; structure,

and concentration profiles from

top to bottom: Fe, Cu, Cu back-

ground.

rwliiJ^r^WK

i£5iiSZs~Z3jM

Fig.5: EDAX: Fe, Cu and Cu

background profiles

10 min of sintering

1100°C.

after

at

208

Fig.6: (a) SEM: structure after 6 hr of sintering at 1100°C,

(b) EDAX: Fe and Cu profiles for scanning line in (a),

209

HEAT TRANSFER TO WATER AND ITS IMPORTANCE FORMETAL CASTING AND HEAT TREATMENT

* **M. Bamberger and B. Prinz

* Technion - Israel Institute of Technology

** Metallgesellschaft AG. West Germany

INTRODUCTIONCast metal properties depend, i.a., on the rate of cooling during

solidification or on heat treatment. A too high rate of cooling downwill lead to thermal s^ressts in the product and consequently to cracksor fracture; on the ether hand, in case of too low a cooling rate, thespecial casting process may lose its metallurgical or economical advan-tages or be a complete failure. Consequently, major importance shouldbe attached to the recognition of coefficients of heat transfer inrelation to different cooling methods, in order to calculate the rate ofcooling and their adaptability to the production process. The prefera-ble method of control over the cooling process is continuous casting.Therefore, most researches deal with cooling methods of this casting,i.e., either by immersion in water or by water spraying. In thisarticle we are going to present a method of measuring heat transfer tothe cooling agent, and the use of continuous casting and heat treatment.

THE PRINCIPLE OF COEFFICIENT OF HEAT TRANSFER

It is impossible to directly measure the heat flow from a hot metalbody to a cooling agent. Therefore the calculation is carried out bymeasuring the body temperature and computation of the temperature field.From these two the heat transfer from the product to the cooling agentis deduced. The experimental set-up is designed in such a way as toestablish conditions of heat transfer in one direction, i.e. theequation, which defines the conduction of heat in a body, in which thetemperature is being measured, is:

where t, X, a, and T are time, coordinate, thermal diffusivity andtemperature, respectively.

In lieu of a continuous description of the temperature field we mayassume the final number N as temperature at discrete points. The heatbalance at every point in accordance with differential equation equals:

where T1 - the previous temperature at point i, and T. , T., T.the present temperature at points i+1, i, i-1. 1

210

At the boundary there exists the relation:

T -TN *N1

X 0K AX ~ Q

With the usual solution 0 is known at the boundary, therefore wehave N equations with N unknown temperatures, which will have to befound by solving a set of equations.

With the cooling experiments the temperature is measured at onepoint T , i.e. with the set of equations we do not know N-l tempera-tures ana the heat flow at the boundary.

The solution of the set of equations will give us the temperaturefield of the product and the heat flow. Control of solution quality iscarried out by means of comparing the calculated temperature field withthe control temperature measured in the body.

With this method the coefficient of heat transfer is calculated bymeans of convections and radiation. In fig. 1 we recognize that this isin accordance with the heat transfer coefficient, which is calculated bythe Sheplan-Bolzman Law. Hence we conclude the method of computation toby correct (1, 2).

HEAT TRANSFER COEFFICIENTS FOR WATER COOLINGHeat transfer coefficient for bars of nickel, aluminium and copper,

when immersed in water, as a function of the surface temperature of thebar, we can see in fig. 2. A definite dependence on the properties ofthe bar, which is in a state of cooling off, can clearly be seen: verylittle with nickel, a little more with aluminium, and the most intensivecooling with copper. Nevertheless, the dependence on surface tempera-tures is identical in these three cases.

In the a/m cases the immersion in water is carried out at 20°C;however, it is well known from literature and heat treatment experim-ents that the temperature of the cooling water influences the rate ofcooling. Therefore experiments with water immersion at various temper-atures have been carried out. In fig. 3 we see that the heat transfercoefficient of a copper bar immersed in water decreased, when the cool-ing water temperature goes up. An increase from 20°C to 60°C broughtabout a drop of 30% in cooling capacity. Furthermore, it can be seenthat the heat transfer coefficient is identical in all three cases.

Spraying water on a hot metal bar surface is accepted cooling pro-cedure, the advantage over immersion in water being the possibility tocontrol the cooling intensity by regulating the intensity of spray. Infig. 4 we show a heat transfer coefficient for cooling by spraying wateron a copper bar. We see clearly an increase of the heat transfer coeff-icient together with an increase of the spraying intensity. Hence, thedependence on the surface temperature is identical with the one of thecooling process by immersion in water.

211

The heat transfer coefficient depends on the following factors:

a) The property of the material to be cooled.

b) Surface temperature.

c) Temperature of the cooling water;

d) Intensity of spray.

Based on a l l previous experiments we arrive at the followingformula:

"spray = 0 " 6 9 lo* <*5> ' f 1 " 4 ^ exP ^ T ^ T > + ^ + «rad

where 'a spray - heat transfer coefficient;V - density of spray (S./m2 min);

X - heat conductivity of metal;p - metal density;c - specific heat of metal;6 - average temperature of coolant;

6 - surface temperature of metal;

0 - evaporation temperature of cooling liquid;

a - heat transfer coefficient during vapour-co

a , - heat transfer coefficient during radiation,rad

RATE OF COOLING AND HEAT TREATMENT

It is an established fact that we can slow down the rate of coolingby adding organic matter to the cooling bath. The heat transfercoefficient drops to 15% of its value during water cooling due to anaddition of 3% of organic matter and arrives at 8% of its immersion-in-water-value, when 10% of organic matter is added.

This change is utilized to control the rate of cooling, when harde-ning - as shown in fig. 5 - a steel plate 30 mm thick.

Quenching in water brings about a structure, which is almost whollymartensite . Adding organic liquid slows down the rate of cooling, andwe, therefore, arrive at a different structure. The more organicliquid reached, the slower the cooling rate, and we have more ferrite asshown by the two curves in the diagram. For comparison there existsanother curve relating to air hardening, where ferrite above 50% isachieved.

212

CONTINUOUS CASTING - RESULTS

Continuous casting of aluminium includes cooling in a water-cooledmetal mould and immersion in water immediately afterwards. From measu-ring the change in temperature of the cooling water in the mould, wederive the complete heat flow from the casting into the mould. It iswell known that the heat flow is relative to the square root of dwellingtime in the mould (3, 4) and thus it is possible to calculate the heatflow from the casting into the mould. During the experiment temperatu-res were measured from a section of an aluminium bar, having a diameterof 190 mm and a solidification front was arrived at. The actual measu-red solidification front and the calculated one are in complete correl-ation (see fig. 6) .

On the other hand there exists a discrepancy between the temperatu-re measured at a distance of 125 mm and 80 mm from the centre and at thecentre itself (fig. 7). The discrepancy is mostly up to 10%, except atthe outflow area of the mould. This error could possibly be the resultof a lack of exact knowledge of the thermo-physical properties of thematerial at high temperatures.

SUMMARY

In this article we have shown the principle of measuring the heatflow from a metal bar to a cooling agent. This method could also beused for measuring the heat flow to various other cooling agents. Aformula has been established for calculating the heat transfer coeffici-ent as a function of cooled-off metal properties, cooling water temper-ature and surface temperature. During experiments a correlation wasestablished between the measured temperature of a cast metal bar and theone calculated by utilizing an established formula in research. Thelatter fact proves the accuracy of the formulating the coefficient ofheat transfer. It provides an optional planning method for the establi-shment of a cooling arrangement for continuous casting and heat treatm-ent.

ACKNOWLEDGEMENTS

The work described was carried out in the Metal1-Laboratories ofMetallgesellschaft AG and with its financial assistance. We would liketo thank the company for allowing publication of this work, the directorof the laboratories, Prof. Dr. Ing. P. Wincierz, for promoting the inv-estigation, and Mr. J. Fachinger and E. Braun for taking part in theexecution of the experiments, as well as Mr. Dohl for his assistance inthe evaluation of experimental results.

"REFERENCES

1. M. Bamberger, R. Jeschar and B. Prinz; Zeitschrift ftlr MettallkundeVol 70 (1979) 9, p. 553.

2. M. Bamberger, B. Prinz - The 47th Foundry Cong. Oct. 1980,Jerusalem

3. R. Alberny et al - Revue de Metallurgie - Juillet-Aout 1976,p.545.

4. G. Vogt and K. WUnnenberg - Klepzigfachbericht 80 (1972)10,p.491.

213

« Cu* Nio A lo vtaiues of N Larrtoert and

M Economopoulos forOxidized Steel

KM 300 <>O0 S00 600 « • 800 900 1000

Surface Temperature <?o m c —-

1 Heat Transfer Coefficient in Air Cooling as

a Funciion of Temperature

100 200 300 400 500 600 700 80C

Surface temperature -&0 in DC

3 Immersion heat transfer coefficient as afunction of cooling water temperature

i00 600 200 (00 500 600Surface temperature in °C »-

2 Heat transfer coefficients in immersion cooling-water temperature 20°C

a) Results of M Bamberger, R Jeschar.B Prin7b) Results of N Lambert. M Economopoulos

105-Material- copper

Cooling inwater

600 ' 1 1 > ' 1 ' 1 r-100 200 300 400 500 600 70C 800

Surfcce temperature TJO m*c — -

^ Heat Transfer Coefficient in Spray Cooling ofCopper as a Function of Temperature

Il

l

-

-

-

-

•••£: \ Ic?""" •

" ' ' -

e

: J"

-/ « * a ic i i

Distance From Center of Cost'ng ((.mi

6 CoTiJuted [JxJ Measured Sol<du5 Lpnem Al'jmrwjm Continuous CaslTig

5. Cooling Curves of 30mm Thick Steel PlateCooled by Immersion into Wpter , Water •Organic Fluid nnd Air Cooling

loa ino 500 TOOTemperature in "c

7. Computed and Measured Temperatures <tAluminium ConlinuOus Costing

214

FAST, NON-DESTRUCTIVE ELECTROCHEMICAL DETECTION OF SURFACE INCLUSIONS INMETALLIC SUBSTRATES

Israel RubinsteinGeneral Electric Research and Development Center

Schenectady, NY 12301

Identification of surface defects in engineering materials is ofconsiderable importance, since such defects constitute a major factorin the determination of the fatigue life of a component. In this con-tribution a fast and non-destructive electrochemical method is presentedfor visual detection of surface inclusions in metallic substrates. Themethod is applicable to non-conductive inclusions, e.g., oxides, carbides,etc., in metals or alloys. It is based on electrochemical coating of thesurface with thin (~1 ym) highly-colored layers. The non-conductiveinclusions remain uncoated and bright, thus becoming highly visible onthe dark background. Very satisfactory results are obtained by employinganodic polymerization of pyrrole in an electrochemical bath.^ Thie

results in the deposition of a black, continuous pclypyrrole film on thesubstrate, in which the flaws appear as bright spots. The method wassuccessfully applied'in the detection of -30-300 ym oxide inclusions innickel-base superalloys. The resultsof typical test experiments arepresented in Figure 1. The two types*of inclusions were (a) sputteredAl2O3 surface flaws, and (b) bulk AI2O3 inclusions, pre-mixed with thealloy by powder-metallurgy techniques. It is evident that the inclusionsbecome clearly visible to the naked eye upon coating with the polymer.3

30 40 50 60 30

Figure 1. Left: Inconel-718 piece with six sputtered AI2O3 inclusions,coated with polypyrrole f i lm. Inclusion diameter, in ym:280 (lower, r ight), 200, 120, 90, 60, 30 (upper, l e f t ) .Right: Rene-95 piece with bulk AI2O3 inclusions, coatedwith polypyrrole f i lm. Average inclusion diameter, 270 ym.

215

References and Notes1. Present address: Department of Plastics Research, Weizmann Institute of

Science, Rehovot 76100 (Israel).2. A.F. Diaz., J.I. Castillo, J.A. Loaan and W.-Y. Lee, J. Electroanal.

Chem. 129, 115 (1981).3. I. Rubinstein, J. Appl. Electrochem. 21> 689 (1983).

216

TELEPHONE TOKENS PRODUCED BY POWDER METALLURGY

Andy Sharon

M.S.H. Sinter Enterprises, Kibutz MefalsimH.P. Hof Ashkelon 79160, Israel

ABSTRACT

The production of telephone tokens by PA' technique is surveyed.The

described process was taylored as to render a good and lasting product

which finds its use in every day life.Manufacturing procedure is described

with emphasis on the P/M advantages over the common procedure of punching

and coining out of a rolled metal sheet.

INTRODUCTION

The manufacture of coins,tokens and medallions by the P/M process is quiteknown(l), although its usage is still not widespread.The P/M has some outstanding advantages over the conventional processmainly when producing tokens of intricated geometrical shape.Telephonetokens are still used in some countries throughout the world.Althoughthis system is inconvenient (tokens shortage), it is efficient in countrieswhere high inflation rates cause frequent changes in prices.This paper emphasises the advantages of the P/M process over sheet punchingand coining.

THE TOKEN

The Israeli token has an intricate geometrical shape with close tolerances.The slot on its diameter(Fig.l) and the 20' slope,makes it difficult tomanufacture it by single punching and coining from a rolled sheet.By usingthe P/M process,the token was produced in almost its final shape duringpowder compaction applying relatively low compacting pressure.

MATERIAL

The material chosen was the cupronickel coinage alloy(75$ Cu - 255? Ni).The powder had the following chemical composition and physical properties;74.655 - 75.155 Copper 0.025% Mn2A.1% - 25.1% Nickel 0.06$ Sulphur0.0155 - Carbon 0.16555 IronFlow rate determined by Hall flowmeter - 28.0 sec/50g.Grain size; +150 microns - 1.2$ wt.

+45 microns - 4-7/5 wt.-45 microns - 51.555 wt.

The powder was fully prealloyed having an irregular particle shape(Fig.2).Prior to compaction,0.7$ Lithium Stearate as lubricant was admixed to thepowder.

217

MANUFACTURING PROCEDURE

Compaction2

Green compacts were pressed at pressures varying between 5 and 7 tons/cmin a rigid steel die. Special adjustments had to be made in order to ensurethe uniform density throughout the various sections of the compacted shape.SinteringOut of the various possible sintering cycles for the Copper-Nickel powdercompacts,the highest possible temperature of 1120"C was finaly chosen.Theeffective dwelling time at the sintering temperature was established at30 minutes.The protective atmosphere was dissociated ajnmonia.

CoiningTumbling of the sintered parts render a smooth and shiny surface,a greatcontribution to the quality and appearence of the final product.The closureof the surface pores following this process,enhances the corrosion resistanceof the P/M part up to the level of the one produced by casting and rollingof metal sheets. „The coining process carried out at 10 tons/eir. renders a final product witha density of 98% of the theoretical density.The close tolerances specifiedby the customer were met.

MECHANICAL PROPERTIESThe mechanical properties of the sintered Cupronickel coinage alloy weretested on standard M.P.I.F. tensile specimens.Comparing the U.T.S. and the Y.P. data for the sintered specimens with thecoresponding data for the wrought material(3),it can be seen that theseproperties are higher in the sintered material.The elongation and impactproperties are louver as can be expected for a sintered material (Table I).

Material Property

U.T.S. (N/mm2)

Y.P. (N/mm2)

Elongation {%)

Hardness (HB)

Table I . - Mechanical P

Rolled Metal Sheet

290

100

35

85

roperties of the Sintered

Sintered Material

300

110

19.5

60

Material vs. the

Wrought One.

MICROSTRUCTURE

The microstructure of the sintered and coined material was compared to thewrought and coined one(Fig.3)The basic phase and twins inside the grains can be seen in both materials.The main difference is the mean grain size.The mean grain size of the sinteredmaterial is 45-50 microns compared to 200 microns in the rolled metal sheet.

218

This difference is of cardinal importance when the token has to activatean electronic device which identifies the composition of the material byfrequency response such as in some vending machines and public telephones.This change in frequency is affected mainly by the number of grainboundaries - grain size present in the same geometrical shape, providedthe porosity of the sintered part is low (less than 3%). This problemcan be otherwise overcome by increasing the nickel content in thealloy, thus altering its paramagnetic properties.

CONCLUSIONS

This paper emphasizes the possibilities of the P/M technique to producetokens and coins in large quantities at high densityes having suitablemechanical properties. By using this process the production cost per unitis considerably lower, having the following outstanding advantages:

1. No scrap left comparing to the punching process,2. Tools life considerably longer - over 5,000,000 units for compacting

tools and over 4,000,000 units for the coining tools.3. Lower pressures required for compacting and coining enables the

producer to make use of cheaper machinery.

The versatility of the P/M process and the above mentioned advantages canpush forward the usage of this technique for manufacture of coins ormedallions.

REFERENCES

1. "Coins, Tokens and Medallions Made by Powder Metallurgy", W.V. Knoopand J.D. Shaw, Int. J. of Powder Metallurgy 3 (1) 1967.

2. Unpublished work - A. Sharon3. DIN Standard No. 176704. "Laminated Cupronickel/Copper Coin Blanks From Metal Powders"

B.G. Harrison and T.R. Bergstrom, Int. J. of Powder Metallurgy3 (4) 1967.

219

SCALE 5:

Figure 1: Production Design for the Telephone Token

Figure 2:

S.E.M. Micrographs

of the Pov.'der Grains.

220

tFigure 3: Mictostructure of

a] Sintered Material b) Wrought material

221

AUTOMATIC DETECTION OF FLUORESCENT INDICATIONS

K.M. Jacobsen

Brent Chemicals International PLC, Ridqeway, Iver. Bucks., England

Fluorescent Penetrant Inspection (FPI) is a well established and provenNon Destructive Testing (NDT) technique which is used extensively inmany industries. Its reliability as a means of finding surface discon-tinuities is affected by the correct choice of penetrant processmaterials, the equipment, and the visual acuity of the inspector.

A choice of penetrant materials may be made from a wide range of excellentcommercially available products. Process parameters appropriate to theintended application are established by experience and experiment, andFPI system sensitivity is verified using test pieces with known defectsas a point of reference.

Fully automated process lines which guarantee that pre-determined FPIprocess parameters are strictly and consistently observed, have been inoperation for more than twenty years. This ensures that parts emergeat the end of the process in optimal condition for visual inspectionunder black light. The equipment has demonstrated its ability toaccommodate parts of varying geometry and surface condition. The formeris achieved by careful jig design which ensures that parts are positionedfor optimal coverage, washing and drainage. Different surface condi-tions are dealt with by providing alternate process sequences, offeringoptions to use penetrants of various sensitivities, variations in hydro-philic emulsifier concentrations and contact times, as well as a rangeof different wash times.

The point of greatest weakness and major source of unreliability of r"PIis the actual visual inspection under black light. Inspectors sit inthe dark, and handle each part individually, carefully examining all itsfaces for fluorescent indications. An experienced inspector with 20/20vision can perceive minute indications on the basis of which he makes anaccept/reject decision, making use of the ability of the humar eye todistinguish high spots of intense fluorescence from the duller backgroundfluorescence due to residual penetrant. Such a high level of perfor-mance can only be maintained for short periods, and is not guaranteedfrom day to day. It is not possible to monitor short-term fluctuations,which go undetected and can lead to indications being missed.

Real improvement of FPI reliability is possible only if the human eyecan be replaced by an instrument able to detect indications in thepresence of background fluorescence.

222

TMThe Automatic Electronic Optical Scanner - AEOS - developed byArdrox, a Division of Brent Chemicals International, has achieved thatobjective. A pre-production prototype machine has been built capable ofinspecting automatically jet-engine blades and similar parts up to 25O mmlong at a rate of about three a minute.

The basic principle is that of a flying spot scan produced by a collima-ted laser beam reflected off oscillating mirrors. Fluorescent lightgenerated by the beam passing over entrapped penetrant is collected byphotomultipliers. The signals are amplified and digitised and processedby a microprocessor. They are used to activate mechanical handlingequipment and displayed on a TV monitor.

AEOS is a go:no-go system which rejects all parts with indications,however produced. These require subsequent visual inspection underblack light. Typically, accepted parts requiring no further inspectionrepresent j-7O-8O% of the through-put; with a corresponding reduction ofthe inspection workload.

The AEOS system may be considered as consisting of three parts: theinspection module, the mechanical handling module and the microprocessor.The inspection module consists of the lightproof inspection cabinet andoptical module containing the laser.

The light source is a He-Cd laser producing a high intensity collimatedbeam approximately 1.5 mm in diameter, in a single wavelength of 442 nm,without any interfering spectral contribution.

The laser beam is reflected off oscillating mirrors. The first mirroroscillates about a horizontal pivot at a selectable frequency betweenlOO Hz and 400 Hz, so as to generate a flying spot in the vertical mode.The second mirror oscillates at a much lower frequency about a verticalaxis and thus induces a horizontal drift and fly-back into the scan.

From the mirrors, the beam is reflected into the inspection chamber,where it arrives as a spot of 1.5 mm instantaneous diameter which scansan area at the required testpiece location of approximately 250 mmheight and lOO mm width.

The spot diameter and mirror oscillation frequencies are chosen so as toensure the illumination of all positions within the active area, whichcan be readily increased or decreased using suitable stops and baffles.To ensure that no stray visible light from the laser excitation mechanismgets into the system, a filter is mounted in the light path.

The testpiece is held in the centre of the inspection cabinet with itsaxis coincident with the central vertical scan line. When the beammomentarily passes over entrapped penetrant it causes it to fluoresce,emitting visible light, which is collected by an array of 6 photo-multipliers placed so that they can pick up light from any point within

223

the scanned area. They are positioned so as to avoid the existence ofshadow regions as might occur with a single photomultiplier. Opticalfilters remove any unwanted reflected or scattered light and ensure thatonly fluorescent light is recorded.

The Mechanical Handling System is designed to present all faces of thetest object to the scan. A small robot removes the test object fromthe fixture on which it is held for penetrant processing, and locates itin a vertical holder. There it is rotated through 360 in steps,pausing at each one to allow the beam co scan each face. The number ofsteps is determined by the geometry of the test objects, three beinggenerally sufficient for round objects, while four steps are necessaryfor objects having a more or less rectangular cross section. Ifrequired the rotation can be performed in a larger or smaller number ofsteps.

A horizontal gripper then ad^^nces and grips the test object midwayalong its vertical axis. The holder releases the test object and thegripper rotates it through 9O in a vertical plane presenting one end tothe scanner, and through 180 to present the other end face. If theshape of the test object requires it, the angle can be varied to overcomeany masking effect due to test object contours. The gripper thenreleases cind deposits the test object on a conveyor belt which carriesit out of the inspection cabinet.

Fluorescent light collected by the photomultipliers is converted into anelectrical signal which is first fed into its associated amplifier andthen into a common adding amplifier, where the individual signals arecombined for any instant in time. The signal then passes through ananalogue to digital converter to the microprocessor, which averages thesignal response over repeated scans, collecting a complete signal intime and dividing it into the elements of an M x N matrix,where M and Nare numbers of the order of 250. This effectively divides eachaveraged vertical scan into M sections, and the horizontal scan into Nsections.

The sensitivity of the system is determined by the Threshold Value whichcan be varied in 64 steps. If an element records a high signal levelwith reference to a predetermined standard - the Threshold Value - ascommanded by the program, then an "indication" is registered. At itslowest setting it will cause the system to trigger when an extremelysmall quantity of light, such as may be due to a small amount of bluefluorescence of small particles of dust, is detected.

The co-ordinates of the <.;.fending matrix element are recorded, and thesignal used to activate the mechanical handling equipment which manipu-lates the test object. For practical purposes the Threshold is set ata value which will ensure that the system is triggered by the smallestindication which the inspection is required to detect. This generallylies within the quartile above the lowe=t of the range.

224

The surface being scanned may be divided into a maximum of four separateareas to each one of which different Threshold Values can be applied.This facility permits the application of different acceptance standardswithin a face thus avoiding miscalls due to a higher than necessaryThreshold setting for less critical areas, while maintaining thehighest level of sensitivity where that is called for.

When an indication above the Threshold is detected, the microprocessorcompares the response of each matrix element with those of its neighbours,and triggers a rejection only when the averaged detected signal isgreater than that of the surrounding area, thus distinguishing wantedindications from background fluorescence. In a case where an indicationextends over several matrix elements, this method would cause the signalto be treated as due to background and thus wrongly accept the testobject. Provision is made for such an eventuality by writing into theprogram an overriding absolute ceiling value, which would cause thesystem to trigger on a signal of that intensity being detected, irrespec-tive of the intensity of the signals of the surrounding area.

The outline of the test object is detected by separate photomultiplierssensitive to the incident wavelength reflected by the back of theinspection chamber when the scanning beam passes over the edges of thetest object. The image of the test object thus produced is displayedon a television monitor in real time, and detected indications are madeto pulsate in their true locations relative to surface geometry,facilitating subsequent more detailed inspection by other techniques.The images may be stored on a disc and re-displayed on a separatemonitor as an aid to conventional visual inspection under black light.

AEOS has been subjected to intensive evaluation in aero engine plants.Compressor and turbine blades from production lines were double inspectedby the system and skilled inspectors, and the findings recorded andcompared. The trials demonstrated that the system will detect consis-tently very small indications in the presence of normal levels of back-ground fluorescence. Therefore it satisfies the two major requirementsfor any NDT procedure, sensitivity and repeatability, and it eliminatesreliance on the human eye for the detection of indications, thus improv-ing greatly the reliability of FPI.

AEOS may be used to detect fluorescent indications irrespective of themethod by which the indications are produced. Thus ferromagneticobjects, treated by a fluorescent magnetic particle process are equallysuitable subjects for AEOS.

225

CRYOFORMING OF 301 AND 302 STAINLESS STEEL

T. Livni, S. Bar-Ziv, A. Rotem. Rafael, Haifa

A. Rosen, Technion, Haifa

INTRODUCTION

Metastable austenitic stainless steels transform to martensite whenbeing plastically deformed. The volume fraction of martensite is afunction of increasing deformation and decreasing temperature.At cryogenic temperature, relatively small amount of deformation isrequired in order to acheive almost a complete transformation Fromaustenite to martensite. This results in a very drastic increase in theyield strength of the material. The effect in indeed used inmanufacturing high strength components with good corrosion resistance[1,2]. The process is known as "Cryoforming" and is used for producingpressure vessals and springs.

The aim of this investigation was to establish the relationshipbetween the parametersof the process and the mechanical and some corrosionproperties of 301 and 302 steels.

EXPERIMENTAL

The experiments were carried ût on three types of specimens:

A. Flat tensile specimens, made of AISI 301 sheet, (1.6mm thick, 6.25mmwide, G.L.=35mm). Specimens were taken in the longitudinal direction,and were vacuum annealed. (1010°C, 5min, Argon cooling).

B. TIG Weidments(both longitudinal and transeverse) were made on piecestaken from the same sheet tensile specimen were prepared, the sameway as in. A.

C. Round tensile specimens, <j>6mm, were prepared from AISI 302 annealedrods.

All specimens were plastically deformed,in tension,to various amounts(10 to 21 pet.). Cryogonic tension device, containing liquid Nitrogen(-ig6°C) was used (Combmed with an Instron machine). Two strain rateswere employed - 0.25x10 2 and 0.25x10 3sec"1. Following this "Cryoforming",martensite content was measured by x-Ray diffraction method, and tensileproperties of all specimens were tested in room temperature.

Cryoforming to 13 and 17 pet. elongation in type A and C specimens,respectively, resulted in optimal combination of tensile propertiess(cTy=l50kg/mm2, P+=10%). Fracture toughness (on type A specimens), Saltspray corrosion resistance (on type A and B) and notch senstivity -NTS/oy (on type C) tests were carried out on specimen cryoformed toacheive the above mentioned properties.

226

RESULTS

Tensile tests - uriwelded material

Fig. 1 demonstrates a typical stress-strain curve of specimen whilebeing deformed in L.N., together with the variation of martensite contentwith the plastic strain. Immersion in L.N. resulted in imidiate formationof 5 pet. martensite. No increase in martensite content was detected inspecimen deformed bellow 2.6 pet. Further elongation causes martensltictransformation. After 20 pet. elongation, an almost fully martensiticmaterial is accepted.

Room temperature tensile properties of specimens deformed cryogenicallyto various amounts of elongation are shown in Fig. 2. Since the strainrate during the cryoforming process had only a negligable effect on theseproperties, Fig. 2 demonstrates results obtained from cryoforming in onestrain rate (0.25x10 ^ " 1 )

The insensitivity of the martensitic transformation progress to strainrate in this steel was confirmed by Bakofen et al [5] for room temperaturedeformation.

180-

170-

150-

IJO-

O3 S IO 15 2OCRYOG€NIC PLASTIC DEFORMATION |%|

T ' ' ' ' I 1—1—r

Fig 1. Typical Stress-Strain Curve ofCryogenic Deformation and theRespective Martensite Content.

Fig 2. Room Temperature TensileProperties of Sheet (Type A)Specimens Cryoformed to

: Various Amounts ofElongation.

227

Few conclusions can be drawn from Fig. 2. Increasing the amount ofcryogenic deformation, increases strength levels, and decreases the gapbetween 6y and ^ u , e.g. s t ra in hardening coeff ic ient in reduced.

Cryogenic deformation bigger than 16 Pet. of sheet (type A) specimensresults in very poor ductility at room temperature.

For the reasons mentioned above, optimal elngation for Cryoformingprocess was chosen to be 13^14 pet. for strain rate of 0.2x10 2sec l.

It is to be noted that using the suggested cryoforming process, tensileproperties similiar to those of Cjstom 455 (H-950 condition) steel can beobtained.

Welded specimens

Room temperature tensile properties of welded specimens are shown inFig. 3.

Oj; o»

190-

180-

170-

160-

150-

140-

130-

120-

. 6r

7 8 9 10 11 12 1J 14 .K 7°

Fig,3 : Room Temperature Mechanical Properties of Welded Specimens, vs.Cryogenic Elongation.

For determination of yield strength, elongation was measured withextensiometer fixed in the close vicinity of the welded zone. Thedependence of mechanical properties on the amount of cryogenic deformationwas found to be similar to that obtained for unwelded specimens. Themajor difference was that slightly smaller cryogenic deformation in weldedspecimens was required in order to obtain a given yield strength.

228

It is to be noted that welding has no deterious effect on mechanicalproperties. For a given value of yeild strength, similiar values ofelongation are obtained, both for welded and unwelded material.

Specimens were fractured randomly - inside or out of the weld zones.

Toughness and corrosion data

Fracture toughness (Kc) and notch sensitivity (NTS/ay) data wasobtained, for specimens crvoformed to yield strength of 150kg/mm2.NTS/ay tests were carried out on type C specimens, Kc tests were performedon type A specimens, fatigue precracked in L-T direction.

Results are summerized in Table 1.

Cryoformed Steel Custom ^55 Steela =150kg/mm (H-950)

MTS/ay (Kt=A.7) l^7 1.53

K m [ (kg/mm2) l'mm] 100 260

Salt spray corrosion tests were carried out on annealed material,and cryoformed specimens (both welded and unwelded). No difference wasobserved in the corrosion behaviour of the different specimens after 3month period.

DISCUSSION

Tensile properties of the cryogenically formed steel are clearly afunction of the martensite content.

In the current literature there are two different approaches todescribe the strengthening effect of martensite in partially transformedsteels. One approach claims that the martensitic phase acts as dispersoid,pinning dislocations and creating unsurmountable barriers [3]. The otherapproach simply assumes that the anstenite-martensite mixture behaves likea composite material. In order to examine which of the above theoriesapplies, room temperature yield strength of cryoformed specimens wasplotted vs. the volume fraction of martensite, see Fig. k.

The linear relationship between yield strength and martensite contentis clear, and can be described in the following equation, which appliesto composite materials:

<j ;CT - yield strength values of martensite and austenite, respectively,

f - volume fraction of martensite.

229

wo-

Fig. k: Yield Strength vs. VolumeFraction of Martensite inCryoformed Specimens.

*>-

—r~20

1 1—to 4o

V . i l i i i i . I r . n r i

T~

ii • " " "'.in

—IUO

The martens!tic transformation was found to be strain induced, sinceno change in martensite content was observed in the elastic region.

The instability of austenite - i.e. the ease of cryoforming inducedtransformation can be derived from the deformation - transformationfunction, which can be expressed by the following equation [A],

f

1-f=AeB

where f,,, is the volume fraction of martnesite, e is the amount of plasticstrain, and A and B are material's constants.

"A" represents the ease with which on austenite structure can undergoa strain induced transformation to martensite, and "B" represents themeasure of the autoc^talitic nature of the transformation.

Fig. 5 exhibits the variation of fm/'-fm vs. e in a log-log scale.The plot is indeed linear, and the values of material's constants are:

B=2.38

230

Fig. 5: log-log Curve of theVariation of MartensiteContent with CryogenicStrain.

A=45OB = slope=2.38

C o l O.1 1.l o g S T R A I N

CONCLUSIONS

The amount strengthening o f metastable s ta in less s tee l is governedby the volume f r a c t i o n of mar tens i te , formed during cryogenic deformat ion.The dependence o f room temperature y i e l d st rength on mar tens i te content is1inear.

The pr inciple results of the investigation are:

1) There is a poss ib i l i ty to obtain high strength steels with re lat ive lygood d u c t i l i t y and toughness, combined with good corrosion resistancecharacterist ics.

2) Welding has no detertous effect on mechanical properties of cryoformedmaterials.

3) The cryoforming process can be controlled by the amount of cryogenicdeformation, in order to achieve desired mechanical properties.

REFERENCES

[ I ] Packner, Bernste in - Handbook o f Stain less S tee l s , McGrow-Hill 1977,p. 47-6.

[2] NASA Cr 61251.

231

[3] Mangonon P.L., Thomas J.G., Met. Trans. Vol. 1 1970, p. 1577-

[k] Ludwigson D.L., Berger J.A., J.I.S.I. 1/1969, p. 63.

[5] Powell, Marshall, Backofen - Trans, of ASM Vol. 50, 1958.

232

AUSFORMING OF H-ll MOD. STEEL

G. Rlkabir* and A. Rosen**

* RAFAEL, Armament Development AuthorityP.O.B. 2250, Haifa

** Dept. of Mat. Eng. - Technlon

INTRODUCTION

Ausforming is a Thermo-Mechanical Treatment (TMT) intended to improvemechanical properties of steels. Generally, ausformed steels have abetter combination of strength and fracture to toughness. In addition,ausforming allows use of relatively simple steels instead of expensiveones, such as maraging steels.

The aim of this investigation was to study the ausforming process inorder to use it in the future in manufacturing real parts such as rocketmotor casings. The H-ll Mod. steel was selected since its use in ausfor-ming technology in widely known. Although there is no plan today to usethis technology in producing thin walled tubes due to lack of very expen-sive foundations, it is a very interesting and important alternative.

EXPERIMENTAL

H-ll Mod steel, having the composition of 0.38%C, 4.94%Cr, 1.20%Mo,0.60%V, 0.17%Ni and 0.79% Si have been supplied in the form of a 10 mmthick plate. After the standard acceptance tests the plate was cut intopieces of 120imr. long and 45 mm wide. Ausforralng was performed by meansof hot rolling. The details of the process are listed below:

a) Austenization treatment (20 min. at 1050°C).b) Fast cooling to 530°C.c) Hot rolling at 530°C to 40%, 60% and 80% reductions.d) Oil quenching.e) Double tempering (60 min. at 450°C).

All the platelets were grinded to a final thickness of 1.3mm in order toobtain scale and decarb free, flat and smooth surfaces.

The following tests were performed:

(i) Hardness measurements, (ii) TensLle tests, (iii) Metallography and(iv) Fractography.

For the sake of comparison two alternative treatments were also carriedout:

(a) Oil quenching from 1010°C, double tempering at 510°C for 2 hourseach. According to the literature this treatment results in maximumstrength.

233

(b) Oil quenching from 1010°C, double tempering at 620°C for 2j hours toa hardness of 35 Re and cold rolling to 40%, 60% and 80%reductions.

RESULTS

Mechanical properties.

The following table suramerises the average properties for the threedifferent treatments. For the case of ausforming and cold rolling treat-ment we report only the results obtained after 80% reduction, since thelower reductions gave inferior results.

Treatment

Ausforraing

Quenchingand Tempering

Pre-temperedand Cold rolled

M?a

2530

1980

1580

°uMPa

2610

2040

1600

£

%

3

5.5

5.5

HardnessRe

65

56

48

Metallography

The microstructure of the quenched and tempered specimens is normal tem-pered martensite. The former austenitic grain size was ASTM 5-7. Nospecial features were observed. The pre-tempered and cold rolled speci-mens exhibited a fiberous, heavily deformed microstructure. The aus-formed specimens showed a rather complex microstructure of branched de-formation bands and fine dispersion of carbides. It was impossible tomeasure former austenitic grain size.

Fractography

The fracture surface of the quenched and tempered specimens can be cha-racterized as quasi-cleavage, i.e., a mixture of smooth cleavage surfacesand wide, shallow dimples. The pre-tempered and cold rolled specimfiisexhibited a typical brittle fracture, characterized by severe delamina-tion. The fracture surface of the ausformed specimens were similar tothat of the cold rolled specimens, however, delamination was less severeand the fraction of the dimpled areas was larger.

DISCUSSION

This investigation has proved without doubt that ausforming can signifi-cantly improve the strength of H-ll Mod. steel compared to standard heattreatment, or to pre-tempering and cold rolling. Actually, the yieldstrength and UTS obtained after ausforraing of this steel Is higher thanthat of raaraging steels.

234

The increase of strength is accompanied by a certain reduction of ducti-lity which is probably unavoidable at these high strength levels. On theother hand it is possible to temper the ausformed steel at a higher tem-perature than was done in this research (450°C) and obtain higher elonga-tion on account of losing strength.

At this stage of the investigation we do not have unough structural evi-dence which can point out the reasons for the improvement of strength.One has to study the micros tructure in much finer details, e.g., bytransmission electron microscopy, which is quite out oE bounds for thisproject.

The results reported in this paper give the best combination of strengthand ductility. Much more combinations were studied, such as variousdegrees of reductions by rolling, different quenching as well as ausfor-ming temperatures and finally various tempering temperatures. The impor-tant fact is that variation of the above factors did not result in markedchanges of strength. This fact is encouraging since upscaling of theprocess will require certain tolerances of heat treatments and deforma-tion temperatures.

CONCLUSIONS

The recent investigation of TMT of H-ll Mod. steel reviewed the followingtendencies:

1. Ausforraing under certain conditions improves both yield and ultimatetensile strength by 30% but causes an almost 50% decrease in tensileelongation, compared to quenched and tempered material.

2. The microstructures and fracture surfaces of ausforraed or quenchedand tempered materials are similar.

3- Small variations of ausforming conditions do not effect the end-strength of the specimens.

235

THERMODYNAMIC AND KHST3TIC PHENOMENA IN ADSORBED LAYERS

M. Grunze

Laboratory of Surface Science and Technology andDept. of Physics and Astronomy, University of Maine

Orono, Maine 04469, U.S.A.

and

Fritz-Haber-Institut der Max-Planck-GesellschaftFaradayweg 4-6, 1000 Berlin 33, West Germany

ABSTRACTThe relationship between thermodynamics and kinetics of adsorption

will be outlined for the system N_/Ui(11O), for which an extensive studyinvolving several modern surface science techniques were performed overthe last few years. With respect to kinetics of adsorption, a briefsummary of "precursor" concepts and their experimental tests will be givenfor molecular nitrogen adsorption on nickel and rhenium surfaces. Finally,seme recent experimental and theoretical results for the identification ofthe intermediate states in nitrogen dissociation on an Fe (111) -surfacewill be presented.

INTRODUCTIONThe elementary steps in the interaction of gases with solid surfaces

are studied in many laboratories to provide an understanding of the basicphysical and chemical phenomena relevant to modern technology, e.g. heter-ogeneous catalysis, materials modification, corrosion resistance or micro-electronics. Equilibrium measurements of the surface coverage in therespective gas ambient are carried out to determine the heat of adsorption(Q) and the entropy in the adsorbed layer (S , ) , both important quantitiesin the description of the adsorbate phase. TRe isosteric heat of adsorp-tion (Ogm) is derived by equilibrating the chemical potentials of the gasphase (y ) with the chemical potential of the adsorbate phase (y^) andsolving for the temperature dependence of the equilibrium pressure at con-stant coverage 8/1/:

0

= ^ § . (1)Equation (1) applies whether or not the properties of the substrate

change upon adsorption, since all changes in enthalpy of the system areincluded in Q. However, thermodynamic equilibrium between the solid andthe gas has to be established, which also requires that gas and solidhave the same temperature. This requirement is often neglected in ad-sorption studies on single crystals, although it was discussed in theearly literature by Ehrlich /2/. As shown elsewhere, using kineticarguments /3/, the error introduced by a temperature difference betweengas and solid is typically small ($ 2 % of Q ), provided the stickingcoefficient of the gas does not depend on gas temperature.

Fran the isosteric heat of adsorption, the temperature of the system

236

and the equilibrium pressure the differential (and integral) entropy ofthe adsorbate layer can be calculated. The differential entropy is givenby

bad = Sg_E - (2)

nea'.ecting again a possible temperature gradient between gas and solid /A/.The entropy in the adsorbed layer is a particularly useful quantity, sinceit allows a comparison with statistical mechanical models of the adsorbatephase /1/.

In steady state equilibrium between solid and gas the rates of ad-sorption and desorption have to be equal. The respective activationbarriers for adsorption and desorption determine the adsorption energy.Fran the knowledge of the absolute adsorption and desorption rates, theequilibrium constant K , and thus the free energy of adsorption can becalculated. Thus, the Jcmetic results can be compared to the equilibriumdata.

The above thermodynamic and kinetic quantities include only indirectinformation on the interaction between the gas species and the substrateand on lateral interactions between the adsorbed particles. For a com-plete description of the adsorption system and a theoretical approach in amodel calculation, complementary data on the structure of the adsorbatephase and the bonding geometry of the molecules are required. This infor-mation can be obtained by surface sensitive spectroscopic techniques suchas low energy electron diffraction (LEED) /5/, x-ray and uv-photoelectronspectroscopy (XPS and UPS) /6/ or vibrational spectroscopy (Infrared-Reflection Absorption Spectroscopy, IRAS /!/ or high resolution electronenergy loss spectroscopy, HRELLS /!/).

MOLECULAR NITROGEN ADSORPTION ON Ni(llO)Molecular nitrogen adsorption on Ni (110) has been studied by many

modern surface science techniques over the past years and probably is bynow one of the most thoroughly studied adsorption systems, not consideringthe extensive data available for OO-adsorption on metals. Although mole-cular N2 adsorption on Ni(110) might seem to have no direct technical re-levance, the adsorption of N_ comprises one of the elementary reactionsteps in the ammonia synthesis reaction or in the nitride formation ofmetals. We will discuss this adsorption system in order to outline thethermodynamics and kinetic phenomena occurring in adsorbed layers.

In Fig. 1, we show the isosteric heat of adsorption determined byeq. 1 from several sets of adsorption isobars using the N1s core levelphotoemission band as a monitor of surface coverage /8/. In addition,the change in work function, the attenuation of the Ni-d band emission dueto N 2 adsorption, and the decrease in the low energy ion scattering signalfrom the Nickel substrate upon N^-adsorption were used as N 2 coveragemonitors to record adsorption isobars /9/. The isosteric heat of adsorp-tion determined with the above techniques (except for the work functionresults /9/) falls well within the error bars of the N1s-core level data,and thus gives a consistent picture of the energetics of adsorption. Asshown in Fig. 1, Q remains approximately constant (Q ~42 kJ/mole) upto e =0.47, and tfien drops steeply to a value of Q _ ~ 20 kJ/mole. Thecoverage calibration was obtained from LEED, the steep drop in Q occursat completion of a (2x1) over layer. This (2x1) phase then transforms,with increasing coverage, into an incannensurate c( 1.4x2) solid phase viaan intermediate fluid phase /10/. Such a commensurate-fluid-inconitiensurate

237

Fig- 1. Q S Ton Ni(11O).

for N-

60-

iO-

o

40-

30-

20-

10-

0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8

solid phase transitionon a rectangular sub-strate lattice waspostulated by recenttheories of two-dimen-sional phase transi-tions /11/, and wasfirst observed experi-mentally in the KL/Ni(11O) system/10/.Thus, the change inQ S T at 6 ~O.5 is as-sociated with a two-dimensioual phasetransition and we can expect this phase transition to be also reflectedin the entropy of the adsorbed phase and in the kinetic data. ^

Figure 2 shows the differential entropy in the adsorbed N- layer S ,(eq. 2) as a function of coverage for Ni(11O) /9, 12/ and Ni(10O) /3/. teinclude the KL/Ni(1OO) data /3/ to demonstrate that the entropy reflectsthe specific changes in degrees of freedom in the adsorbed phase and isthus sensitive to the structure cf the substrate and the change of lateralinteraction between the adsorbed molecules. On Ni(1OO) , the isosteric heatis similar to bhat on Ni(11O) up to a coverage of 5.5x10 molec/cm(corresponding to 8 = 0.5 on Ni(11O)). It remains constant almost up tosaturation coverage, when a drop in Q__ is observed, most likely due tothe formation of antiphase domain walls in the c(2x2) over layer /3/.

As mentioned in the introduction, the entropy in the adsorbed phasecan be correlated with statistical mechanical models /1/. We plotted thedifferential entropy calculated by assuming /3/

a) a lattice gas model, considering only the configurational entropyin the adsorbed layer plus the vibrational entropy associated with the sixnormal modes of a N_-Ni surface complex (A) and

b) a two-dimensional (mobile) Volmer gas, including the vibrationalmodes of the individual molecules (o).Clearly, both model calculations do not explain the experimental data, inparticular the high coverage region where the commensurate-fluid phasetransition occurs. Presently, there are no statistical models which ex-plain in detail the high entropy in the fluid phase. For the coverageregion below 6 ~ 0.5, we concluded that collective excitations in the ad-sorbed phase contribute to the high entropy, since considering only thedegrees of freedom of individual molecules cannot explain the experimentalresults /3/.

The thermal desorption data recorded for the N?/Nj|11O) system show asingle desorption peak up to 6 = 0.5 /9,13/. At higher initial coverages,a second peak at lower temperatures develops indicating a lowering of theactivation energy of desorption. The activation energy of desorption upto 3 = 0.5 is, v/ithin the error bars, identical to the isosteric heat ofadsorption, which means that adsorption of N ? is a non-activated process.For the higher coverages, present data are not accurate enough for a com-

238

Fig. 2. Differential entropyin the adsorbed layer for N- onNi(110)(«), Ni(100)(°) andstatistical mechanical calcula-tions (see text).

parison of activation energy ofdesorption and isosteric heat.By integration o± the desorptiontraces, a plot of coverage versusexposure can be made, from whichthe sticking coefficient, s =f(9), can be calculated;

S(6) =893Ex

250-

225-

200

175-

_ 150-

125-

100-

75-

50-

25-

0.8 1.6 2.4 3.2 4.0 4.8 5.6 6.4 7.2 8.0molec

cm2

where Ex is the integrated flux ofmolecules impinging onto the sur-face (exposure).

In Fig. 3, the stickingcoefficient as a function of ex-posure is displayed. S is con-stant up to 9 = 0.48, then dropsto a value of s ~ 0.3, increasesagain and diminishes at satura-tion coverage. This behaviour is not only observed at low temperatures,but is also evident at all temperatures where the catmensurate-f luid phasechange is observed in the LEED experiments. At the higher temperatures ex-periments involving adsorption and subsequent desorption cannot be carriedout, since the residence time of the molecules on the surface becomes muchtoo small. Therefore, we used detailed balancing arguments to extract thesticking coefficient from equilibrium measurements. Under equilibriumconditions, the rate of adsorption equals the rate of desorption:

rad = s(e)

r, = v-N ,des ad (3)

Fig. 3. Sticking coefficient forN 2 adsorption on Ni(HO) as afunction of coverage (T = 87 K).

s

1.0'

CD

* i i

} • • ,.5 1

0 •

239

and TEqua-

0

Fig. 4. Sticking coefficient of ôn Ni(110) as obtained from isobars

+p = 1x10~ mbar,

= 5x10~ mbar,

*" -7on = 1x10 mbar.

The adsorption rate is given by theimpingement rate multiplied by thesticking coefficent; the desorptionrate is expressed1 in the usualArrhenius form. An isobar gives thecorrelation between P, N(T = substrate temperature).tiSn 3 can be solved for s(6) and,knowing the desorption rate para-meters exactly, s(6) can be calcu-lated. Fig. 4 shows the result fors(8)(note, that s(8) is not an iso-thermal quantity anymore!) obtainedfrom three isobars. The agreementwith the s(8) data from the desorp-tion experiments is surprisinglygood, considering that the data weretaken in different UHV chambers andno normalization of p, T or N , wascarried out for the two ixperiments.In particular, we note that at 9~ 0.5 also under equilibrium conditions adiscontinuity in s(e) occurs, which thus is clearly related to the con-mensurate-fluid phase transition.

The constant sticking coefficient up to saturation coverage can beexplained by assuming a mobile "precursor" state, i.e. a molecule imping-ing onto the surface is not reflected into the gas phase, but moves overthe surface for a sufficient time to accanodate its translational and in-ternal energy (T = 300 K, T < 200 K) and becomes chemisorbed. A dis-tinction is made^Between "intrinsic" and "extrinsic" precursor states, theformer trapped over the bare surface, the latter on top of a chemisorbedlayer. The kinetic formalism relating these precursor states to theexperimental sticking coefficient has been reviewed in several articles/14,15/. A spectroscopic characterization of these precursor states wasonly carried out at low temperatures for a few adsorption systems wherethe chemisorbed state is formed by dissociation of the respective gases/16,17/. For dissociative adsorption the "precursor" state is identifiedas the molecular state, which loses its translational and internal energybefore dissociation.

The results of our study on precursor states for molecular N^ adsorp-tion on Ni(100) and Re(0001) /18/ are schematically shown in Fig. 5.These experiments were carried out at T ~ 20 K, as a spectroscopic tech-nique, XPS, was applied. On the clean metal (5a), the molecule adsorbs andconverts within the time it takes to record a spectrum (ca. 5 min.), intothe linearly bonded chemisorbed state, which is the state observed alsoat the higher temperatures. This result implies, that a possible activa-

240

(hi N 9 GasN 2Gas

(b) 2

1

Gas

Phys. N 2 Layer

Fig. 5. Schematic model for precursor state in molecular N« adsorptionon Nickel and Rhenium surfaces (see text).

tion barrier between a metastable, frozen-in state (e.g. a molecule lyingflat on the surface) has to be less than E ~ 3 kJ/mole /18/. As long asthe coverage in the chanisorbed state is less than 2/3 of saturation cov-erage (erei <2/3} i molecules impinging on occupied sites will find emptychemisorption sites during the time they need to accanodate all theirkinetic and internal energy, i.e. s remains constant. For 0 > 2/3,molecules are found to condense on top of the chemisorbed layer (5c) i.e.they lose their energy in the second layer before they find a chemisorptionsite on the metal surface. This results in a decrease of the stickingcoefficient into the chemisorbed state, but the net trapping coefficient(chemisorption and condensation) ranains unity, and multilayer formationtakes place (5d)„ Completion of the chemisorbed layer is hence onlyobserved at exposures exceeding those necessary to complete a monolayer,at T > 80 K. Thus, a molecule chemisorbed in the second layer acts as an"extrinsic precursor". The heat of adsorption of this precursor in thesecond layer is about twice the heat of bulk condensation /19/, whichincreases the residence time on the surface considerably above the valueexpected for a purely condensed species.

Our low temperature experiments explain the observed constant andhigh sticking coefficient for N 2 on Ni(11O) up to half monolayer coverage.The discontinuity at 9 > 0.5 cannot be explained within the framework ofpresent precursor models, since all these models would give a continuouschange in s(6). Presently, computer simulations relating the kinetics ofadsorption and desorption to the geometry and energy changes in the ad-sorbed layer at 6 > 0.5 are being carried out to develop a microscopicmodel for the behaviour of s(e).

N2 DISSOCIATION ON Fe(lll)Ammonia synthesis on iron is a structure sensitive reaction with the

Fe(111) surface being the most active low index single crystal plane /20/.Using various methods, it has been demonstrated for clean iron surfaces,

241

that the rate determining step for the overall reaction involves nitrogendissociation, since the subsequent hydrogeneration of atonic nitrogen toammonia proceeds at much higher rates. The kinetics of dissociative nitro-gen adsorption on Fe(111) have been studied by Ertl et al. /21/ in detail,and it was concluded that this process proceeds through a molecularly ad-sorbed "precursor" state (α-state) with a rather low activation barrier.Recently, it was shown /22/ that this α-state is preceded by a more weaklyheld physisorbed γ-state, which desorbs already at about 85 K. More im-portantly, the XPS/HRELLS study summarized below showed that the α-stateis a ir-bonded molecule, with both N atoms interacting with the surface andthus forming the immediate precursor to dissociation /24/.

Fig. 6. N 1 S core levelspectrum for a: the α-state,b: γ-state, c: atonicβ-state.

-

-

tn .

CD

[a

rb

.

J -a> -c

c -

c

b

a . -.' /

405 9 ...

' " ' ' " ' 39

397 0 '-"-" •-:-••'.•,.

4012

412 406 400

energy belou E-FermI [ eU)

394

In Fig. 6 we show thecharacteristic N1s spectraof the three distinguishablenitrogen states on Fe(111).Spectra 6a shows the doub-let structure of two bandsat 405.9 and 401.2 eV be-low E ,, observed for theweakly bonded γ-state. Theobservation of a doublet structure for a single species is due to finalstate effects in the photoemission process; the peak at lower bindingenergies corresponds to emission from a "screened" final state, the oneat higher binding energy to emission from the "unscreened" final state/23/. A comparison of spectrum 6a with those of N_ bonded on Ni(11O)/8/ orRe(COO1)/18/ shows that intensity is transferred to the "unscreened" finalstate peak, indicating a weak interaction with the substrate which is alsoevident from the low heat of adsorption of Q $ 24 kJ/mole /22/. This T-state then transforms, as a function of time and temperature, into the «-state producing the XPS-spectrum shown in Fig. 6b. Isotopic exchange ex-periments show /21/ that, in the α-state, the molecular unity is pre-served. The XPS-data and the HREELS results ,/24/ reveal, however, a stronginteraction of the two nitrogen atoms with the surface. A comparison ofthe v

N stretching frequency observed by HREELS to those in -bonded

inorganic dinitrogen complexes, .as well as CMDO-CI calculations of the N1sspectrum show that a) the molecule bonds with its molecular axis parallelto the surface, b) that bonding to the surface involves charge donationfrom the 1 IT orbital of the molecule to the substrate and a 1 " * back-bonding from the substrate to the molecule and c) that the charfe transferinto the antiJDonding (with respect to the dinitrogen bond) T^* orbitalincreases the N-N bond length to ~1.25 A , compared to the N-S bond dis-tance of 1.1 A in the free molecule. Thus, the effective activation bar-rier for dissociation is lowered and dissociation of N_ is observed atlow temperatures (T ~ 120 K ) , leading to the formation of the character-istic N1s-core hole emission of atonic R nitrogen (Fig. 6c).

242

Fig. 7. Potential energy diagramfor N2-dissociation on Fe(111) asderived frcm experiment. inn]

ft 2N.,From a careful study of theN1 S spectra as a function of gaspressure, temperature and time, apotential energy diagram cis shownin Fig. 7 can be constructed. Inthe upper part, we show a schema-tic view of Y, a and atomic g ni-trogen on the Fe(111) surface,the potential energy diagram in-cludes the respective heats of ad-sorption and activation barriers.The individual rate constants forthe reaction are indicated in theinsert.

The microscopic mechanism ofdissociation of N~ on Fe(111) wasalso investigated theoreticallyby calculating the potential ener-gy surface for this reaction /25/. The theoretical treatment is describedin ref. /26/. It involves calculation of the total energy E of thesystem as a function of the height, h, of the molecule above the surfaceand the intramolecular distance, d /26/. The Fe(111) surface is approxi-mated by a four atom cluster and the molecule is lowered onto the sur-face with different orientations of its molecular axis with respect tothe substrate. The bestagreement between theoryand experiment is ob-tained when the molecule be-comes adsorbed with itsaxis parallel to the surfacein the precursor state fordissociation (Fig. 8).

N2/FE(1r

Fig. 8. Potential energysurface, E (h, d ) , for Ndissociationon Fe(111). Thedissociation path is indi-cated by the arrows, theequipotential lines have aseparation of 0.2 eV.

0.5

243

In this intermediate a-state, charge is transferred into the \ * or-bital of

QN

2, leading to a calculated increase of the N-N distance of

d 1.5 A.TSy surmounting a potential energy barrier of E . ^ RJ0.2 eV =20 kJ/mole (compared to experimental E ^ = 28 kJ/mole) ,^he moleculedissociates and forms the atomic β-state. From this calculation it alsofollows that the height of the activation barrier, E^. , is critically de-pendent on the charge transfer into the TIT * orbital O P N

2. Charge dona-

tion into this antlbonding molecular orbital by substrate electrons weakensthe intramolecular bond, and thus decreases the activation barrier for dis-sociation.

The simple dissociation reaction discussed above is the first examplein which the microscopic mechanism of the heterogeneous fission of an in-tramolecular bond has been identified both experimentally and theoretical-ly. From a general point of view, it may be concluded that partial occu-pation of antibonding molecular orbitals by substrate electrons comprisesone of the crucial steps in dissociation. Details of the geometry and elec-tronic overlap in the "precursor" of dissociation will, however, depend onthe particular system under study.

ACKNOWLEDGEMENTSThe experiment described were performed at the Fritz-Haber-Institut

der Max-Planck-Gesellschaft and at the Free University of Berlin. Finan-cial support from the Deutsche Forschungsgemeinschaft, Sonderforschungsbe-reich 6, is gratefully acknowledged.

REFERENCES1. A. Clark, in "The Theory of Msorption and Catalysis", Academic Press,

1970, and other textbooks.2. G. Ehrlich, J. Chem. Phys. 36 (1962) 1499.3. M. Grunze, P.A. Dowben, and R. Jones, Surf. Sci., in press.4. M. Procop and J. Vblter, Surf. Sci. £7 (1975) 514.5. see for example: M. Henzler, "Electron Diffraction and Surface Defect

Structure", in: "Electron Spectroscopy for Surface Analysis", ed.H. Ibach, Topics in Current Physics, Springer Verlag 1977.

6. B. Feuerbacher and B. Fitton, "Photoemission Spectroscopy", ibid.7. H. Froitzheim, "Electron Energy Loss Spectroscopy", ibid.8. M. Golze, M. Grunze, R.K. Driscoll, and W. Hirsch, Appl. Surf. Sci.

6 (1980) 464.9. M. Grunze, M. Golze, and W.N. Unertl, to be published.10. M. Grunze, P.H. Kleban, W.N. Unertl, and F. Rys, Phys. Rev. Lett.

51 (1983) 582.11. see for example: S.N. Coppersmith, D.S. Fisher, B.I. Halperin, P.A.

Lee, and W.F. Brinkmann, Phys. Rev. Lett. 46 (1981) 549.12. M. Grunze, M. Golze, and W.N. Unertl, Proc. 3rd Symp. on Surface

Science, Obertraun/Austria 1983, ed. P. Braun, HTU-Druck, Wien, p. 218.13. M. Golze, M. Grunze, and W. Hirschwald, Vacuum 31_ (1981) 697.14. D.A. King, in "Chemistry and Physics of Solid Surfaces", ed. R.

Vanselow, Vol. II, CRC Press, (1979).15. A. Cassuto and D.A. King, Surf. Sci. 102 (1981) 388.16. G. Ehrlich and F.G. Hudda, J. Chem. Phys. 35 (1961) 1421.17. R. Gcmer, in "Field Emission and Field Ionisation" Harvard University,

Cambridge, MA. 1961.18. M. Grunze, J. Fuhler, M. Neumann, J. Behm, C.R. Brundle, and D.J.

Auerbach, Surf. Sci., in press.19. J. Fuhler, Diplcmarbeit, University of Osnabruck, 1984, and to be

244

published.20. For a review, see: M. Grunze, in:"The Chenical Physics of Solid Sur-

faces and Heterogenous Catalysis", ed. D.A. King and D.P. Woodruff,Vol. 4, Elsevier (1982), p. 143.

21. G. Ertl, S.B. Lee and M. Weiss, Surf. Sci. 114 (1982) 515.22.a)M. Golze, W. Hirschwald, M. Grunze, and M. Polak, Proc. 3rd Symp.

on Surfac3 Science, Obertraun/Austria 1983, ed. P. Braun, KTU Druck,Wien, p. 254.

b)M. Grunze, M. Golze, J. Fuhler, M. Neumann, and E. Schwarz, Proc. 8thInt. Congress of Catalysis, Berlin (1984), in press.

23. K. Schonhammer und 0. Gunnarson, Surf. Sci. 89 (1979) 573.24. M. Grunze, M. Golze, H.-J. Ereund, H. Pulm, U. Seip, M.C. Tsai,

G. Ertl, and J. Kiippers, submitted.25. D. Toraanek and M. Grunze, to be published.26. D. Totnanek and K.H. Eennemann, Surf. Sci. 127 (1983) L111.

245

SUPPORTED SILVER CATALYSTS HAVE SOME IMPORTANT PROPERTIES INCOMMON WITH ROUGH SILVER FILMS: ELLIPSOMETRY AND RAMAN DATA.

P.H. McBreen*. D. Hall**, J. Lalman and M. Moskovits

Department of Chemistry, University of Toronto,Toronto M5S 1A1, Canada.

*Department of Materials Engineering, Ben-GurionUniversity of the Negev, Beer Sheva, Israel.

Deparment of Chemistry, Ben-Gurion Universityof the Negev. Beer Sheva, Israel.

ABSTRACT

Rough silver films absorb much more visible light than do smooth silverfilms. Ellipsometry data for this phenomenon may be interpreted by attri-buting a strong optical absorption feature, centered at 550nm, to theoptical properties of a surface roughness layer. This feature is at leastpartially the reason why the Raman signal for adsorbates on roughenedsilver substrates is often up to a million times larger than expected. Theso-called SERS (Surface Enhanced Raman Spectroscopy) effect is used inthe present study to observe the behaviour of carbon on the surface ofsilver catalyst particles.

INTRODUCTION

Six years ago the surprising discovery was made that Raman spectroscopycould be used to study adsorbed species. Previously it was believed thatthe combination of the low cross-section for Raman scattering and thesmall concentration of molecules probed at the surface (10 /cm ) wouldpreclude its use as a routine surface spectroscopic technique. However,it was found that on some surfaces the observed Raman scattering signalassociated with an adsorbate was often a million times larger thanexpected. This enormous effect is exploited in surface enhanced Ramanspectroscopy (SERS), and has been the subject of intense study in recentyears. In a SERS experiment a monochromatic beam of laser light isdirected at the sample. Some of the scattered light is collected and isanalyzed on the low frequency side of the incident beam. Observed sharpspectral features may be directly related to the excitation of molecularvibrations through the Raman scattering process. Such spectra may be usedto determine the identity and the structure of adsorbed species. Since,in general, SERS involves a beam of visible light as a probe it has theimportant capability of being suitable for in-situ studies of electrodeor catalyst surfaces. The magnitude of the surface Raman enhancement isdependent on the substrate metal; the effect is most pronounced foradsorbates on silver. Further, there is a correlation between SERSactivity and the structure of the substrate. Silver films deposited ontoa cold ( lOOK) substrate are SERS active, whereas continuous silver

246

films prepared at room temperature are found to be SERS inactive. DataDbtained using an ellipsometric technique show that the optical proper-ties of the two types of silver film are markedly different. The filmdeposited at low temperature is strongly absorbing, compared to the roomtemperature film, in the 320-600nm region. Calculated spectra show (Fig.l)that the observed difference may be accounted for in terms of the differ-ence in optical properties between a film with a smooth surface and afilm with a rough surface. As illustrated in the inset of Fig. 2 thesurface roughness is simulated by a layer of metal spheres resting on aflat surface. £ is the effective dielectric constant of the layer. Thedashed line spectrum is very similar to measured B spectra for vapordeposition of a rough silver film on top of a smooth film. The (0) - andsolid - line spectra are similar to B spectra measured for partiallyannealed rough silver films. Thus it appears that the SERS effect isrelated to the optical properties of the surface roughness layer. Such aroughness layer may be considered as a two dimensional aggregate of metalârticles. With this in mind we wished to see if SERS could be appliedto a thre^ dimensional aggregate of metal particles such as a dispersedsilver catalyst. Supported silver particles are unique in their abilityto catalyze the reaction of ethylene with oxygen to yield ethylene oxide.

Fig. 1.Calculated spectra forthe difference in opticalabsorbance between a silverfilm with a smooth surfaceand silver films withprogressively (solid, circleand dashed line) rough sur-faces. The quantitity B isrelated to the opticalabsorbance and is defined inreference 2.

EXPERIMENTAL

The experiments were performed using two types of supported silver samp-les which differed both in the nature of the support and in the manner ofpreparation. Type A. samples were prepared by a novel technique whichenabled the electrodeposition of metals into the porous surface ofanodized aluminium foil. Type B catalysts were prepared by filtering sil-ver colloid solution through glass fibre paper. The silver impregnated

247

100 cm-1 3000Fig. 2. SERS spectra of silver particles, of less than 2 urn diameter,

on a glass fibre support. The solid line was obtained at 300Kand displays peaks associated with hydrogenated surface carbon.At 460K [dashed line) the carbon related peaks are severelyattenuated.

glass fibre paper was taken as a model Ag/SiC^ catalyst.


In terms of SERS per-se the striking result of these experiments is thatunder appropriate condition the technique may be used to study silversamples at high temperatures. As shown in Fig. 2, SERS spectra were ob-tained for a sample held at 460K. In fact, type B samples which werereduced in H at 775K subsequently displayed surface Raman enhancement.Similarly, a type A sample which was heated to 770K in oxygen did notlose its SERS activity. In general SERS is not observed for silver samplesat or above room temperature due to the loss of the requisite morphologythrough sintering. However, for type A and type B samples the supportserves to minimize sintering. SEM pictures of a type B samples clearlyshow silver particles dispersed on glass fibres. The ability of SERS toprobe such samples under conditions of high pressure and high temperaturemakes it an important tool for studying catalysis. For instance, thecommercial production of ethylene oxide is carried out in the 520-600Krange using silver particles supported on low surface area refractorymaterials, and should therefore be amenable to SERS studies.

We will now concentrate on the behaviour of carbon on the silver particlesas revealed by SERS spectra. The observed intense carbon bands

248

at ^ 1370 cm" and ^ 1590 cm" (Fig- 2) are commonly observed in SERSstudies of silver surfaces. Using SERS and EELS Tsang ct. al. have iden-tified the features as belonging to amorphous (a) carbon. The feature at-1370 cm" correlates with the Raman spectrum of crystalline graphite,and the feature at 1590 cm is observed in the IR spectrum graphite.Our SERS spectra indicate that the surface carbon is to some extent hydro-genated, as a broad CH stretching band is observed at ^ 2800 cm

All of the freshly reduced samples displayed the intense a carbon Ramanbands. In the 450-550K range these bands disappeared even for a sampleheld under 1 atm. of C_H However, on cooling the sample to 450K thefull intensity of the Raman bands was restored. It was noted that heattreatment, to 750K in the presence of l.atm. of C2H4, caused a markedsharpening of the feature at ^ 1590 cm" . This change may be due toordering of the surface carbon as a result of increased mobility at hightemperature. Such an interpretation is consistent with the dependence ofthe Raman spectrum of carbon on long range order .

Considering the surface science results obtained by a number cf groupsit appears that the observed disappearance of the carbon Raman peaks inthe 450-550K region is attributable to a change in the nature of thesurface carbon. Madix has shown that the interaction of acetylene withoxygen adsorbed on Ag (110) leads to the formation of adsorbed C H at170K. The latter species was then observed to transform to C~ anaamorphous C at 275 and 550K, respectively; a LEED pattern associatedwith adsorbed C? was found to vanish at 550K. In light of these resultsby Madix we attribute the disappearance of the a carbon Raman peaks at ^500K to decomposition to adsorbed atomic carbon. Thus the results for thesupported silver sample are in good qualitative agreement with UHV studiesof a well defined Ag (110) sample.

Before concluding we note that the observed SERS signals in the presentstudy arise from adsorbates on the silver particles. This is confirmed bycomparing the Raman spectra of blank and silver impregnated samples. Inconclusion we have found that SERS may be used to study silver catalystsunder realistic conditions. Further, the results that we obtained for thebehaviour of carbon on supported silver particles are in good agreementwith results for carbon on a silver single crystal surface.

REFERENCES

1. A recent review of SERS is given by R. Miles, Surf. Int. Anal. 5(1983).43.

2. P.H. McBreen and M. Moskovits, J. Appl. Phys.,54 (1983) 329.3. D. Hall and M. Moskovits, presented at the 7 Canadian Symposium on

Catalysts, Edmonton, October 1980.4. J.C. Tsand, J.E. Demuth, P.N. Sanda and J.R. Kirtley, Chem. Phys.

Lett. 76 (1980) 54.5. P.C^ Painter, 0. Mahajan and P.L. Walker, Jr., Extended Abstracts,

14 Biennal Conf. on Carbon, p. 113 The Pennsylvania State Universi-ty (1979) .

6. R.J. Madix, Appl. Surf. Sci. 14 (1982-83) 41.

249

XPS STUDIES OF Si FILMS DEPOSITED FROM SiCl. BY AN RFCOLD PLASMA TECHNIQUE

E. Grossman, A. Grill and M. Polak

Ben-Gurion University of the Negev, Beer-Sheva, Israel

ABSTRACT

Microcrystalline silcon films were deposited in an inductively coupledglow discharge in SiCl.. XPS measurements show that the film has oxidelayers composed of SiOx with 05x$2, 12 A wide. The chlorine concentrationin the interface region is higher (5%) than that in the bulk (2%) , andit acts as a p-type dopant formiiv; Si~Cl bonds.

INTRODUCTION

In the past few years microcrystalline (me-)Si received increasing atten-tion due to its attractive properties (1,2). The relatively high conduct-ivity of the mc-Si can be very useful in amorphous (a-) Si solar cells asa conductive surface layer. a-Si solar ce'.ls are usually prepared fromSiH4 because of the essential role of the hydrogen as a dangling bondcompensator (3). However the use of SiCl- as a starting material, canlead to lower cost a-Si:H,Cl solar cells, with chlorine in the filmacting as a dangling bond terminator (4-6) . The thermal stability of thechlorinated film is better than that of the hydrogenated a-Si, sinceSi-Cl bonds (91 kcal/mole) are stronger than Si:H bonds (77.7 kcal/mole)(6). Several works have been performed on a-Si:H,CI (5-9) with emphasison the physical, electrical and optical properties of the films. In thepresent work we studied the surface and bulk chemical properties of Sifilms deposited from SiCl. in a rf plasma, by X-ray PhotoelectronSpectroscopy (XPS). Special attention was given to the Si-SiO- interfacedue to its importance to the microelectronic industry (10) and to theeffect of the CI in the film. All samples except one were exposed toair for several weeks.

EXPERIMENTAL

The Si coatings were deposited on a stainless steel substrates from gasmixtures of Ar.H^, and SiCl-, and B-H, or PH,, for dopant purposes. Theplasma was initiated in a quartz reactor inductively coupled to a PlasmaTherm rf generator of 27.12 MHz. The silicon films were obtained by thedecomposition of the SiCl- molecules in the plasma and at the stainlesssteel substrate, and their reduction by hydrogen. XPS measurements weremade with PHI model 549 system with double-pass cylindrical-mirroranalyzer operating in the retarding-potential mode. The typical operatingpressure in this system was 2x10" torr. Electrons were excited with aMg k X-ray source operating at an accelarating potential of 10 kV and apower of 400 W. Arjon ion beam sputtering with 0.5 keV ions was used for

250

sputter profiling of the films. The microstructure of the film was deter-

mined by Transmission Electron Microscopy (TEM) (Jeol JEM-200B).


The Transmission Electron Microscopy (TEM) studies indicated that the

films were microcrystalline. Figure 1 shows a TEM micrograph of a Si Film

deposited on a TEM grid. Figure la shows a bright field image of the film.

A dark field image which was taken in the (111) debye ring is shown in

figure lb. The bright areas have a crystalline structure while the dark

areas are amorphous. Figure 2 shows a diffraction pattern and schematic

diagram of the crystallographic planes of the Si film. The average size

of the crystallites is 150 A. About the same crystallite size was

determined by x-ray diffraction (9) . The following results were obtained

for samples exposed to air for several weeks. Figure 3 shows representa-

tive XP spectra of Si 2p, before sputtering (a), and after several

sputtering times. Spectrum (a)consists of two components, one with

binding energy of 99.1 eV, which is due to unoxidized Si, and the other

one at 102.7 eV, which is attributed to SiO,, (11). Sputtering reduces

the oxide component. Figure 4 shows the difference between the spectra

of the external layers (fig. 3 aid), and the spectrum from the bulk

(fig. 3e). In the first difference spectrum (a-e) there is a shift of 3.6

eV between the binding energies corresponding to SiO- and the unoxidized

Si. The peak is asymetric indicating the existence of other components

at inteimediate binding energies. In (b-e) the main component is shifted

3.1 eV, and is attributed to Si-O, (12) Contribution of SiO can be seen

in (c-e) with a shift of 1.9 eV (13). Accordingly the oxidation state of

Si decreases *rith depth, the topmost layer being SiO,,. The overall thick-

ness of the oxide layers was determined by using the expression for the

ratio of the peak intensities of oxidized silicon and of the underlaying

= 0.67

ISiC

r

expf __JL Uox CosS

Hwhere: d - oxide layer thickness, X - electron mean free path, β -effec-

tive angle between normal to the surface and the electron take-off

directions. In this case X was taken as 30 A aiiJ cos^=0.70. The experi-

mental value gave d=12 A. Figure 5 shows the concentrations ratios of

the oxygen and the oxidyzed silicon, as a function of the sputtering time,

for two samples. The first was exposed to air for several weeks, while

the other was exposed for 1 hour. For long exposure, PS seen above,

there was a gradual change of Si oxidation state, starting with SiO . On

the other hand, for short exposure to air, only a thin layer of SiO was

formed. Figure 6 shows the depth profile of the component concentrations

from the surface down to the substrate. At the top layer there is a sharp

decrease in oxygen due to the removal of oxide layers from the surface

region. The chlorine concentration is about 2% and it remains the same

upto the interface with the substrate. At the interface there is an

increase of the chlorine concentration to about 5%. This increase is

probably due to interaction of the stainless steel substrate with the

251

chlorine ions in the plasma, at the beginning of the process. Once a thinlayer of Si is formed, the substrate is no more exposed to the CI ions,and there is a reduction in the chlorine concentration to a constant value.In the interface region there is also an increase in the oxygen concentra-tion, probably due to oxygen which is originaly bonded to the surface ofthe stainless steel substrate. From the chemical shift of the Si peak,measured at the interface res ion, we found that the Si is bonded to theoxygen, forming SiO- layer. The binding energy of the Si 2p in the bulkis 99.1 eV. This value is 0.5 eV less than the binding energy for a-Si:H.This shift should be due to the chlorine present in the film. Because ofthe relatively low concentration of CI (~2%), no chemically-induced shiftis expected to be observed for the Si XPS lines. "Tierefore the apparent0.5 eV shift should be due to displacement of the Fermi level in the gap.Thus, the chlorine presence causes a downward shift of the Fermi level,so that undoped a-Si:H.Cl is behaving like p-type semiconductor (14).Inspection of the chemical-shift observed for Cl(2p) (-2.5eV) indicatesSi-Cl partially-ionic (0.5) covalent bonding.

REFERENCES

1. K.E. Spear, G. Willeke and P.G. LeComber, Physica, 117B $ 118B (1983)908.

2. S. Hasesawa, S. Narikawa and Y. Kurata, Physica, 117B 5 118B (1983)914.

3. M.H. Brodsky, M. Cardona and J.J. Cuomo, Phys. Rev. B16 (1476) 3556'.4. J. Chevallier, S. Kalem, S. al Dallal and J. Bourneix, J. Non-Cryst.

Solids, 51 (1982) 277.5. V. Augelli, R. Murri, S. Galassini and A. Tepore, Thin Solid Films,

69 (1980) 315.6. S. Kalem, J. Chevallier, S. al Dallal and J. Bourneix, J. Appl. Phys.

Coll. C4 suppl. No. 10. 42 (1981) 361.7. G. Bruno, P. Cpezzuto and F. Cramarossa, Thin Solid Films, 106 (1983) .8. R.D. Plattner, W.W. Kruhler, B. Rauscher, W. Stetter, J.G. Grabmaier,

Proc. 2nd E.C. Photovoltaic Solar Energy Conference, Berlin (1978)860.

9. E. Grossman, A. Grill and R. Avni, Plasma Chem. and Plasma Proc,2(4) (1982) 341.

10. W.Y. Chins, Physical Rev. B, 26(12) (1982) 6633.11. B. Carriere, J.P. Derielle, Ch. Burggrat, Analysis, 9(5) (1981) 236.12. CM. Garner. I. Lindau, C.Y. Su.P. Pianetta and W.E. Spicer, Physical

Rev. D. 19(19) (1979) 3944.13. T. Adachi and C.R. Helms, J. Electrochem. Soc, 127(7) (1980) 1617.14. A.E. Delahoy and R.W. Griffith, J. Appl. Phys., 52(10) (1981) 6337.

252

Figure 1TEM micrograph of plasma depositedmicrocrystalline Si film Ca) brightfield image (b) dark field image,taken in the (111) debye ring.

Figure 2Selected area diffraction patternand schematic diagram, at themicrocrystalline Si film shownin figure 1.

Figure 3Xp spectra of Si 2p, before (a),and after spattering (bfe) .

Figure 4Difference spectra curves of theexternal layers (fig- 3 afd), aftersubtracting the spectrum at thebulk Cfig. 3e) .

4

j

long

short

-

exposure

exposure

— 1 I—

SiO2

Si 2 O 3

SiO

10 20 30

SPUTTER TIME,min

§70

I GO

S4O

§30

20

10

0

• -Si•-a•-o

4 \

•»-t--"-T"?-<"

Figure 5Oxygen/oxidyzed silicon concentrationsas a function of sputtering time forthe two atmosphereic exposures

O 3D 30 40 50 60 70 30 330 350 370

SPUTTER TIME.mn

Figure 6Depth profile of the long-exposed Si film

253

QUANTITATIVE XPS OF HaS-ALUMINA:EVIDENCE FOR SODIUM ANISOTROPIC SEGREGATION

Y. Grinbaum and M. Polak

Department of M a t e r i a l s E n g i n e e r i n g ,Ben-Gurion U n i v e r s i t y of t h e Negev,

3eer-Sheva, Israel

ABSTRACT

The surface composition of two faces of the superionic conduc-tor sodium β-alumina has been studied using quantitative X-rayPhotoelectron-Spectroscopy (XPS). The most remarkable resultconcerns the occurance of sodium enrichment at one face asreflected by the relative line intensities of Na and Al fromthe two surfaces. This conclusion is supported by quantitativecalculation of relative intensities expected for crystal facesas perfect bulk planes, and by detailed in-depth compositionprofiles. Possible mechanism for the !Ia anisotropic segregationprocess in the air-exposed crystal and its implication forthe performance of Ha-B alumina as a solid electrolyte, aredi scussod.

Surface and interface phenomena in ionic solids have beenstudied to rather limited extent, as compared to metallicmaterials, in spite of their implications for variousproperties of such materials. A fundamental question concernsthe surface composition and how closely it mirrors the bulkcomposition. Deviations from the bulk compositions can bedue to extrinsic factors, such as interactions with gasesand influence of applied electric fields, or to intrinsicforces causing elemental segregation to the free surface,interface or grain boundries.

cleovoge face

Fig. 1. Schematic diagram ofsingle-crystal !Ja6-alumina;sodium (•) and oxygen (o) ionsconstitute the exposed conduc-tion plane at the cleavage-fact. The spinel blocks separa-ting the conduction planes fromeach other consist of close-packed aluminum and oxygen ions.

/"O'O"O"O'O"O'cy/O'O'C'G-O'O'O/'o'c'c'o'o'o'o/

edge foce

We have been conducting a comprehensive study of surfacephenomena of the super ionic conductor sodium β-alumina

254

(1.2 II a O'llAl 0 ) using elcctronapec tro scopic techniques(XPS, AES). Due to its exceptionally high ionic conductivityat room-temperature, its high mechanical strength and chemicalstabilityj Wa-! alutnitiii has been used as solid-electrolyte inthe high energy dc-i.d i ly sodium-sulfur battery » arid rianyinvestigators tried to <- J u..-i.i •. -• the Ha fast diffusionmechanism fl,2]. Previously v/e reported on the unuaual res-ponse ofNa-6 alumina, crystal to electron bombardment asrevealed by AES experiments [3-5] • When the electron-beamwas directed onto the cleavage face (see Fig. 1 ) , induced highlevels of negative charging at the surface region causedemergence of some sodium through induced cracks and faultsin the spinel blocks. On the other hand, electron bombardmentof the perpendialar edge face (Fig. l) resulted in weaknegative charging which rather easily forced out large quanti-ties of sodium. Thus, in this case electroir.igrat ion of Ha"along the conduction planes to the external edge-face tookplace and could be followed quantitatively via changes inthe Auger intensities to give the bulk diffusion constant[3].tfhile these experiments revealed t?ie electron-beam inducedanisotropic response of the Ila-g ?i i u m: i, a 'Tvstal, M s .-._,_-; u I <• dwith its two-dimensional superionic conauetivity, the presentstudy aimed at exploring possible anisotrophy in the unpertur-bed surface composition of air-cleaved TJa-0 alumina. There-fore, rather than AES with the electron beam strongly affec-ting the mobile sodiuia LO.-JLS in this case comparative XPSmeasurements of the cleavage ;:.::,t e-.ip-j ?a.-.::-j, wcr >_• ... •._•!'.. fined .Observed differences in relative line intensities between thetwo surfaces, had to be compared to calculated values for thecase of surfaces as terminating bulk planes.

EXPERIMENTAL

Air-cleaved slab of single crystal HSL-Q xli..mina. v •. I h ?~S rup.iwide edge-face was chosen, so as to yield reasonable XPSintensities. Yet, signal averaging was needed for the highresolution spectra of both surfaces (Pass-energy 50). Thecrystal was mounted on the sample holder in such a way thateither the cleavage face or the edge-face could be analyzedafter careful alignment in front of the CMA ( PHI 549 system).Mild sputtering was achieved with 500 V Ar+ ions (>ixlOtorr) operated for short time intervals in order to obtainthe near surface compositions.

RESULTS A1ID DISCUSSION

Figs. 2a-b present XP survey spectra of the cleavage and edgefaces. As can be seen, in both cases there is a significantcarbon contamination and different relative intensities ofsodium, aluminum und ox;/een- ^n particular, ' -e .-.7- '.-tri;p of theas-inserted edge fi-cir- (Fig. 2b) roi'.lcjts rtlal:--?jj smallconcentration of aluminum as compared to TTa, O and C. Focussing on the aluminum-sodium ratio, we have chosen for quantita-tive analysis to compare the aluminum intensity with that of

255

F i g . 2. Survey XP s p e c t r a of a s -i n s e r t e d and A r + - s p u t t e r e d s i n g l ec r y s t a l Na 3- a lumina. (Mg Kα IJOCWX-rays; CMA pass-energy 100 V o l t ) .

ttuttceo eogt lote

\

i-X S-M IOC 7ji fJO V * iTO *> fJC 0

BINDING tNtKGf. eV

the sodium Auger transition (KLL) rather than of Na (1$)since the former are much closer in electron kinetic energy,and! therefore are similarly attenuated by any overlayer conta-mination. Most significantly, the Na(KLL)/Al( 2P) peak arearatio measured from the high-resolution spectra of the edge-face is about six times larger than for the clearage-face.This remarkable difference may reflect excess sodium accumu-lation at the surface region of the edge-face,but possibleintensity variation associated with the greatly anisotropiccrystal structure of Na-0 alumina (Fig. l) should be consi-dered before any firm conclusion can be made. We have derivedexpressions for the relative XPS intensities, ^ - / ^ i 'expected for the two faces under the assumptions: a J thecleavage plane as an overlayer contains, due to symmetry,half the number of sodium ions in a bulk conduction plane,and b) the edge-face region has homogeneous and stoichio-metric distribution of Na and Al. The ratio of relativeintensities, (iIJa/IAl) edge/( INa/IAl) cleavage, then reads

Cal = 2"Na ^NaSin9exp( -

Na 2[l-exp( -d1 /> EinQ)]

-\-1

with n and n'jj

a as the numbers of sodium atoms per unit vo-

lume ana per unit area in a bulk conduction-plane, respecti-vely;d is the effective "thickness" of a conduction plane,d' is the distance between adjacent conduction planes (11.3 S

'Na a n d - the mean-free-paths of the

256

corresponding electrons(taken from data for alumina matrix [6])and 6 is the "effective" emission angle of the elcetronscollected by the CMA [7]. With IK, /n' calcualted from thecrystal structure, R 1 for Na( KLL ) / All 2P ) is 0 . U1* , whereasthe measured peak area pave the experimental ratio, Rejjp

= 6±1This significant distinction between the measured ratio andthe ratio calculated for "ideal" surfaces as terminating bulkplanes seems to establish the conclusion concerning Ka segre-gation at the edge-face of Na-B alumina.

Also, the measured relative intensities, 0(lS)/Al(2P) andC(IS)/Al(2P), indicate excess of oxygen and even more of carbonat the edge-face as compared to the cleavage-face composition.In depth composition profiles of the edge-face obtained withshort-time Ar+ sputter-etching are shown in Fig. 3. During thefirst 15 seconds of sputtering mainly carbon was removed thusleading to enhancement of signals from all other elements,'•fhile further sputtering gradually exposed more aluminum andoxygen, the sodium profile reached a maximum and then decrea<-sed similarly to the carbon. Disregarding possible preferentialsputtering effects, this behaviour is consistent with the edgeI!a segregation picture indicated by the quantitative analysisof the unsputtered surfaces. Inspecting the variations in 'relative intensities with sputtering-time (Fig. h) adds furtherinformation. Thus, during the first 30 sec there is rathersharp decrease in the Ila/Al ratio, and also a clear increasein the Al/ 0 ratio, which may indicate that the as-insertededge surface contained some excess oxygen (adsorbed or oxidic)as compared to the bulk of the crystal. This increase can notbe due to the different attenuations of the 0( IS) and Al(2P)signals coiring from the bulk by outer overlayers, since 0( IS)ha;.; shallower sampling depth so that Al(2P^)(lS) should havedecreased upon removal of overlayers with no oxypen. Theexperimental ratio (edge/cleavage) of Na/Al relative intensi-ties, R , decreased significantly upon 13 minutes sputteringof the ecFge-face (see Figs 2b and 2c) from a value of 6 to0.7, which is much closer to theoretical value derived above(RCal = o.U).

Fig- 3. High-resolutionXPS peak area vs. sputteringtime for ;iu-3 alumina edge-face. (500 volt Ar +).

O(IS)

Na(KLl)

2 3 4 t> 6SPUTTERING TIMF <m.n)

25 7

Fig. h. Ratios ofpeak area from Fig. 3.

-

UJ

sUJ(r

t t

. • • • * • ; •

1

__i i i i

ftil2P)/01IS) • • *

*

No(KLL)/CllSI

NotKLL)/flll2P)

«

1 , 1 , 1 , 1 , 1 ^ 1

0 1 2 1 1 5 6 7 B

SPUTTERING TIME ( m m )

•fhile the quantitative investigation shows that the anisotro-pic segregation of sodium at the edge surface is accompaniedby accumulation of carbon and some oxygen, no definiteconclusion regarding the chemical bonding of carbon andoxygen could be drawn from their XPS chemical-shift data (thechemical shifts of the sodium lines are characteristic ofHa+). Nevertheless, we can guess that the preferentialaccumulation of Na+ at the edge-face originates from inter-action (or chemical reaction) between the mobile Na+

and atmospheric gases (e.g. H ?0). Such an interaction mayoccur at the cleavage face as well, but due to the particularcrystal structure, only at the edge-face excess sodiumdiffusing from the near-surface region of conduction planescan be expected to emerge (see arrows in Fig. l).

Suci. anisotropic segregation of sodium may affect the per-formance of polycrystalline sintered Na-f5 alumina as solid-electrolyte in the Wa/S battery. Thus, one can expect thatmass and charge transport through the solid-electrolyte/liquid-electrode interface would be affected by the particu-lar composition of the solid-electrolyte surface.

REFERENCES

1. G.D. Mahan and W.L. Roth Eds., "Superionic Conductors"(Plenum, Nev York, 1976).

2. A. Highe, M. Polak and R.'.{ . Vaughan, in "Proc. Inter.Conf. on Fast Ion Transport in Solids", Wisconsin, 1979(North-Holland, Amsterdam, 1979).

3. A. Livshits and M. Polak, Surface Science,119 (1982) 311*.•» . M. Polak and A. Livshitz. Appl. Surface Science, 10

(1982) kk6.5- A. Livshits and M. Polak, Vaccuum 33 (1983) 2Ul.6. F.L. Battye et. al. Phys. Rev. D. 9 (l97l») 2887.7. A. Livshits and M. Polak, to be published.

258

STRUCTURE OF PROTECTIVE DIFFUSION COATING FOR NIOBIUM ALLOYS

1 2 1 2M. Kazinets , 0. Gafri , L. Zevin , B. Rabin

Institute for Chemistry and Chemical Technology, The Institutes forApplied Research, Ben-Gurion University, Beer-Sheva, Israel

IAI Bedek Aviation, Israel.

INTRODUCTION

Nb alloys, like other refractories alloys, have good mechanical proper-ties at high temperatures but low oxidation resistance at these tempera-tures. Silicide-based, multilayered coatings are used for protection ofNb alloys against oxidation. One kind of coating (R512E) for the niobiumalloy D-43 was found (1) to be composed of six discrete layers, some ofwhich appeared to be single-phased. Five of the layers contained NbSi2or Me5Si3 (Me = Cr, Fe) with slight amounts of NbSi2, and only the sixthwas Nb5Si3. In our preliminary studies such coatings showed differencesin the arrangement of the layers and in their phase compositions. Theresults of further work are presented here.

EXPERIMENT AND RESULTS

The specimens were samples of Nb alloy C-103, with a silicide coating(Si-20Fe-20Cr) formed by the slurry-diffusion process (2) . The metal-lographic structure of the coatings was studied with the scanning electronmicroscope (SEM). A cross section of one sample (Fig. la) shows thelayered structure of the coating. A diagram of these layers is shown inFig. lb. The total thickness of the coating was about 200 mm. The sur-face layer A of the coating is very thin, 10-15 pm. Layer B, which isjust under the surface, is two-phased (I+II) , with the regions adjoiningthe surface richer in phase I. Layer C has approximately equal amountsof both phases. The thickness of layers B+C is about 120 ym. The nextlayer D is about 40 ym thick and seems monophasic. The boundary betweenlayers C and D is sharp. Layer E is about 30 pm thick, is adjacent tothe bulk Nb, and consists of 4-5 thin diffusion layers. The compositionof each layer was determined by EDA X-ray spectrometry (Table I).

Note that phase II contains appreciable amounts of Cr, Fe and Hf, whilephase I contains low concentrations of these elements. The Cr distribu-tion in a cross-section of the coating (Fig. lc) provides more evidencethat phase II contains large amounts of Cr.

The phase composition of each layer was determined by X-ray diffractionfrom the surfaces exposed by consecutive removal of layers. The pene-tration of the CuKa radiation used here is such that the thickness ofthe layer responsible for 90% of the diffraction intensity does not ex-ceed 0.6 vm. This is less than the thickness of the removed layers. Atotal of 30 layers 5-6 urn each were mechanically removed, and 30 X-raydiffraction analyses were done.

259

a) b) c)

Fig. 1. Cross section of coating: a) SEM pattern, b) Diagram,

c) Cr distribution

Table 1. Weight ratios of coirponents in the layers of the coating

Layer

A

B

C

D

E

Phase orsublayer

I

I I

I

II

4

3

2

1

S i

38.8

44.4

33.4

43.1

31.3

30.9

40.1

38.9

28.6

Element, %

Nb

31.1

54.2

30.5

52.5

29.6

37.7

51.1

54.3

61.3

80-90

T i

0 . 6

0 . 2

0 . 6

0.2

0.6

0 . 6

0 . 4

0 . 4

0 . 7

C r

14.4

1 . 1

19.2

2.6

20.0

9 . 1

1 . 4

0 . 8

0 . 9

Fe

11.1

0 . 1

12.5

1.3

14.7

17.7

4 . 2

1.9

3 . 0

Hf

3 .

-

3 .

0.

3.

4 .

2 .

3 .

5 .

9

8

3

8

0

8

,7

,5

Nb/Si

0.80

1.22

0.91

1.22

0.94

1.22

1.27

1.40

2.14

260

From the chemical data we could surmise that the layers are composed ofNb-Si compounds with some Cr (Fe)-Si compounds in regions rich in theseelements. The phase identification was done on the basis of the X-raydiffraction data (3). There are six compounds in the Nb-Si system. They

a r e :

NbSi2, with a hexagonal lattice, a = 4.797A, c = 6.592A, space groupP6222;Nb5Si3 - tetragonal, a = 6.570A; c = 11.884A;Nb5Si3 - tetragonal, high temperature phase, a = 10.018S, c = 5.072A,space group 142 m;Nb5Si3 - hexagonal, a = 7.5368, c = 5.248S;ttt>3Si - cubic, a = 4.211A; andNb3Si - tetragonal, a = 10.2l8, c = 5.19?.In the Cr-Si system the compound Cr Si2 with a hexagonal lattice, a =4.42R and c = 6.35A is known.Some diffraction lines were identified on the basis of the interplanarspacings of NbSi2- It was difficult to identify the rest of the linesbecause none of the known Nb-Si compounds has lines that coincide satis-factorily with the deXp. We therefore indexed these lines on the basisof the high-temperature tetragonal modification of Nb5Si3 (a = 10.018Aand c = 5.072A), with the assumption that the Bravais lattice of thiscompound is not body-centered (4) but primitive. This permits the appea-rance of the lines with uneven sums of Miller indices, such as 102(d = 2.46) and 331 ( d = 2.14). Thus, there are two phases in the sur-face layer, MbSi2 and the high-temperature tetragonal phase of

The X-ray pattern changes at once after the removal of the surface layerA. The intensity of the l 35813 peaks increases. No further changes occur-red in the next five patterns. During this time about 30 ym was removed,which represents the full thickness of layer B. In the sequence of dif-fraction patterns from layer C the intensities of the NbSi2 lines in-creased, while those of NbsSi3 decreased. The surface of this sectionconsisted of two phases with a prevalence of the NbSi2 phase.The intensities of only two pairs of lines belonging to Nb5Si3 begin toincrease again in layer D, those with d = 2O49 (400 ) +2.46 (102) and1.24 (800) + 1.23 (204). The second pair of lines is the second-orderreflections of the first pair. All other lines of Nb5Si3 were relative-ly weak. Therefore, this layer is textured. The NbSi2 peaks decreaseand disappear. The X-ray pattern obtained in the middle of layer D showedonly two peaks, 400 and 800; thus layer D is single-phased, homogeneous,and textured. The diffraction pattern from layer E contains peaks ofboth phases, NbSi2 and Nb5Si3, and also lines belonging to metallic Nb.

DISCUSSION

A comparison of the results obtained from metallographic investigationwith those of X-ray diffraction shows that phase I is Nb5Sx3 and thephase II is NbSi2. Cr, Fe and Hf occur together with NbSi2 (II) and areabsent from the areas of NbqSi3 (I) . The Cr in phase II may substitutefor Nb in NbSi2, because NbSi2 and CrSi2 are isostructural. The hexago-nal phase of compound Nb5Si3 does not exist at all in the coatings. Thisis a very important result, because the hexagonal phase NbsSi3 has a lowoxidation resistance (5), in contrast to MbSi2 and the high-temperaturetetragonal phase of Nb5Si3, which have a high oxidation resistance.

261

The results of the compositional analyses (Table I), show that there isan excess of Si in the coating layers, but we have not yet made an assign-ment for this surplus Si.

There are some differences between the structure of the layer analyzedearlier (1) and that determined here. The most important of them is thatCr and Fe were previously found to be absent from the NbSi2 layers (1) .However, they coexist in significant quantities with NbSi2 in our investi-gation.

REFERENCES:

1. Priceman S. and Sama L. Electrochem. Tech. 6_ 315-326, 1968.

2. Gafri, I., Rabin, B., Materials Engineering Conference 1981, Dec.20-22 Technion, Haifa, Israel.

3. Joint Committee on powder diffraction standards. Powder diffractionfile.

4. Parthe, L. Monatsch Chem. 86_, 385, 1955.

5. Lyndon B. Johnson Space Center, Houston, Texas NASA Tech. Briefs,Spring/Summer, p. 395, 1982.

262

THE EFFECT OF B AND P DOPING ON THE DEPOSITION RATE AND CHLORINE

CONCENTRATION OF Si FROM SiCl4 IN RF PLASMA

R. Manory, E. Grossman, R. Avni and A. Grill

Materials Engineering Department, Ben-Gurion University of the Negev

Beer-Sheva, Israel

ABSTRACT

Silicon films were deposited by a cold RF plasma from a mixture of

SiCl4 + H2 + AΓ and were doped in situ by introducing diborane or phosphine

in the gas feed. The film thickness was measured by Scanning Electron

Microscope and the content was determined by Energy Dispersive X-ray

Analysis.

The deposition rate and the chlorine content were reduced with

addition of diborane and both were enhanced with addition of phosphine to

the gas mixture.

The described phenomena have been found to be independent of substrate

material or deposition period. The morphology of the deposit has been

found to be different in the two types of films. B-doped films show an

uniform growth proceeding by uniform coverage of the substrate, while

P-doped films show an erratic growth proceeding by agglomeration of solid

clusters at preferential sites. The different deposition rates are

explained by differences in the surface morphology.

INTRODUCTION

In a series of papers published almost simultaneously ten years ago

(1-5) several workers reported the inverse effect which dopants from the

3A and 5A groups have on the deposition of silicon in CVD and plasma

reactors. All these reports showed similar findings: that the deposition

rates of Si are diminished in the presence of doping gases containing As

or P(AsH3 or PH3) and increase in the presence of boron containing gases,

such as B2H6 or BCI3 . Eversteyn and Put (1) performed CVD of silicon

using silane as the silicon gas, while Ray-Choudhury and Hower (2) used

both silane and silicon tetrachloride (SiCl4) in CVD and obtained similar

results. Hall and Koliwad (3) showed that boron lowered the temperature

required for chemical vapor deposition of silicon from silane by more than

two hundred degrees C and increased the deposition rate probably due to a

catalytic effect on the decomposition of silane. Similar findings were

reported by Yasuda et al. (4), while Farrow obtained the same behavior of

the deposition rates in silane plasma (5).

Recently, Knights (6) reported on the catalytic effect of diborane

during plasma deposition of a:Si-H from a gas mixture of SiH4 + Ar.Several

attempts were made to explain this behavior. Chuang [7) has presented

one of the more accepted theories about this behavior. He explained the

inverse effects of phosphine and boron by the different electronic

configuration of the depositing molecules and the surface. According to

[7), both SiH4 and SiCl4 molecules have Si+-H" and Si

+-Cl" ionic bonds

respectively, would be attracted by a positive surface potential

(p-type Si) and repelled by a negative surface potential (n-type Si).

263

The growth rate of the layer is thus affected in the adsorption stage.These effects of increasing and decreasing the rate of silicon depositionby B and P respectively have become since then a well known and acceptedfact.

Therefore, details about deposition rates obtained during doping inthe plasma are now scarcely reported.

In our deposition system we employed RF plasma to deposit si l iconfrom a gas mixture of SiCl4 + Ar + H2, and B2H6 and PH3 were used asdopant gases. The deposition rates of silicon on graphite and glasswere invetigated as a function of the plasma macrovariables, togetherwith the CI content of the films and their moxphology.

In this paper we present the effect of doping gases, diborane orphosphine on the deposition rate and composition of Si films depositedby an rf plasma of Ar + H2 + SiCl4. The difference between our resultsand the quoted data (1-7) will be discussed.

EXPERIMENTAL

The deposi t ion was ca r r i ed out in the experimental setup shown inFigure 1. This system has been previously descr ibed elsewhere ( 8 , 9 ) , anda few d e t a i l s are repeated here for reasons of c l a r i t y :

The b a s i c gaseous mixture consisted of SiCl4(99.9% pure suppl ied byMerck Gmbh."), hydrogen and argon (UHP grade-supplied by Matheson, USA) .High p u r i t y diborane and phosphine di luted in argon were d i r e c t l yintroduced i n t o the gas mixture and monitored by a mass flow c o n t r o l l e r .(Matheson-MFC). The films were grown at various conditions on graphi teor glass subs t ra tes using dopant to S i C ^ concentrat ions which var ied from

2-10-3 to 2 .10" 2 . The s i l i c o n to hydrogen r a t i o in the gas feed was 1:3.The pressure in the reac tor was 1 or 2 mbar, the input power was between100 and ISO Watts and the pos i t i on of the subs t r a t e in the r eac to r wasupstream, a t the beginning of the RF coi l (pos i t ion H - Ref. 10,11) . Theexperiments were carr ied out in p a i r s , using the same conditions fordeposition of p-type or n-type layers . Therefore, the external parametersof the plasma such as p res su re , subst ra te p o s i t i o n , input power andperiod of deposition affected bath types of layers in the same way, andthe observed differences are a t t r ibu ted only to the effects of the dopinggases.

Thickness measurements were made by Scanning Electron Microscope (SEM)and the chlor ine content of the film was measured by Energy DispersiveX-Ray Analysis (pDAX). For samples grwon on g l a s s , the percent of CI vs .Si in the film could not be determined, because the subs t ra te i t s e l fcontained s i l i c o n


A SEM micrograph of two layers deposited one upon the other i s shownin Figure 2A, while Figure 2B showed the l ine-scan p ro f i l e of the CIconcentration for the same sample. The upper l ayer was P-doped and thelower one was B-doped, and each layer was deposited during two hours at150 Watts input power and a pressure of 2 mbar. The dopant to SiCl4 r a t i oin the gas stream was 10~3 for both PH3 and B2H6 gases. The r a t i o betweenthe thickness of the B and P-doped layers in t h i s sample was measured tobe 1:2.8 respec t ive ly . This fac to r (approximately 3) was found a l so forthe r a t i o of the chlorine content in the two layers shown in Figure 2B.Similar values for the two r a t i o s were obtained a lso when the order of

264

doping was changed and the deposition time was reduced to 1 hour foreach layer. Table I summarizes the findings for several samples depositedon graphite and on glass substrates. The growth rate on the glasssubstrates was slower than on the graphite substrates. The filmsdeposited on glass were grown separately and the comparison was madebetween different samples. This comparison shows that the ratio ofdeposition rates between B and P doped samples remains 1:3 even thoughthe dopant concentration was changed by one order of magnitude. Thedifference in deposition rates between undoped and n-tyne layers depositedon glass was found to be 1.3 "" '„ value shows that phosphine enhancesthe deposition not only as _^mpared to the effect of diboTane, but alsoas compared to the growth of undoped layers.

The data presented in Table I also indicates that the effect of thedopant gases on the growth rate is not affected by the nature of thesubstrate.

These findings, first presented in (11) were later confirmed by othergroups (12,13) working with halogeneous silanes (SiF4 and SiClj.respectively) in plasma. The described behavior, which is inverse to thebehavior reported for CVD (1-5) or for plasma of SiH4(6) indicates thatthe effect of the doping gases is related to phenomena in the plasma phaseand to the plasma-solid interaction for plasmas of halogenated gases.

Figure 3 presents SEM micrographs of films deposited for short timeson glass substrates. One can observe the uniformity of the growth of theB-doped film, as compared to the erratic surface of the P-doped and theundoped layers. The latter presents a growth pattern based on pile-upsat preferential sites. These pile-ups serve as three dimensional growthcanters while the B-doped films grow uniformly on unoccupied sites on thesurface. The different morphology of the growing films can account forthe different deposition rates considering that pile-ups can grow intoall directions above the surface, while the B-doped film grows uniformlyon unoccupied sites on the surface. This morphological difference canexplain the almost constant ratio between the deposition rates with thetwo dopants, for dopant concentrations in the gas as low as 10~3f andother deposition parameters.

Figure 4A shows the behavior of the chlorine concentration vs.deposition time for undoped Si films, deposited on graphite substrates,and Figure 4B shows the behavior of the deposition rate for the samedeposition conditions. A detailed discussion of these figures may befound elsewhere (14). As i t can be observed the deposition rate and theCI concentration behave similarly with time, indicating a correlationbetween these two parameters. A similar correlation is observed alsoduring doping, as indicated by Table I.

The ratio between the chlorine concentrations in the two kinds offilms deposited under similar conditions, is the same as the ratio ofthe growth rate of these films. We conclude therefore that the changesin the CI content in the doped films are a direct result of the differentdeposition rates of these films. It should be stressed that the presenceof chlorine in the deposit is a significant difference between Si filmsobtained in plasma of SiCl4 as compared to Si films obtained by CVD fromSiCl4 or by plasma of SiHg. I t is therefore possible that the presenceof CI in the films causes the difference between the effect of the dopantson the growth rate as observed in plasma deposition and CVD. This effecthas s t i l l to be studied.

265

REFERENCES

CI) F.C. Eversteyn and B.H. Putt, J. Electrochem. Soc. 12(3, 107 (1973).(2) P. Ray-Chouahury and P.L. Hower, J. Electrochem. Soc. 120, 1761

(1973) .(3) L.H. Hall and K.M. Koliwad, J. Electrochem. Soc. 120_, 1438 (1973).(4) Y. Yasuda, K. Hirabayoshi and T. Moriya, Proc. 5th Conf. on Solid

State Devices, Tokyo (1973).(5) R.F.C. Farrow, J. Electrochem.Snc. 121_, 899 (1974).(6) J. Knights, J. Non-Cryst. Solids, 35_ and 36_, 159 (1980).(7) Chin-An Chuang, J. Electrochem. Soc., 123, 1245 (1976).(8) E. Grossman, M.Sc. Thesis, Ben-Gurion University, Beer-Sheva

(in Hebrew), (1981).(9) R. Avni, U. Carmi, R. Manory, E. Grossman and A. Grill, Proc. 6th

Int. Symp. on Plasma Chem., Montreal, M.I. Boulos and R.J. Munz Eds.,Vol. 3 p. 820 (1983).

(10) A. Grill, E. Grossman, R. Manory, U. Carmi and R. Avni, ibid, p. 843(1983).

(11) R. Manory, E. Grossman, R. Avni and A. Grill,Phys. S oc, Bulletin of IPS, _2£, p.32 (1983).

(12) B. Pratt, private communication (1983).(13) P. Capezutto, G. Bruno, F. Cramarossa, Proc. 6th Int. Symp. on

Plasma Chem., Montreal, M.I. Boulos and R.J. Munz, Eds., Vol.3,p. 815(1983).

(14) E. Grossman, A. Grill and R. Avni, Plasma Chem. Plasma Proc, 2,p. 341 (1982). ~

Ann. Meeting Israel

PRFPARATItW OF n- AND p-lYPE HICHOTKYSTALUNE Si BY THE COLD PLASMA TECHNIQUE

1. gases introducing system 2. mass flow controller 3 . mixing chamber

4. Teactor 5. sample holder 6. matching unit 7. Z7.12 MU RF generator

B. induction coil

D. selection valve H. heater M. pressure gauge P. pumping system

T. cold trap

Fig. 1: Experimental setup for plasma deposition.Fig. 2: A.

B

SEM micrograph of two Si layers on graphite Sample 219)Thick layer - n type Si, Thin layer •• p-type Si. 2hr deposi-tion timeEDAX Line scan profile of CI concentration vs. Si for sample219 (For CI concentration values see Table 1).

266

Figure 3: SEM micrographs of Si samples grown on glass:A - undoped B - n-type C - p-type.

Figure 4: A - Behavior of deposition rate vs. deposition time forundoped Si deposited on graphite.B - Rate of CI incorporation ( CC1) in the coating vs.deposition time.

3 4 5 6 7 8 9TIME, hours

3 4 5 6 7 6

TIME, hours

10

Table I. Deposition Rates and CI Concentration for Several Si Layers.'

SampleNo.

219

220

223

504

505

S06

DopingOrder

I-BII-P

I-BII-P

I-PII-B

none

P

Bi

Substrate

graphite

graphite

graphite

glass

glass

glass

AverageDep.Rate

A/sec

514.2

4.913.9

11.74.2

2.2

2.8

0.9

depositionRate Ratio

CB/P)

1:2.8

1:2.8

1:2.8

(504/505)1:1.3

C506/S05)1:3

CI vs. Si(%)

3.511

4.8IS

7.82.6

CIContentRatio

1:3.1

1:3.1

1:3.4

Dopant:SiCl4Cone.

ID"3

ID"3

ID"3

—

2-10"2

2-KT2

CDExperimental conditions: for graphite: 150 Watts input power,p=2 mbar, SiCl4 : 3.7 v/o; total flow 130 SCCH.for glass: 100 Watts input power, p=1 mbar, SiCl4 : 5 v/o;total flow 60 SCCM.

26:

SILICON NITRIDE COATINGS BY THE LOW PRESSURE RF PLASMA TECHNIQUE

U. Carmi, A. Raveh, A. Inspektor and R. Avni

Nuclear Research Center-Negev, POB 9001, Beer-Sueva, 84190 ISRAEL

ABSTRACT

Silicon nitride coatings have been deposited in a low pressure r.f.plasma on stainless martensitic steels (AISI-410), from gas mixtures ofsilicon tetrachloride (SiCl4) and amcnia (NH3) in an argon plasma. Thesubstrates temperature between 230-440°C, depending mainly on the inducedr.f. power. The coatings obtained were identified by x-ray diffractionand analyzed by scanning electron microscopy, optical microscopy andEDAX for morphology, thickness, and silicon and chlorine contents, respec-tively. The influence of parameters such as the pressure in the reactorand the r.f. power input on the properties of the coating such as thedeposition rate, morphology and microhardness were investigated.

The correlation between plasma parameters and properties of silicon ni-tride is shownSome discussion on the nature of the reactions mechanismis presented.

INTRODUCTION

Silicon nitride had been chosen as a fine ceramic coating material dueto its special properties. Silicon nitride is characterized by its highhardnessCll, its high resistance to thermal shock aid to chemical reac-tions. It has high fracture thoughness(2) and high Vickers microhard-ness (3) . These propertied make silicon nitride suitable for high tempe-rature wear resistant coatings on parts such as turbine blades.

Several methods are available for the deposition of silicon nitride coa-tings, each has its virtues as well as drawbacks. Amongst others one canname Chemical Vapour Deposition (CVD)(4), glow discharge(5), and sputte-ring (6) (reactive and non-reactive).

R.F. plasma coatings offers a technique whereby some of the disadvantagesof the above methods are avoided. In this technique the coating is car-ried out at low temperature (300-500°C) the substrate is coa": 1. on allsides including inside pores and holes, at relatively high depositionrate and adhesion.

The deposition and coating of silicon nitride by the r.f. plasma techni-que will be discussed with regards to working parameters and resultingproperties.

EXPERIMENTAL

A full detailed description of the experimental setup is given elsewherehowever for convinience of the readers some details are given below.R.F. plasma is induced inside a pyxex reactor by r.f. coil wound aroundit. Silicon nitride is produced on a metal substrate from silicon chlo-ride and amonia in argon atmosphere. The coated metal was examined forits microhardness chlorine content and deposition rate. Also the morpho-logy and crystaljographic structure were examined by SEM and XRD

268

respectively.

RESULTS AND DISCUSSIONS

Primary investigations of the silicon nitride deposition indicated that thesample to be coated should be placed in position H in the plasma (ups-tream) and that the sample should be electricaly floating. The position-ning of the sample was fixed so that the chlorine content in the samplewould be minimal (8). The electrical connection were determined from -.the comparison of deposition rate on substrates biased and floatingThis was corrob rated by the results obtained by EPR (Electron Paramagne-tic Resonance) sampling of the plasma, showing that the (SiC14+NH3)/Arplasma had a large concentration of radicals(10).XRD identification ofthe coatings shows the existance of hexagonal silicon nitride with cellunits a=7.76^ and C=5.62A". Microhardness, deposition rate and chlorinecontent were determined as a function of working parameters i.e. pressure,power and NH3/SiC14 ratio. Figures 1-3 represent the changes in the aboveproperties varying the pressure. It is noted (fig.l) that the microhard-ness increases with pressure up to 3 Torr, where it obtains its maximumvalue of 2400kg/mm2. At higher pressures the microhardness decreases.k similar behaviour is obtained for the deposition rate (fig.2) and forthe chlorine content (fig.3) . It should be noted here that chlorinecontent is in percentage of silicon only as measured by EDAX. All theabove results show that at 3 Torr the reaction reaches it optimum.This optimum may be explained considering kinetics of the reactions. Atlow pressure the reacting particles have larger mean free path and thushiger energies for reaction. Increasing pressure increases reactantconcentration but reduces its energy. Thus ancptimum regarding the pro-pagation of the reaction is obtained as seen in figs 1-3.

Figures 4-6 show the changes in the microhardness, deposition rate andchlorine content as a function of the input power. As expected fromthe previous explanation, increasing rf power increase the value of the•nicrohardness to a value of 3000kg/mm2,fig.4. The deposition rate hasa somewhat unexpected behaviour, in fig.5 the round points curve risesto 100 watts and then drops moderatly with increasing power.

rhis may be explained by looking at the deposition rate on an earthedsubstrate, (square points),which rises with power. The increase in powercauses particles to be more energetic, and the rate reaction by ion mole-cule mechanism is increased on account of radical molecule mechanism.\s expected, from the above explanation chlorine content increases withincreasing power as shown in fig.6. Figure 7 describes the microhardnessas function of feed ratio. It can be seen that increasing NH3 concentra-tion 'increases microhardness to a maximum and then drons. This is becauseDf the stoichiometry of the reactants, and the formation of compoundsjther than Si3N4 at non adequate concentration.

A fructured britlle silicon nitride, indicating its high microhardness,is shown is the SEM micrograph of fig.8. Ttcan be observed that the coa-ting is of substantial thickness (*>»8 Urn) and it is well adherent to themetal substrate. The film is dense with high homogeneity and has nopores in it. The surface of the film shows that its build-up followsthe topography of the metal surface, again indicating the adherent natureof the coating.

269

REFERENCES

1. Engineering Property Data on Selected Ceramics. Vol.1, Nitrides,

Battelle Colombus Lab. NTIS AD 23773 p. 47 (1976).

2. L.J. BawenandT.G. Carruthers, J. Mater. Sci. 13, 684-687, (1978).

3. K. Niihara and T. Hirai,J. Mater. Sci. 12,1243-1252,(1977).

4. F. Galasso, U. Kuntz and W.J. Croft,J. Am. Ceram. Soc. 55,431,(1972).

5. W.A. PlisMng.J. Vac. Sci. Technol. 14, 1064,(1977).

6. C.J. Mogab and E. Lugujjo J. Electrochem. Soc. 127, 1853,(1980).

7. A. Raveh, Y. Hornik, U. Carmi, A. Inspektor and R. Avni,

To be published in plasma chem. Plasm. Proc. (1984).

8. R. Manory, A. Grill, U. Carmi and R. Avni, Plasma Chem. Plasma Process.

3, 235, 0-983)

9. Y. Ron, A. Raveh, U. Carmi, A. Inspektor and R. Avni3 Thin Solid

Films, 107, 181, (1983)

10. N. Mayo, U. Carmi, I. Rosenthal and R. Avni. To be published

J. Appl. Phys. Apr. (1984).

2 3 4 5

GAS P R E S S U R E ( T D I I )

270

FIG 1FIG 2

1 2 3 4 5

GAS PRESSURE ( T o r i I

Fig.l. Microhardness (Vicker<?' hardness at a load of 50gf) of SiN films

vs. pressure in the reactor for a floating substrate (430 C)

(SiCl4-to-NH3 ratio, 1/3).

Fig.2. Growth rate (dh/dt) of SiN filr vs. pressure in the reactor for

a floating substrate (430°C) (SiCl4-to-NH3 ratio, 1/3).

FIG 3

1 2 3 4 i 6GAS PRESSURE , lo i r

</>2D0Du>ILJ

aa.<i

oa looo

oi

FIG 4

100 150

RF POWER Watt

Fig.3. Chlorine content Ccl of SiN film vs. pressure it: the reactor for afloating substrate (430°C) (SiCl4-to-NH3 ratio, 1/3).

Fig.4. Microhardness (Vickers1 hardness at a loadof 50 gf) of SiNfilm vs. r.f. powe* for ajfoating substrate (pressure 5 Torr;SiCl4 - to - NH3 ratio, 1/3).

271

4a.

u>o0.

FIG 5

100 150RF POWER (Watt)

FIG 6

100 150

R F P O W E R ( W a t t )200

Fig.5. Growth rate (dh/dt) of SiN film vs. r.f. power (pressure; 5 Torr;SiCl4-to-Ml3 ratio, 1/3): grounded substrate; floating substrate.

Fig.6. Chlorine content Cci of SiN film vs. r.f. power for a floatingsubstrate (pressure, 5 Torr; SiCl4-to-NH3 ratio, 1/3.

iiooo

FIG7

3 4NH 3 /SiCI 4 RATIO

Fig.7. Microhardness (vickers' hardness at a load of 50 gf) of SiN filmvs. NH3-to-SiCl4 ratio (floating substrate, pressure, 5 Torr:r.f. power, 100 W).

Fig.8. An SEM micrograph showing a top view of a fractured SiN coatingon steel at an inclination of 60 .

272

BORIDATION OF STEELS IN LOW PRESSURE RF PLASMA

A. Raveh, A. Inspektor, U. Carmi, E. Rabinovitz and R. Avni

Nuclear Research Centre - Negev, P.O.Box 9001, Beer-Sheva 84190, ISRAEL

Prepared for Presentation at the Second Israel Materials EngineeringConference, Ben-Gurion University, 1984.

ABSTRACT

Iron bor ide films were produced by the bo r ida t i on of m a r t e n s i t i c s t a i n -less s t e e l in low temperature environment ( l e s s than 500°C). The processwas c a r r i e d out by the r . f . plasma of BCl3/H2/Ar a t low p r e s s u r e . Theformation r a t e , microhardness and chlor ine content of the f i lms weredetermined and cor re la ted to t he working parameters . A maximum of forma-t ion r a t e [4-5 pm/h) and micxohardness (2600 kg/mm2) was found.

1. INTRODUCTION

Boride films provide hard coatings(1,2) with high oxidation and chemicalresistance C 2~ 3)• Their properties make these materials attractive forus^ in corrosive and abrasive environment. Today iron boride films areproduced by the electrolytic process(4), by chemical vapour deposi-tion^^) or by the pack-cementation technique(1). In these methodsboron diffuses into the ferrous material at high temperature (above800OC) and reacts with the substrate material. At these conditions ofhigh temperature damage to the substrate material and its pre-treatmentwill occur.

The r.f. plasma technique is a convenient environment for low temperaturesurface boridation(7). Boron is supplied to the plasma from the decompo-sition of a boron compound such as diborane (B2Hg) or boron trichloride(BCI3) in an argon or argon-hydrogen mixtures.

This work sets to present the r.f. plasma boridation of stainless mar-tensitic steel (AISI-410), and to check the influence of the workingparameters on the surface boridation.

2. EXPERIMENTAL

The experimental set-up is shown elswhere(8), however for the convenienceof explanation the main features are outlined below. The substrate(AISI-410) was place inside a pyrex reactor (40 mm long and 8 mm in dia-meter) , in which plasma was induced by a r.f. coil wound around it andenergized by 0.5 MHz generator. The reactor was pumped down to workingpressure by rotary vane pump. 10 vol.% BC13 (Matheson C.P) premixed with^rgon (Matheson, U.H.P) are further diluted by argon and mixed with hydro-gen in a mixing chamber and flown into the reactor. Before boridationprocess the steel samples were mechanically polished to 25 ym (SiC) or to0.05 at' (AÔj), ultrasonicly cleaned in organic solvent and pretreatedin avgon plasma for 50 min.

273

In the boridation process a mixture of BCl3+H2+Ar introduced to theplasma. The concentration of BCI3 in the gas mixture was 4 vol.% whilethe total gas flow rate was kept constant at 200 seem"1, during 2hprocessing. The following parameters were varied:i) total gas pressure in the plasma reactor, between 2-6 Torr;ii) r.f. power input expressed as current flowing in the r.f. coil,

between 30-50 A.The iron boride films were analyzed by x-ray diffraction for theircrystallographic identification, by scanning electron microscopy (SEM)for thickness and morphology and by energy-dispersive analysis of x-rays(EDAX) for chlorine contents in the films. Microhardness were determinedon Vickers scale at 50 g constant load.


Two crystalline structure of iron boride were identified by x-ray diffrac-tion: FeB, orthorombic (a0 = 4.053 A, b 0 = 5.495 A and Co = 2.946 A) andFe2B, tetragonal (a0 = 5.099 A and bo = 4.240 A). Figure 1 shows theiron boridefilms on a martensitic stainless steel (AISI-410) substrate.It can be seen that the film is well adherent to the substrate and hasa non-porous structure. The surface of the boride film shows that itconsists of small microscopic spherical shapes, and reproduce thegeometric features of the substrate.

Table 1 shows the anlaysis of the iron boride films at different locationsin the plasma reactor witn regards to r.f. coil: in position H (-1.0 cmaway from the begining of the F 1.0, upstream), in positi^:1 G, in themiddle of the r.f. coil and F 1.0 cm away from the end of the r.f. coil,downstream. The formation rate of the films do not strongly differbetween the location H and G, but the chlorine content reaches a maximumin position G. The substrate temperature was the highest in position G,where the highest r.f. energy is delivered to the nlasma and substrate.

From table 1 and from previous work(8) it was evident that position H isfavorable for higher formation rate and minimum chlorine content in thefilms. Therefore, the results which follow will be presented only forsteel substrate located in position H.

A maximum rate of boride formation was obtained at gas pressure of 3 Torrand shown in Fig. 2. It should be mentioned that the boride films wasformed on both sides of the steel substrate, and the film thickness wasabout 60% on the steel side facing the gas flow and 40% on the backe.

The maximum formation rate was obtained for the steel samples polishedto 25 jjm by SiC i.e. for a rougher surface [Fig. 2).

The 25 ]im surface had a higher surface area available for the gas-metalinteraction. Since in the formation of iron boride the sole source foriron is the metal substrate, a higher process rate expected.

The variation of the formation rate as a function of the current in ther.f. coil is shown in Fig. 3. Figure 3 shows a similar behaviour forpressures of 3 Torr and 5 Torr. A maximum of 4.5 ym/h was obtained at41 A and 3 Torr and at a substrate temperature of 440±20°C (Figs. 2 and3). This behaviour indicate that the formation rate reaches a maximuma 41 A, and is not increases with increasing the r.f. power.

274

The behaviour of the chlorine content in the boride film is shown in Figs.4 and 5 as a function of gas pressure and current in the r.f. coil, res-pectively. The higher chlorine content in the boride was obtained on thesteel substrate polished down to 0.05 ym. In other words, in order to re-duce the chlorine content a rough surface is favorable. The chlorineconcentration increases linearly with increasing the current in the r.f.coil, as shown in Fig. 5. It should be noted that increasing the currentfrom 37 A to 46 A increases temperature of the substrate from 370°C to500°C, and thus diffusion rate of chlorine content increases too.

Figure 6 shows the behaviour of the microhardness of the films as a func-tion of the total gas pressure. At 2-3 Torr the microhardness reachesits maximum values, and is an order of magnitude higher compared to thenon-boridized surface of the AISI-410 steel (i.e. 2600 kg/mm2 vs.210 kg/mm2, respectively).

4. CONCLUSION

Iron boride fi'ms was identified by X-ray diffraction as Fe2B and FeB atboridation of stainless martensitic steel (AISI-410) in a low pressurer.f. plasma. The starting materials were gas mixtures of boron trichlo-ride (BC13) and hydrogen in an argon plasma. A maximum of formation rate(4.5 ym/h) and microhardness (2600 kg/mm2) obtained at current flowingin the r.f. coil of 41A and gas pressure of 3 Torr, whereas the substrateswere maintained at relatively low temperature (440±20°C) compared toother techniques used for boridation.

REFERENCES1. S.C. Singhal, Thin Solid Films, 45 (1977) 321.2. R.H. Biddulph, Thin Solid Films, 45 (1977) 341.3. E. Randich, Thin Solid Films, 63 (1979) 309.4. H.C. Fiedler and R.J. Sicraski, Met. Prog. 99 (1979) 101.5. H.O. Pierson and A.W. Mullendore, Thin Solid Films, 72 (1980) 511.6. L. Randich, Thin Solid Films, 72 (1980) 517.7. A. Raveh, A. Inspektor, U. Carmi and R. Avni, Thin Solid Films 108

(1983), 39.8. A. Raveh, Y. Hornik, U. Carmi, A. Inspektor and R. Avni, to be

published in Thin Solid Films, 1984.

Table 1Properties of Iron Boride films at different locations.Boridation time 2h, current in r.f. coil 39A, feed composition BCI3(4 v/o)+ H? (17 v/o) + Ar (79 v/o), gas flow 127

Substrateposition

gaspressure

(Torr)

Filmthickness

(ym)

chlorinecontent

(w/o)

substratetemperature

C°C)

microhardness

(kg/mm2)

HGFH*

4441

7615

24503010

440550400451

20007009001500

* H2/BC13 = 6/1, No argon

+ Values are in weight percentage of Fe in the iron boride films.

275

2 3 4 5

gas pressure (Torr)

Fig.l. An SEM micrograph showing a top view of an iron boride film on

a martensitic stainless steel (AISI-410) substrate.

Fig.2. Formation rate (dh/dt) of iron boride films vs. pressure (r.f.current of 44A and 3.9vol % BCI3); - - -samples polished to25pm by SiC; samples polished to 0.05™ by AI2O3.

k 2 -

40 44

current rf coil (A)

Formation rate (dh/dt) of iron boride films vs r.f. currentat 3 Torr and 5 Torr, and at 3.9 vol % BCl,.

276

na- pressure (Ten-)

6

o

2

1

_ I 1 "'

I /yx

y ~

-

current rf CDII (A)

Fi'_».4. Chlorine content in bo-ride films vs. pressure (r.f.current of 44A and 3.9 vol %BC13);—Samples polished to 25 pm

by SiC.Samples polished to 0.05ymby Al203.

Fig-5. Chlorine content in boride films vs. r.f. current at 3 Torrand 3.9 Vol% BC13.

Fig.6. Microhardness in boride films vs. pressure at r.f. currentof 44A and 3.9 vol% BCI3.

277

LASER INDUCED COPPER ELECTROLESS PLATING

S. Tamir and J. Zahavi

Israel Institute of Metals, Technion, Haifa, Israel

INTRODUCTION

During the last few years, laser induced plating processes have beenfound to be a very attractive technology in producing high rate,highly selective deposits, mainly for the electronics and micro-electronics industry. High rate metal plating from aqueous solutionson thin metallic substrate was reported by R.J. Gutfeld et al (1-5).Plating rate of more than 1000 times the rate of conventional platingwas found (1) in the irradiated area when argon laser was used duringcopper electroplating process on copper thin fiim substrate laid onglass. It was proposed [2) that primarily thermal and heating effectsassociated with the laser radiation were responsible for the highplating rate.

This work aimed at studying laser induced copper electroless platingprocess on polymeric substrate such as polyiir.ide (Kapton). Polyimideis used as a substrate for flexible printed circuit boards.

EXPERIMENTAL

A schematic representation of the experimental set-up is illustrated inFig. 1, Two types of laser were used in .nis study. The first onewas a continuous argon laser operating at wavelength of 0.5145 microns,while the second one was a Nd+3/YAG laser used in the pulsed mode atwavelength of 0.53 microns, pulse duration of 150 ms and repetitionrate of 1-5 KHz. The plating solution used was a commercial copperelectroless plating solution, while the plating processes were carriedout through various stages, consisting of specimen cleaning followedby immersion in sensitization and activation solutions, and finallyimmersion in the electroless plating solution. In the first part ofthe work, argon laser was used when the radiation of the substratetook place during its immersion in the plating solution. In thesecond part of the work, pulsed Nd+^/YAG laser was used to radiatespecimen surfaces before the plating process, while being immersed indeionized water immediately after the activation process. Scanningelectron microscopy (SEM) and Auger electron spectroscopy (AtiS) wereused to examine deposit morphology structure and composition.


Typical SEM observation of laser induced electroless copper depositionline is shown in Fig. 2. Deposit lines of this type were built usingargon laser radiation on the specimen during its immersion in the

278

electroless plating bath. This irradiated line produced at 7.9 x 103

watt/cm2 showed cracks and microcracks probably in the polyimidesubstrate. Copper deposit up to 60C A0 in thickness was detected atthe irradiated line zone by AES examination. AES surface analysis ofthe deposit line and of the polyimide between lines are shown aftersputtering in Fig. 3. Copper was found in the bright irradiated linewhile no copper was found in the background. The sputtering was per-formed in order to eliminate the copper which might have remained fromthe plating solution. AES depth profiles of laser deposit line aud theline-free polyimide surface are shown in Fig. 4. Deposit line containedabout 70% atomic concentration of copper and its thickness reachedvalues up to 600 A0 (as shown in Fig. 4). Deposit thickness on non-irradiated surface areas was found to be around 150 A0. The calculatedplating rate was four times greater at the irradiated areas compared tolaser non-irradiated areas.

In the second part of the work, copper deposit areas were producedwith pulsed laser system. A pulsed Nd+3/YAG laser was used atrepetition rate of 5 KHz, pulse width of 150 nsee and intensity of0.7 watt (experiment No. 31). Immersion time of the specimen in thewater was around 10 minutes, while plating time was around 5 minutesin the electroless solution. SEM observation of copper deposit linesis shown in Fig. 5. Copper deposited at the irradiated lines whileno deposition occured at the non-irradiated area. Deposit structurewas characterized by fine and dense grains, where the grain size didnot exceed half a micron, as shown clearly in Fig. 5b. Thesegrained copper deposit lines had sharp edges with continuous structure.AES depth profile of the deposited line is shown in Fig. 6.

The analysis showed copper concentration profile with thickness rangingfrom as much as 60% atomic concentration up to around 20% atomic concen-tration (AC) at thicknesses of 200 A° and 1600 A°, respectively.However, at the non-irradiated areas (Fig. 7") , the copper concentrationranged from 30% (AC) to 10% (AC) at depths of around 100 A° to 700 A°,respectively.

CONCLUDING REMARKS

Highly selective high rate electroless copper deposition on polyimidesubstrate was obtained with C.W. Argon laser irradiation.

Laser irradiation took place at the final stage of the plating, i.e.,after the substrate was immersed in sensitization and activationsolutions.

Irradiated zones were electroless plated at very high rate of deposi-tion in the value of about 120 pm/sec compared to about 10 pm/min inconventional electroless. Furthermore, the electroless plating wasselective and took place primarily at the local zones that wereirradiated with the laser beam.

The major effect of laser irradiation was probably the generation oflocal heating at the substrate - electroless plating solution inter-face resulting in high rate deposition. While the electroless solution

279

was kept at relatively low temperature of 15°C to 18°C, its platingreaction rate was very low and almost non-active. However, localselective laser radiation resulted in local temperature rise followedby high rate of electroless deposition as was observed.

Laser irradiation on the activated substrate surface resulted inplating deposition on laser irradiated zones while the noa-irradiatedzones were not plated during the plating stage. Pulsed laser irradia-tion on specimen surface was carried out while the activated substratewas immersed in distilled water. It might be that the activatormaterial was washed out by the water while laser irradiated zonesmaintained their activator material followed by electroless deposi-tion in these zones.

REFERENCES

1. R.J . Von Gutfeld, R.E. Acosta and L.T. Romankiw, "Laser-EnhancedPla t ing and Etching: Mechanisms and Applicat ions" , IBM, J . Res.Develop. Vol.. 26, No. 2, March (1982).

2. I .C. Puippe, R.E. Acosta and R.J. Von Gutfeld, " Inves t iga t ion ofLaser Enhanced Elec t rop la t ing Mechanism", J . Electrochem. Soc.Vol. 128^ No. 12, 2539-2545 (1981).

3. R.J. Von Gutfeld and L.T. Romankiw, "Laser-Enhanced P la t ing -Applications t o Gold Pat terning", Gold Bull . 15 (4) , (1982) .

4. R.J. Von Gutfeld, E.E. Tynan, R.L. Melchet and S.E. Blum, "LaserEnhanced Elec t ropla t ing and Maskless Pat tern Generation", App.Phys. Lett 35 (9) , November 1979.

5. L. Kulynyck, L. Romankiw and R. Von Gutfeld, "Laser-EnhancedExchange P la t ing" , IBM Technical Disclosure Bul le t in , Vol. 23,No. 3, August 1980.

280

Fig. 1: Schematic illustration ofexperiments1 set-up.

CU

BLACK LINEAFTER SPUT

600 840 1080

KAPTON AREAAFTER SPUT

Bo .1C ,ISO 360 600 840 1080

KINETIC ENERGY. Ev

Fig. 3: AES surface analysis oflaser induced copper electro-less plating on polyimidesurface after sputtering90 nm.(A) Laser line irradiated

area on polyimidesurface. The presenceof copper is shown.

(B) Non-irradiated areaof polyimide surface.No copper was detected.

Fig. 2: Lines of laser inducedcopper electroless platingon polyimide substrate.(A) Irradiated line showing

morphological damagein form of microcracks.

(B) Enlargement of (A).

281

0 9

100

t°60

20 -

B 0

18 27 36 45 54 63 72 81SPUTTER TIMEtMIN.)

-

w

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Iβ 27 36 45 54 63 72 81SPUTTER TIME (MIN.)

Fig. 4: AES depth profile analysis of laser induced copperelectroless plating on polyimide.

(A) Laser l ine i rradiated area.(B) Non-irradiated area.

Fig. 5: SEM observations of pulse laser induced copper electrolessplating on polyimide substrate. Laser irradiation took placeafter activation process.

(A) Typical deposit at various areas.

(B) High magnification of the deposit area shown in (A).

282

10 15 20 25 30SPUTTER (MIN.)

35 40

Fig. 6: AES atomic concentration depth profile ofpulse laser copper electroless plating onpolyiraide. Laser irradiation took placeafter activation process.

10 15 20 25 30 35SPUTTER TIME (MIN.)

40

Fig. 7: AES atomic concentration depth profile ofpulse laser copper electroless plating onpolyiraide. Non-irradiated area.

283

SURFACE HARDENING OF STEEL BY BORIDING IN A COLD RF PLASMA

I. Finberg, R. Avni and A. Grill

Materials Engineering Department, Ben-Gurion University of the NegevBeer-Sheva, Israel

T. Spalvins and D. Buckley

Tribology Branch, NASA-Lewis Research Center, Cleveland, Ohio, U.S.A.

ABSTRACT

Samples of 4340 steel have been plasma borided in cold rf plasma ofdiborane and argon. The plasma was initiated in a gas mixture of B2H^+Arinductively coupled to 27.12 MHz rf generator. The steel samples weretreated in the plasma at a net power input of 500W for 5 hours, the bulktemperature of the samples reaching values of 550-600°C. The influence ofthe position of the sample in the reactor and of the gas pressure wasinvestigated.

The plasma borided surfaces were studied by scanning electronmicroscopy, X-ray diffractometry and microhardness measurements. It wasfound that as a result of the decomposition of the diborane in the plasma,boron is deposited on the surface of the steel substrates and twocrystalline phases are formed: tetragonal Fe2B and orthorhombic FeB. Dueto the formation of the boride phases, the surface microhardness increasedfrom the original value of 270 Kg-mm-2 to a maximum value of 790 Kg-mm-2.

INTRODUCTION

Boriding, or boronizing, has become an effective means for wearprotection in many fields of industry (1-3). In addition to the increasingof the surface hardness, the boriding improves the resistance of steelto certain types of corrosion facids and molten light metals) (2,3) andimproves the oxidation resistance up to 800°C (4). The high hardness ofthe borided layers is attained directly through the formation of boridesand the process does not require subsequent quenching (5).

Boriding is performed in a gas, molten salt or pack cementationprocess at temperatures above 900°C (2,5,6). At these temperatures theeffect of previous heat treatment will generally be lost. In order toobtain the required bulk properties,heat treatment has to be done afterboriding, with the possibility of spalling of the hardened surface layer.Boriding at lower temperatures could therefore be advantageous. Duringboriding of steel two phases are formed, Fe2B and FeB. The simultaneouspresence of two phases affects adversely the quality and wear behavior ofthe borided samples (1,2) and attempts are therefore made to obtain onlythe less brittle

284

Boriding by a cold plasma process can be advantageous from twoaspects:(i) Since the availability of the active species in the plasma is

determined by the electrical discharge, the temperature of thosubstrate can be controlled independently and can be kept atrelatively low temperatures, determined by the required diffusionrate of boron into the steel.

(ii) By controlling the plasma parameters, it may be possible to produceonly one boride phase, like in the ionitriding process (7), resultingin an improved wear behavior of the borided samples.

The aim of the present study is to investigate the feasibility ofplasma boriding of steel and to determine the effect of the plasmaparameters on the boriding process and properties of the borided surface.

EXPERIMENTAL

The experimental set up for plasma boriding is similar to the onedescribed elsewhere (81• The plasma was initiated in a premixed gasmixture of 2.87% B2H6 in Ar with a 27.12 MHz rf generator, inductivelycoupled to the quartz reactor through an impendance matching unit. Thetreated samples were prepared from 4340 steel sheet and polished with1 pm diamond paste. Before plasma boriding the samples were cleaned inacetone and ultrasonically cleaned in ethanol. The bulk temperature ofthe substrates was measured with a chromel alumel thermocouple in contactwith the substrate.

X-ray diffractometry was used to identify the crystallographicphases formed at the treated surface. Its morphology was analyzed byscanning electron microscopy, SEM. The microhardness of the treatedsurface was measured, normal to it, with a Vickers diamond indenter at aload of 20 g. Preliminary Auger Electron Spectroscopy, AES, analysis wasalso performed to determine the concentration of the elements, close tothe surface of the borided samples.

The decomposition of the diborane, I^Hg, in the plasma, thedeposition of boron on the steel substrate, and its diffusion into thesteel are strongly dependent on the plasma parameters. In the presentrfork the effects of two plasma parameters were investigated:

(i) The position of the sample relative to the rf coil and direction ofgas flow:H - at the entrance to the rf coil (upstream).G - at the center of the rf coil.F - at the exit of the rf coil (downstream).

(ii) The total pressure of the gas in the reactor which was varied between1 and 5 mbars. The other plasma parameters were kept constantthrough this study, namely:Total flow = 5 seemNet Power Input = 500 WPower Density = 2 W-cnr3

Treatment Time = 5 hours

The temperatures reached by the substrates during plasma boriding were550-600°C.

2S5


Figure 1 presents the surface morphology, typical for the plasmaborided samples. The appearance of the surface indicates that anoverlay coating was deposited onto the substrate, where the growthproceeded through solidification of spherical globules. The size ofthese globules varied between 0.6 and 1.8 ym, however no systematicdependence of the size of the globules and either position of substratein the plasma, or gas pressure could be determined from the availabledata. Further investigation is required to find whether the plasmaparameters affect in a systematic way the morphology of the surface.

Preliminary AES analysis indicated that the layers close to thesurface of the plasma treated samples are rich in boron (more than 90%).Depth profiling indicated that the concentration of boron decreases andthe concentration of iron increases into the sample as shown in Figure 2.

Figure 3 presents an X-ray diffractogram obtained from the surfaceof a plasma borided 4340 steel sample. As indicated in the figure, X-raylines identified as belonging to the tetragonal Fe2B and orthorhombic FeBphases are obtained from the plasma borided sample. Although, asindicated by the AES measurements (Figure 2) , boron rich layers exist onthe surface of the treated samples, only X-ray diffraction lines of 4340steel, Fe2B and FeB have been obtained. This indicates that the boronatoms not bonded in Fe2B or FeB form an amorphous boron layer.

The results of SEM, AES and X-ray diffraction indicate that:(i) As a result of the excitation of the plasma in the B2H5 -i Ar

mixture, the diborane decomposes and boron species deposit on thesteel substrates. These boron species may be either BnHm fragments(n < 2 , m < 6) or boron atoms. Identification of the BnHm speciesrequires mass spectrometrical study of the plasma,

(ii) During the continuing deposition of the boron species on the steelsubstrates, a diffusion process of both boron and iron atoms occurs,resulting partly in the formation of the two crystallographic phases,

and FeB.

The microhardness of the treated surfaces has been measured normallyto the surface. A though the measurements made normal to the surface willgive values which to a certain extent are averages between the hardenedlayer and the bulk, these values can allow a qualitative comparisonbetween surfaces treated under different conditions.

In addition to the contribution of the bulk to the measured micro-hardness values, these values are also affected by the morphology of thssurface. Due to the roughness of the surface as shown in Figure 1, themeasured microhardness values are lower than the microhardness of thematerial itself. The microhardness values of the treated surfaces shouldbe compared to the microhardness of the untreated steel samples, ofVHN = 270 Kg-mm-2. In order to check for an eventual heat treatmenteffect on the microhardness, the microhardness was measured also on thesurfaces of the treated samples which were not exposed to the plasma. Itwas found that the microhardness on the unexposed surface was.the sameas that of the untreated sample, thus excluding any heat treatment effectson the microhardness of the plasma borided samples.

Figure 4 presents the microhardness values as a function of theposition of the sample in the reactor, for different boriding pressures.The results indicate that the hardening effect decreases from position H(entrance of the rf coil) towards position F (exit from the rf coil). Thevalues of the surface ir.icrohardness attained at a given position depend on

286

the gas pressure during the plasma boriding. As can be seen the highestvalue of 790 Kg-mm" 2 (as compared with 270 Kg*mm~2 for the untreatedsample) was obtained at position H and 3 mbar.

Taking into consideration the effect of the microhardness of thebulk and of the surface roughness, as mentioned before, the microhardnessof the treated surface layers is higher than 790 Kg-mm-2. This indicatesthat, as a result of the exposure of the steel samples to the diboraneplasma, a significant hardening of the surface is obtained. Thehardening of the surface layer occured at bulk temperatures lower than600°C, as compared to conventional boriding temperatures of 900°C.

Continuing studies are in progress to find the plasma parameters forfurther increase of the surface hardness and for the formation of asingle boride phase in the plasma borided layer.

ACKNOWLEDGEMENT

The authors are grateful to the US-Israel Binational Science Foundationfor supporting the work by a research grant 2813/82.

REFERENCES

1. K.H. Habig, Materials in Engineering, 2 (2), 1980.2. P. Goeuriot, F. Therenot and J.H. Driver, Thin Solid Films, 78, 1981,

67.3. R.H. Biddulph, Thin Solid Films, 47_, 1977, 341.4. L.S. Lyankhovich and S.S. Bragilevskaya, in "Protective Coatings on

Metals", ed. G.V. Samsonov, Consultants Bureau, NY 2, 1970, 123.5. T.S. Eyre, Wear, 34, 1975, 383.6. N.N Golego, A.P. Epik, V.D. Derkach and'V.F. Labunets, Idem 4, 5,

1973, 257.7. B. Edenhofer, Heat Treatment of Metals, 1974, 59.8. R. Avni, U. Carmi, R. Manory, E. Grossman and A. Grill, Proc. ISPC VI,

Eds. M.I. Boulos and R.J. Munz, p. 820.

Figure 1Surface Morphology of PlasmaBorided 4340 Steel p = 3mbarposition - H

287

Figure 2.Composition of plasmaborided surface asdetermined by AES withdepth profiling.

o

40 120 200 280 360 440 520 600

SPUTTERING TIME, mm

Figure 3.X-Ray diffractograms of 4340 steel .

Figure 4.Surface uicrohardness ofplasma borided 4340 s tee las a function of positionin the reactor.

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288

CURRENT METALLIZATION ISSUF-S IN MICROELECTRONIC DEVICES

K.N. Tu

IBM Thomas J. Watson Research CenterYorktown Heights, New York 10598

The current trend in very-large-scale integrated Si devices is miniaturiza-tion. Since it is not accompanied by a reduction of the temperatures used infabrication and operation, there are acute structural and compositional instabili-ty issues due to thermal stress and diffusion. Three issues on metallization willbe briefly reviewed. They are the delamination in polycide gates induced bythermal stress, the failure in conducting lines due to electromigration, and thedegradation of ohmi' or Schottky contacts caused by interdiffusion. Theseissues are generic and persistant in the miniaturization of device structures.

The application of silicide in FET gate has been dominated by the com-bined use of refractory metal silicide and heavily doped poly-Si as the so-calledpolycide gate.1-2 The silicides such as WSi2 and TaSi2 must be a good conduc-tor, chemically stable with the poly-Si so that the gate oxide will not be affect-ed, and capable of surviving a high temperature process without transforminginto oxides. For these requirements, the conduction and oxidation behavior3 ofrefractory metal silicide have been studied. The oxidation of silicide is alsointeresting for the formation of a surface oxide layer of gate insulation. Whena composite of refractory metal disilicide (with excess Si) and poly-Si is oxi-dized, a layer of SiO2 grows on top of the disilicide by consuming the excess Siand the poly-3i. Since Si is the dominant diffusing species in the refractorymetal disilicides, the transport of Si will cause a reverse flux of vacancies tocondense at the interface between the disilicide and poly-Si interface, weaken-ing the adhesion across the interface. Therefore, peeling occurs if the gate ishighly stressed. The gate tends to experience a high thermal stress since a hightemperature (~1000°C) annealing is required to improve the conductivity ofthe disilicide film. The disilicide film is typically amorphous in the as-depositedstate. Upon annealing, the amorphous film crystallizes into a highly defectivestructure, and it is only after the high temperature annealing that the defectsare reduced and the conductivity greatly improved.4 At present, how we canlower the annealing temperature without sacrificing the conductivity is an

289

important issue of the polycide gate. Also, the choice of disilicide is unsettledyet.

Electromigration in Al thin films can be retarded by adding a few percentof Cu.5 To be effective, the Cu addition has to go specifically to grain bounda-ries in the Al film but not to the interior of Al grains. This is because electro-migration at low temperatures ~100°C in Al films occurs mainly along grainboundaries, and any additive of Cu into Al grains will raise line resistance. Thegrain boundary diffusional flux due to electromigration can be decreased byreducing the grain boundary diffusivity or the effective charge number6, or byincreasing the grain size. The beneficial effect of alloying Cu to Al has beenfound largely due to the reduction of grain boundary diffusivity of Al. Howev-er, this effect alone will be insufficient, especially in resisting eiectromigrationdamage in one micron or submicron lines. For this reason, additional improve-ments of using bamboo-type microstructure, very fine grains, and sandwichstructure have been developed. In the sandwich structure, the middle layerserves as a diffusion barrier in preventing openings from running across theentire film,7 so its lifetime is improved. Clearly, the issue of electromigrationwill remain and it will become more serious at places where current crowdingoccurs such as a stepped surface.

Interdiffusion occurs naturally in a multilayered thin film structure. Theconsequence of interdiffusion, for example, of Al to a PtSi/n-Si contact is tochange Schottky barrier height. It can be detected readily by current-voltagemeasurements and analyzed by the concept of parallel contacts.8'9 To preventthe Al-pene*ration, most often a diffusion barrier is introduced to slow downinterdiffusion.10'12 However, we must understand the unique behavior ofinterdiffusion in thin films in order to solve the problem satisfactorily. Thinfiim interdiffusion tends to form a single intermetallic compound rather than toform all of them simultaneously as in bulk diffusion coupies. The phenomenonof "single" intermetallic compound formation has been verified by using highresolution transmission electron microscopy of atomic resolution.13 A kineticexplanation has been given assuming that the growth is limited by diffusion aswell as interfacial reaction.14 The consequence of "single" compound forma-tion is that the boundary conditions of its growth kinetics are simple and wecan analyze it with diffusion markers to unravel the intrinsic diffusioncoefficients.15 Furthermore, the marker analysis can be repeated in each of thecompounds formed in sequence. The selection of the first compound and thesequence of the followers have been found not to depend on driving force (freeenergy change)16 but rather on kinetics.17 This has been supported by study-ing interdiffusion in Al-Cu thin film bilayers; the compound CuAU which hasthe fastest interdiffusion coefficient is the one which forms first1 \ Knowingthe intrinsic diffusion coefficients, the diffusion flux and in turn the failure ratecan be estimated. However, to reduce the diffusion flux and to develop aneffective diffusion barrier are still important technological issues.

290

REFERENCES

1. B.L. Crowder and S. Zirinsky, IEEE Trans. Electronic Devices, ED26,369 (1979).

2. S.P. Murarka, J. Vac. Sci. Technol. 17, 775 (1980).3. J.E.E. Baglin, F.M. D'Heurle and C.S. Petersson, J. Appl. Phys. 54, 18 49

(1983).4. T. Tien, G. Ottaviani, and K.N. Tu, J. Appl. Phys. 54, 7047 (1983).5. F.M. d'Heurle and P.S. Ho, chapter 8 in "Thin Films - Interdiffusion and

Reactions".. Ed. by J.M. Poate, K.N. Tu and J.W. Mayer, Interscience -Wiley, N. Y. (1978).

6. H.B. Huntington, in "Diffusion in Solids - Recent Developments" Ed. byA.S. Nowick and J.J. Burton, Academic Press, N.Y. (1975).

7. J.K. Howard, R.F. Lever, P.I. Smith and P.S. Ho, J. Vac. Sci. Technol.73,68 (1976).

8. I. Ohdomari and K.N. Tu, J. Appl. Phys. 57, 3735 (1980).9. J.L. Freeouf, T.N. Jackson, S.E. Laux and J.M- Vvoodall, J. Vac. Sci.

Technol. 21 (2), 570 (1982).10. M-A. Nicolet, Thin Solid Films 52, 415 (1978).11. M. Wittmer, J. Noser and H. Melchior, J. Appl. Phys. 52, 6659 (1981).12. K.N. Tu, Chapter 7 in "Preparation and Properties of Thin Films", Ed. by

K.N. Tu and R. Rosenberg, Academic Press, N.Y. (1982).13. K.N. Tu, G. Ottaviani, U. Goselu and H. Foil, J. Appl. Phys. 54, 758

(1983).14. U. Gosele and K.N. Tu, J. Appl. Phys. 53, 3252 (1982).15. U. Gosele, K.N. Tu and R.D. Thompson, J. Appl. Phys. 53, 8759 (1982).16. K.N. Tu, G. Ottaviani, P..D. Thompson and J.W. Mayer, J. Appl. Phys.

53, 4406 (1982).17. H.T.G. Hentzell, R.D. Thompson and K.N. Tu, J. Appl. Phys. 54, 6923

(1983); H.T.G. Hentzell and K.N. Tu, J. Appl. Phys. 54, 6929 (1983).

291

ELECTRICAL PROPERTIES OF Ti :W SC1IOTTKY BARRIER CONTACTS TO SILICON

M.O. Aboelfotoh and K.N. Tu

IBM Thomas J. Watson Research CenterYorktown Heights, New York 10598 (U.S.A)

ABSTRACT

Ti:W alloy films sputter-deposited on chemically -leaned surfaces of n- andp-type Si(100) have been annealed at temperatures up to 7 5 0 ° C . Schottky-barrier height measurements were performed over a wide range of temperatureusing forward c . tent-vol tage characteristics. The barrier height on n-type Siwas found to increase and that on p-type to decrease with increasing theannealir g temperature. These changes in the barrier height can be interpretedin terms of the changes in the interfacial oxide on the Si surface upon anneal-ing. The barrier height on n-type Si was also found to decrease with increasingmeasurement temperature, and that on p-type Si to be independent of themeasurement temperature. The temperature dependence of the barrier height isshown to be nearly equal to that of the energy band gap in Si.

INTRODUCTION

Films of Ti:W alloy have found application as a diffusion barrier betweenplatinum silicide contacts and aluminum interconnecting lines in very-large-scale integrated devices.1 This alloy is also attracting increasing interest2 forapplication as low Schottky-barrier contacts to silicon with stable electricalcharacteristics at high temperatures. In this contribution the measurements ofthe Schottky-barrier height and its dependence on temperature for Ti-W alloyon n- and p-type silicon are reported. The effect of annealing at temperaturesin the range from 4 0 0 ° C to 750°C on the current-voltage characteristics of iheTi:W-Si Schottky-barrier contacts is also reported.

EXPERIMENTAL PROCEDURE

The Ti:W alloy films were prepared by d.c. magnetron sputtering. Films1000-1500A thick were deposited on O.OO5-£2cm n + and p + (100) oriented Siwafers with 2-jum-thick, 10-ficm epitaxial layer. Buffered H F was used to

292

clean the Si wafers immediately before they were loaded in the depositionchamber. Prior to deposition, the deposition chamber was evacuated to lxl0"7

Torr before admission of lOmTorr of argon of 99.99 percent purity. TheTi3oW7O target was given a 30 min sputter clean against a shutter at 3 W/cm2

before deposition. The Ti:W was sputtered at 0.5 W/cm2 resulting in a depos-ition rate of 40A/min. The wafers were grounded during the Ti:W deposition.The resistivity of the Ti:W alloy films was about

During the same deposition, two types of samples were made: bare Sisamples for x-ray diffraction, Rutherford backscattering spectrometry (RBS)and Auger electron spectroscopy (AES) analysis, and SiO2 covered and pat-terned Si samples for I-V characteristics. For the I-V measurements thedeposition was made through a metal mask having 2.5-mm diameter openingsbut the active area of the contacts was defined by oxide windows with diameterof 130, 250, 500, and 1000 micrometers. The samples were annealed attemperatures in the range from 400°C to 750°C in a furnace with flowinghelium purified by hot (950°C) Ti particles.

Composition analysis of the Ti-W alloy film and detection of compoundformation upon annealing were made using MeV 4 He + Rutherford backscatter-ing spectrometry. The structure of tho phases in the film was determined byglancing-angle x-ray diffraction in a Seeman-bohlin geometry diffractometry.

The forward I-V characteristics were analyzed according to the thermionicemission theory3 of Schottky-barrier current transport:I = aT A* exp ( - q<£BO/kT) exp (qV/nkT), where V>3kT/q is the forwardapplied voltage, <£BO is the effective barrier height at zero voltage, A* is thetheoretical effective Richardson constant for Si, T is the measurement tempera-ture, a is the contact area, and n is an ideality factor introduced in order totake into account deviations from ideal diode behavior. The value of $ B Q wasdetermined from an extrapolation of the linear portion of a In I-V curve to zeroapplied voltage. The extrapolation was made using a curve fitting program,which fits the experimental data points by exp (qV/nkT). The slope of thelinear portion yielded the value of n. The I-V measurements were made withthe samples held at temperatures between -90°C and +20°C. The measure-ments were made on all four diode areas.

Experiments were performed to investigate the importance of leakagecurrents at the edge of the Ti:W-nSi contacts. The edge leakage occurs as aresult of the high field due to accumulation of the n-type Si surface by thepositive charge in the SiO2 surrounding the contact.4-5 It has been shown4'5

that for contacts to n-type Si, this edge effect dominates the I-V characteristicsfor reverse-bias voltage and for low forward-bias voltage. We have plotted ourmeasured reverse currents at -0.1V against contact diameter and found that thedata lie along a straight line with a slope equal to 1.7 indicating that almost all

293

of the current is flowing across the whole area of the contact. If, on the otherhand, the reverse current was dominated by the edge effect, the data would liealong a straight line with a slope equal to unity. In Ti:W-pSi contacts, the edgeeffect is absent because the surface of the Si at the edge of the contact isdepleted in this case. The depletion does not give rise to high fields as doesaccumulation in the case of n-type Si. Indeed, a plot of our measured reversecurrent against contact diameter for Ti:W-pSi contacts showed that the data liealong a straight line with a slope equal to 2, showing that the reverse currentsare proportional to the contact area. We concluded that the edge effect is notimportant in determining the I-V characteristics of the Ti:W-Si contacts.

EXPERIMENTAL RESULTS AND DISCUSSION

X-ray analysis on the as-deposited Ti:W alloy films showed the presence ofa bcc solid solution of Ti and W2 . This phase persisted upon annealing up to700°C.

01 0.2 0.5 04FORWARD BIAS (V)

Fig. 1. Forward I-V characteristicsTor 400°C annealed Ti:W Schottkybarrier contacts to n-type Si as afunction of the measurement temp-erature.

0.5

294

The forward I-V characteristics of Ti:W schottky-barrier contacts to n- andp-type Si annealed at 400°C for 1 hr. are presented, respectively, in Figs. 1 and2 as a function of the measurement temperature. The results show that theideality factor n is slightly temperature dependent. The deviation of n fromunity can be caused by image force and the presence of an inlerfacial layerbetween the metal and the semiconductor. The contribution of field depend-ence of image-force lowering of the barrier height to n does not exceed onepercent for barriers with the Si doping concentration (10 1 5 cirr 3 ) underconsideration.3

Auger depth profile analysis revealed that an interfacial layer containingoxygen is present between the Ti-W and Si even after annealing at high temp-eratures. The presence of an interfacial layer is well known to cause fielddependence of the barrier height3. The existence of a field dependence of thebarrier height must then represent the largest contribution to deviations of nfrom unity. Since the barrier height obtained by extrapolation to zero voltageis affected by the electric field in the depletion region, the zero-electric-fieldbarrier height and not the zero-voltage barrier height should provide a bettercharacterization of the metal-semiconductor Schottky barrier. Several authorshave modeled the electric-field dependence of the barrier height.6"8 Morerecently, Wagner et al.9 derived a simple expression relating the zero-voltagebarrier height 4>BO and ideality factor n (as determined from the forward I-Vcharacteristics) to a zero-electric-field barrier height c/>BF defined under flat-band condition:

l n (^j n - *ype Si

(1)

p — type Si

where N c and N v are, respectively, the effective density of states in theconduction and valence band. The change in N c and Nγ with temperaturevariation is calculated assuming a dependence law of T 3 / / 2 ' " . N D and N A are,respectively, the donor and acceptor concentration. The values of the barrierheights 4>gp and <£gF based on equation (1) and deduced at various measure-ment temperatures from the curves of Figs. 1 and 2 for the 400°C annealedsamples shown, respectively, in Figs. 3 and 4. It is seen that the barrier heighton n-type Si decreases with increasing measurement temperature, and that thebarrier height on p-type Si is independent, within the experimental errors, ofthe measurement temperature.

295

£ 10

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In order to understand the temperature dependence of <£BF, a comparisonis made between the experimentally determined values of <£gF and thosecalculated on the assumption that it is related to the dependence of the Sienergy band gap Eg on temperature. The calculated values are normalized tothe experimental value of <£BF determined at -82°C. The reported temperaturedependence of Eg in Si by Bludau, Onfon, and Heinke11 is used. The resultsshow that the decrease in <>Sp with increasing temperature (-2.2xlO"4eV/°K),in the temperature range -82 C to +20°C, is very close to the change in Eg.11

Hence it can be concluded that the temperature dependence of the barrier"weight of Ti:W on n-type Si, is nearly equal, within the experimental errors, tothe temperature dependence of Eq.

In Figs. 3 and 4, the values of the barrier heights 0 B F and <>BF for theas-deposited samples and for samples annealed at 500°C to 700°C are alsoshown as a function of the measurment temperature. It is seen that the valueof <£BF increases and that of <£gF decreases with annealing up to 700°C, with

296

0RF + ^BF ecJual t o Eg- In the as-deposited and all annealed samples, $ B F isindependent of the measurement temperature, and the change in $ B F is nearlyequal, within the experimental errors, to the change in Eg. These results implythat the Fermi level at the Ti:W-Si interface is pinned in relation to the valenceband edge. Auger depth profile analysis revealed that the oxide on the Sisurface is gradually reduced and the oxygen is dissolved into the Ti:W, and thatintermixing of Ti:W and Si at the interface does not occur upon annealing up to700°C. It is likely that this change in the interfacia! layer (regarding thegradual reduction of the oxide on the Si surface) can change the pinningposition of the Fermi level and hence the barrier height (Fig. 3) upon anneal-ing.

RBS and x-ray analyses showed, after annealing at 75O°C, the formation ofa ternary silicide phase (TixWj_x) Si2(x~0.2) with a hexagonal WSi2

structure.2'12 The values of the barrier heights 0 B F and <>BF for the Ti:Wsilicide are also shown, respectively, in Figs. 3 and 4 as a function of themeasurement temperature. As in the case of the Ti:W alloy, <|>gF is independ-ent of the measurement temperature, and $ B F decreases with increasing themeasurement temperature with a coefficient of -2.7xl0'4eV/°K. This is againnearly equal to the change in Eg. These results, imply that the Fermi level atthe Ti:W silicide-Si interface is also pinned in relation to the valence bandedge. In addition, the results of Figs. 3 and 4 show that the values of0g F and $ B F for Ti:W after annealing at 700°C, are essentially the same asthose when the silicide phase is formed. Auger depth profile analysis revealedthat oxygen is expelled from the interface and is dispersed throughout thereacted layer as a result of silicide formation. The result that changing theinterface from that of Ti:W on Si to that of Ti:W silicide on Si does not changethe value of the barrier height indicates that the barrier height is pinned at avalue which depends on the Ti:W alloy and not on the silicide phase. Weconclude that it is the Ti:W alloy which is important in determining th electricalproperties of the Schottky barrier in the Ti:W silicide-Si system.

ACKNOWLEDGEflENTS

The authors gratefully acknowledge helpful discussions with P.S. H o . It isalso our pleasure to acknowledge R . D . Thompson for valuable assistance inmaking the electrical measurements , J .E . Lewis for A E S da ta , P. Saunders forRBS data and CSS Materials Labora to ry at Yorktown for t he Ti :W deposi t ion.

REFERENCES

1. P.B. Ghate, J.C. Blair, C.R. Fuller and G.E. McGuire, Thin Solid Films53, 117 (1978).

2. S.E. Babcock and K.N. Tu, J. Appl. Phys. 53, 6898 (1982).

297

3. E.H. Rhoderick, Metal-Semiconductor Contacts (Clarendon, Oxford,1980).

4. A.Y.C. Yu and E.H. Snow, J. Appl. Phys. 39, 3008 (1968).5. A.M. Cowley, Solid-State Electron. 12, 403 (1970).6. J.D. Levine, J. Appl. Phys. 42, 3991 (1971).7. C.R. Crowell, Solid-State Electron. 20, 171 (1977).8. J.M. Shannon, Solid-State Electron. 19, 537 (1976).9. L.F. Wagner, R.W. Young and A. Sugerman, IEEE Electron Device Lett.

EDL-4, 320 (1983).10. S.M. Sze, Physics of Semiconductor Devices, Wiley-Interscience, New

York (1981).11. W. Bludau, A. Onton and W. heinke, J. Appl. Phys. 45, 1846 (1974).12. F. Nava, C. Nobili, G. Ferla, G. Iannuzzi, G. Queirolo and G. Celotti, J.

Appl. Phys. 54, .7.434 (1983).

TEMPERATURE C O-80 -60 -40 "20

0 6 5

- 500-C

S 0.60

0.55

750-C

A—

" AS

. 0A

- a_ O

V- v

-DEPOSITED —'

AS-DEPOSITED4OO°C- Ihr5OO°C-Ihr600°C-lhr700°C-lhr75O°C-3hr

1 I l 1 I

0 65

0 60

0 50 —

200 250TEMPERATURE(K)

Fig. 3. Dependence of barrier height onmeasurement temperature in Ti:W-ntype Si Schottky-barrier contacts ann-ealed sequentially from 400°C to750°C.

0 45'150

-100T"

TEMPERATURE C O-80 -60 -40 -20

i r

Oo

A AA V-

V V

O AS-DEPOSITEDA 400°C-lhrV 700°C-lhrT 750°C-Jhr

roID00

200 250TEMPERATURE (K)

Fig. 4. Dependence of barrier height onmeasurement temperature in Ti:W-p typeSi Scholtky barrier contacts annealedsequentially from 400°C to 750°C.

299

INTERFACIAL REACTIONS IN LASER ANNEALED Ni/GaAs CONTACTS

A. L a h a v ( a \ T. B r a t ( a \ C. Cytermann (b) and M. E i z e n b e r g ( a ' b )

(a) Dept. of Materials Engineerings Technion, Haifa(b) Solid S ta te I n s t i t u t e , Technion, Haifa

ABSTRACT

Thin films of Ni on (100) GaAs substrate were annealed by Nd-YAG pulsedlaser (A=0.53ym). Morphology and phase composition of the reactionproducts as a function of laser energy density were studied by a numberof analytical techniques, including Transmission Electron Microscopy,Auger Electron Spectroscopy and X-ray diffraction. The results show thelack of adhesion between melted Ni-film and GaAs, indicating that thistechnique is unsuitable for preparation of Ni/GaAs contacts.

Nickel is usually added to the eutectic composition of Au and Ge in thefabrication of ohmic contacts to n-type GaAs by short duration meltingat a temperature of 450°C (1,2). It was found that Ni forms (Ni,Ge)Ascompounds at the interface with GaAs (2,3) and thus it is considered tobe a "wetting agent" which prevents "balling" and improves surfacemorphology of the contacts. The curiosity about the role of Niinitiated a research on the interactions between Ni and GaAs both insolid and liquid states. Here we present the results on liquid phasereactions, and these on the solid state will be reported somewhere else.

Pulsed laser annealing seems to be a good technique for investigation ofmetallurgical aspects of the interaction of melted Ni with GaAs, since themelting temperature of Ni (1453°O is higher than that of GaAs (1238°C).tfelting of the localized area on the surface of the sample by laser beamfollowed by very fast solidification is an extremely nonequilibriumprocess but in some cases it improves contact morphology and electricalproperties (5-7).

Ni films 1300 A in thickness were evaporated by an electron gun on clean(100) GaAs single crystal substrates at a vacuum of 1.10 torr. Laserannealing was performed by Q-switched Nd-YAG laser (wavelength 0.53]im)tfith pulse duration of 150 nsec. The laser beam was focussed on thesample surface with a lens of f=75 mm, resulting in a beam diameter(at 1/e intensity) of 64ym. The average laser energy density was about6 J/cm^ and it was attenuated by a polarizer. It should be noted thatthe laser power was not stable enough and the energy density of individ-ual pulses varied by *v< 10% of the average value. Area coverage wasobtained by computer controlled movement of the sample on a X-Y table.

Characterization of the reaction products was carried out by Nomarskinicroscopy, Scanning Transmission Electron Microscopy (STEM), X-raydiffraction, Auger Electron Spectroscopy (AES) and Energy DispersiveX-ray Spectroscopy (EDS). TEM specimens were prepared by chemicalthinning from the back side of the sample.

X-ray diffraction pattern of the as deposited sample (Fig.la) showspeaks of polycrystalline Ni and a very wide peak, the (400) reflection,of single crystal GaAs. TEM examination (Fig.2a) reveals the average

300

grain size in the Ni film to be of the order of 200A. AES depth profil-ing shows that the interface between Ni and GaAs is abrupt and free fromcarbon and oxygen (to the limit of detection of this technique *v> 1 at.%).

Morphological examination of the annealed films by Nomarski microscopeshows that at laser energy densities less than 0.4 3/cvP- the Ni filmremains smooth and undisturbed. At an energy of 0.4 J/cm^ small hillocksof l-4ym in diameter appear on the sample surface. TEM examination ofsuch a hillock reveals an increase in the average grain size up to lOOOA(Fig.2b). Electron diffraction pattern from this area yields only Nirings, and EDS analysis in the STEM mode shows that this area containsonly Ni. The above results suggest that these hillocks represent areaswhere Ni film was melted and recrystallized.

In some places, where the laser energy was larger than the average, thecontinuity of Ni film is broken (Fig.3a) and it is folded to the periph-ery of the melted spot, leaving almost a flat surface of GaAs in thacentral area. Such a behaviour of melted Hi film indicates the lack ofadhesion between it and GaAs, possibly due to large surface tension, ordue to undetected- interfacial contamination.

At laser energy density of 0.6 J/cm^ the whole surface is covered withmelted spots. The enlarged Nomarski image of a few spots is shown inFig.3b. It can be seen that the central area of a spot is quite flatbut covered with tiny ripples. The spots are separated by walls of3-5]im in thickness. X—ray diffraction of this sample (Fig.lb) yieldspeaks of Ni (much weaker than in the as-deposited sample), polycrystal-line GaAs5 and also additional peaks that can be indexed as belongingto a hexagonal phase with lattice parameters a=3.78 and c=5.01 A.A similar phase, the ternary compound Ni2GaAs, forms as a result of solidstate reaction between Ni thin film and GaAs (4,9). AES high resolutionanalysis shows that the central area of the laser spot is almost free ofNi and contains only Ga and As (Fig.4a), whereas the walls around thelaser spot contain about 50% Ni as well as Ga and As, and are probablyresponsible for the X-ray reflections of the ternary phase.

TEM bright field images of the characteristic structures in the centralarea of laser spot are shown in Fig.5. Fig.5a shows dark ripples on awhite background. Electron diffraction from the background givespatterns of GaAs in single crystal and polycrystalline form. EDS analy-sis of the background and the ripples confirms the AES results (absenceof Ni), and shows that the ripples contain about three times less As thanthe background. Thus it may be assumed that the ripples are formedduring the regrowth of melted GaAs accompanied with Au evaporation fromthe surface. At some places the ripple area has a cell structure(Fig.5b) with Ga segregation on the cell boundaries, probably due toconstitutional supercooling (8). The area of the walls around the laserspot couldn't be analyzed by TEM, since it was too thick to be trans-parent to electrons.

In conclusion, it was demonstrated that the lack of adhesion betweennelted Ni film and GaAs makes the method of pulsed laser annealingunsuitable for preparation of Ni/GaAs contacts. The situation is quitedifferent in the case of solid state reaction, where Ni was found to bevery reactive and it forms an epitaxial compound at the interface with3aAs at 150°C (4,9). It should be noted that unlike nickel, germanium

301

being pulsed laser annealed at the same conditions does not have anyadhesion problems and forms a uniform reacted layer (4,7).

ACKNOWLEDGEMENT

The authors gratefully acknowledge the help of R. Brener in AES analysis

REFERENCES

1. N. Braslau, J. Vac. Sci. Technology, 19, 803 (1981).

2. G.Y. Robinson, Solid-State Electron, 18, 331 (1975).

3. M. Ogawa, J. Appl. Phys. 51, 406 (1979).

4. Authors unpublished data.

5. G. Eckhardt, in "Laser and Electron Beam Processing of Materials",C.W. White and P.S. Peercy, eds., p.467. Academic Press, New York(1980).

6. 0. Aina, J. Norton, W. Katz and K. Rose, in "Laser and Electron-Beam Interactions with Solids", B.R. Appelton and G.K. Celler, eds.(1982).

7. J.S. Williams, in "Laser Annealing of Semiconductors", J.M. Poateand J.W. Mayer, eds., p.383. Academic Press, New York (1982).

8. A.G Gullis, in "Laser Annealing of Semiconductors", J.M. Poate andJ.W. Mayer, eds., p.147. Academic Press, New York (1982).

9. M. Ogawa, Thin Solid Films, 70, p.181 (1980).

Fig.l. X-ray diffraction patterns, a) as deposited, b) annealedat 0.6 J/cm2. o - Ni, A - GaAs, li- Ni-GaAs.

302

a

Fig.2. TEM bright field image of Ni film.a) as deposited, b) annealed at 0.4 J/cm

Fig.3. Nomarski micrographs of annealed surface by anenergy density of a) 0.4 J/cm^, b) 0.6 ^

303

450 720 990 1260KINETIC ENERGY , (EV)

Fig.4. High resolution Auger spectra of sample annealed at0.6 J/cm^. a) central area, b) wall area.

a

Fig.5. TEM bright field images from central nrea of laser spotat 0.6 J/cm2. a) Ga-rich dark ripples in GaAs singlecrystal, b) cell structure in the ripples.

504

ENHANCED OPTO-ELECTRONIC ACTIVITY OF SEMICONDUCTORS

BY MEANS OF (PHOTOELECTROCHEMICAL ETCHING

R. Tenne and V. Marcu

Department of P l a s t i c s ResearchThe Weizmann I n s t i t u t e of Science

Rehovot 76100, I s r a e l

The performance of e l e c t r o - o p t i c a l devices such as so l a r c e l l s ; l i g h temitting diodes; radiation detectors e tc . is determined, in part , byelectron-hole recombinations which occur at thei r interfaces. Varioustechniques were employed for the passivation of surface states andrecombination centers like dipping in metal ions solutions (1); bombard-ment with atomic hydrogen beams (2) and different etching techniques.

We have discovered (3) that a short photoelectrochemical etching ofCd-chalcogenide electrodes leads to a considerable improvement in the per-formance of photoelectrochemical cells based on these semiconductors. Thistechnique is of special importance to devices which are based on thin filmlayers of semiconductors where surface defects and grain bounderies haveadverse effect on the performance.

For yet uncomprehensible reason short photoelectrochemical etchingremoves these defects preferentially, whereas chemical etching f a i l s .

Figure 1 shows the I-V curve of an electroplated CdSeQ g 5T e

0 %<-electrode in polysulfide electrolyte. The beneficial role of thephotoetchning is clearly exhibited.

I t i s noteworthy that well over ten semiconductors which belong tofive different families have shown positive response to photoetching.These families include: Cd-chalcogenides; Zn-chalcogenides; oxides;ternary semiconductors and laminar materials. Whereas light was necessaryfor the photoetching of n-type semiconductors forward bias was used inthe electrochemical etching of p-type semiconductors.

Various spectroscopic techniques have been used recently in order toconfirm the selectivity of the photoetching towards surface s tates andrecombination centers. They include spectral response (4,5), electronbeam induced current (6); cathodoluminescence (7); electroluminescence;pho-tovoltage transients; photocapacitance (8) and electrolyte electroreflec-tance (9) • Although each technique employed different semiconductor anddifferent etching conditions were used, the emerging consequence was thatindeed surface states and recombination centers, which are not accessibleto regular chemical etchants, were removed upon short photoetching. Thecathodoluminescence signal of CdTe prior to and after photoetching isan example. The two-fold increase in the signal is mainly at tr ibuted tothe increased lifetime of minority carriers which is caused by the re-moval of surface detects.

305

Apart from the favorable effect of the (photo)electrochemicaletching on the performance of semiconductor devices there seems to besome fundamental questions which are related with this phenomena. First ,oxidative decomposition of the semiconductors seems to be more facileand selective than reductive decomposition by surface electrons. BothCdTe and InSe, which can be prepared as n or p-type materials byappropriate doping, were selectively etched by surface of p-type InP incomparison to i t s n-type counterpart, under illumination.

Second problem which is probably of even more fundamental natureis the appearance of a unique morphology of photoetched surfaces. Thedense pattern of etch p i t s (>109 cm" ) i s in general not sensitive tothe circumstances under which the photoetching was carried out. Recentlyi t was shown, however, that the etch p i t s density depends on the donordensity (conductivity) of the crystal (n-type) (12). Figures 2 and 3show the morphology of photoetched CdSe with two different donor densi-t i e s . Figures 4 and 5 show the morphology of photoetched CdTe with twodifferent donor densities.

We suggest (13) that the etch p i t pattern is the resul t of a non-uniform flow of charge carriers across the space charge layer of the semi-conductor during photoetching. This non-uniform flow results from non-uniformities in the electric fields which are induced by the localizedpositive charges of the donors (n-type) within the space charge layer.Evidence in support of this model comes from different sources. Forexample, if one applies a strong forward bias during the photoetching,the space charge layer diminishes and the density of the etch pi ts isreduced considerably. Also, the etch p i t density is similar fordifferent faces of the same crystal and i t is not sensitive to theduration of the photoetching; light intensity (over a certain threshold);electrolyte used, e tc .

The implications of this theory are not limited to Schottky barr iers .In fact non-uniform Fiicroscopic fields are expected in any semiconductorjunction and they influence the charge flow of both majority andminority carriers across the space charge layer.

Recent experiments in which the wavelength of the exciting lightis varied are in accord with the present model.

REFERENCES

1. B.A. Parkinson, A. Heller and B. Mi l l e r , J . Electrochem. Soc. , 126,954 (1979).

2. C.H.Seagar. and D.S. Ginley, Appl. Phys. L e t t . , 34_, 337 (1979).3. R. Tenne and G. Hodes, Appl. Phys. L e t t . , 37_, 428 (1980).4. R. Tenne, Y. Mirovsky, Y. Greenstein and D. Cahen, J . Electrochem.

S o c , U9_, 1506 (1982).5. Y. Mirovsky, R. Tenne, G. Hodes and D. Cahen, Thin Sol id Films, 9 1 ,

349 (1982).6. G. Hodes, D. Cahen and H. Leamy, J . Appl. Phys . , 54_, 4676 (1983).7. R. Tenne and A.K. Chin, Mater. L e t t . , 2, 143 (1983).

306

8. R. Haak, C. Ogden and D. Tench, in "Photoelectrochemistry: FundamentalProcesses and Measurement Techniques", eds. W.L. Wallace et a l . ,Elect-ochem. Soc. Ser. , 82-3, 486 (1982).

9. M.Tomkiewicz, W. Siripala and R. Tenne, J . Electrochem. S o c , inpress.

10. a) N. Miller and R. Tenne, Appl. Phys. Lett . , 39_, 283 (1981);b) R. Tenne, ib id . , 43, 201 (1983).

11. R. Tenne, B. Theys, J . Rioux and C. Levy-Clement, submitted to J.Appl. Phys.

12. R. Tenne, V. Marcu and N. Yellin, submitted to Phys. Rev. Lett.13. a) R. Tenne and G. Hodes, Surf. Sc i . , 135 , 453 (1983);

b) R. Tenne, H. Flaisher and R. Triboulet, Phys. Rev. B., in press.

V(mV).-i pt

Figure 1

(+) The electrode ?.fter preparation;(x) The electrode after mild etching;(o) The electrode after photoetching.

307

i&fiFigure 2 Figure 3

The photoetching of CdSe COOl) from Cleveland Crystals was performedin Na2S03 1M, at +1V v s . SCE and AMI l i g h t i n t e n s i t y .

Fig. 2: p = 1 £2-cm Fig. 3: p = 10 n> cm

Figure 4 Figure 5

The photoetching was performed in HNO,; HC1; H90 [1:4:10 v o l j , at +1V vs.SCE and AMI l igh t i n t e n s i t y .

Fig. 4: p < 5 Fig. 5: p ~ fl-cm

308

LASER INDUCED METAL DEPOSITION ON GaAs SUBSTRATES

J. Zahavi

Israel Institute of Metals, Technion, Haifa, Israel.

INTRODUCTION

Recently, much attention has been paid to laser-induced chemical andelectrochemical plating from aqueous solutions on metallic (1.) ,semiconductor (2-5) and polymeric (4) substrates for possible use inmicroelectronic circuits and devices. Use of C.W. laser to enhanceelectroplating or electroless plating of nickel,copper and gold fromaqueous solutions on thin metallic layers evaporated on glass sub-strates was carried out (1) for maskless microcircuit repair and designchanges.

Pulsed Q-switched Nd /YAG laser irradiation induced metal or alloydeposition from aqueous electroplating solution or semiconductorsubstrates for potential formation of conductive lines, as well asohmic and schottky barrier contacts, have recently been reported (2-5).These investigations demonstrated that semiconductor substrates suchas Si (2), InP (3) and GaAs (4,5) could be selectively plated withPt, Au, Ni-Pt, from aqueous solutions, making use of laser beamirradiation without external electric current and without maskingprocedures.

Mechanisms and processes associated with metal deposition via pulsedlaser irradiation on semiconductor substrates immersed in standardplating solution without external electric current could be character-ized by three major stages:

a. Chemical reduction of metal ions from the electroplating solutionby photogenerated electrons (7,8) followed by metal atom depo-sition at laser irradiated zones.

b. Surface heating followed by interdiffusion of deposited atomsinto the substrate material (2,5).

c. Termination of laser pulse followed by rapid cooling rateresulting in formation of alloys or compounds of the semiconductorwith deposited diffused elements to give, for example, Pt or disilicides with silicon substrate (2) .

This study aimed at characterizing and studying factors affectinglaser-induced, highly selective, high-speed gold deposition fromstandard electroplating solution on undoped, n-type-and p-type GaAssubstrates.

309

EXPERIMENTAL AND PROCEDURE

A schematic outline of an experimental system used in conducting laserinduced plating is described in Fig. 1. The system consisted of aQ-switched Nd:YAG pulsed laser system, x-y computerizsd table and anelectroplating cell, all of which were controlled and monitored by acomputer (Fig. 1) .

Semiconductor GaAs flat specimens 15 mm x 15 mm in size were preparedfrom (a) intrinsic undoped GaAs wafer with resistivity of 10 Q cm,(b) N-type Tellurium doped GaAs wafer with resistivity of 2.1 x 10n.cm and (c) p-type zinc doped GaAs wafer with resistivity of5.5 x 10"3 fi.cm.

A commercially available potassium gold cyanide electroplating solu-tion produced by Lea-Ronal Company, New York (known as Aurospeed CVDelectroplating gold solution has been used in this study.

Specimens immersed in the plating solution were irradiated with pulsedlaser beam at wavelengths of 1.06 yin and 0.53 pm, pulse duration of15 ns, beam diameter of 0.5 mm and energy densitizer of 0.1 joule/cm^to 0.5 joule/cm2. Continuous deposit lines or squares were producedon the various substrates.

Gold deposit surface morphology structure and composition were examinedby microscopic and surface analysis techniques, such as scanningelectron microscopy (SEM), transmission electron microscopy (TEM).Auger electron spectroscopy (AES), Electron spectroscopy for chemicalanalysis (ESCA), Rutherford back scattering (RBS), and x-raydiffraction analysis.


Laser induced high speed highly selective deposition of gold on GaAssemiconductor substrate immersed in aqueous plating solution of goldwithout masking and external electrical current have been achievedin this work. Lines or squares of gold deposition were formed usinglaser energy densities of 0.1 to 0.5 joule/cm2 at wavelength? of1.06 ym and 0.53 ym, laser pulse duration of 15 us with pulse over-lapping of 95% with beam diameter of 0.5 mm. Gold deposit surface wascharacterized by the presence of craters or cell-like events of severalmicrons in size. These events resulted probably from heating ormelting processes associated with laser irradiation upon the GaAssubstrate surface, as shown clearly in Fig. 2. Furthermore, depositsurface roughness increased with increase of laser energy densitiesas observed clearly in typical SEM micrographs (Fig. 2). Typicalstructures of gold deposit produced via pulsed laser of GaAs wereexamined by TEM as shown in Fig. 3. Typical gold deposit structurewas characterized by small gold polycrystallites ranging from lessthan 100A° to a few hundred angstroms, as revealed by bright anddark field pictures in Fig. 3A and Fig. 3B, respectively. A typicalelectron diffraction of the gold deposit (Fig. 3C) revealed thepresence of pure metallic gold in the GaAs substrate. The presence

310

and the amount of gold denosit were also obtained by RBS techniaues.Typical RBS snectrum indicated the presence of gold on GaAs laserirradiated zones, as shown in Fig. 4. It might be noted that thepresence of elemental gold metal at the laser irradiated zones wasobtained by ESCA and by x-ray diffraction examinations (data are notshown).

The amounts of gold deposits were determined by RBS examinations. Theresults were correlated with laser energy densities as shown, for exanrale,in Fig. 5. Gold deposit concentration increased linearly with increaseof laser energy density. Furthermore, it was found that gold depositon N-GaAs subst.ate was found to be one order of magnitude higher in com-parison to gold concentration found on P-GaAs substrate under the samelaser operating conditions.

Typical results of current-voltage curves characteristic of Schottkybarrier contacts and of ohmic resistance are shown in Fig. 6 for laserinduced gold deposits on n-type Tellurium doped GaAs. Deposit resistancevalues were obtained from the slopes of 1-V curves (Fig. 6B) which weremeasured and recorded at various deposit lines. Deposit resistancevalues decreased linearly with increase of laser energy density as can beseen, for exanrnle, in Fig. 7.

CONCLUDING REMARKS

A. It has been demonstrated that gold could be deposited on GaAs semi-conductor substrates via pulsed laser irradiated upon the GaAsimmersed in standard electroplating solution.

B. Laser induced high rate highly selective gold deposition on GaAswas conducted in direct one ster> eliminating masking and externalelectric current.

C. Laser induced deposition took r>lace primarily through photoelectro-"hemical deposition, followed by inter-diffusion of gold atoms intothe GaAs substrate through a molten surface laver which solidifiedupon the termination of the laser pulses.

D. Laser gold deposit showed Schottcky behavior contacts on n-typeGaAs under the operational conditions employed in this work.

REFERENCES

1. T.C. Puippe, R.E. Acosta, R.J. Van Gutfeld, J. Electrochem. Soc,Vol. 128_, No.12, pp. 2539-2545 (1981].

2. Y.C. Kiang, J.R. Moulic, J. Zahavi, IBM Technical DisclosureBulletin, Vol. 26_, No.l, p. 327 (June 1983].

3. R.F. Karlicek, V.M. Donnelly, J. Appl. Phys. 53 (2), pp. 1084-1090(1982] . ' '

311

4. J. Zahavi, "Laser Induced Electrodeposition on Polyiraide andGaAs Substrate", Research Report, Israel Institute of Metals,Technion, Haifa, Israel; AFWAL-TR, 1983, Wright-Patterson AFB,OH 45433, U.S.A. (August 1983), to be published.

5. J. Zahavi, "Laser Induced Gold Deposition on GaAs Semiconductor",Research Report, Israel Institute of Metals, Technion, Haifa.Israel; AFWAL-TR-1983, Wright-Patterson, AFB, OH 45433, U.S.A.(August 1983), to be published.

6. J. Zahavi, S. Tamir, "Laser Beam Technology Promoting High Speedand Selective Plating Processes", Annual Technical Report, IsraelInstitute of Metals, Technion, Haifa, Israel; AFWAL-TR-83,Wright-Patterson AFB, OH 45433 (September 1983), to be published.

7. R.H. Michaels, A.D. Darrow, R.D. Rauch, Appl. Phys. Lett.,Vol. 39_, No. 5, pp. 418-420 (September 1981).

8. F.W. Ostermayer, Jr., P.A. Kohl, Appl. Phys. Lett., Vol. 39,No. 1, pp. 76-78 (1981).

312

Computer

Specimen^

Laser

W////M

PlatingSolution

• X-Y Table

Fig. 1: Schematic description of experimental set-up.

LASER PLATED QOLD ON N-Qui: Ta (100)

LASER WAVE LENGTH -1.08 pn ( * Iβ)

TEM OBSERVATIONLASER QOLD DEPOSIT ON UNDOPED QaAB

LASER WAVE LENGTH =0.53 pm:ENERGY = 0.23 JOULE/cm2

1.0 PULSE/SPOT: 60x60 PULSES OF0.5mm BEAM DIAMETER

0.11 JOUU/te1

Fig. 2: Typical SEM observationof laser plated gold onGaAs

Typical TEM observa-tion of laser platedgold on GaAs

313

RBS. PROFILELASER PLATED GOLD ONN-TYPE TELLURIUMDOPED GoAi (100)SUBSTRATE

1.200 l.GOO 2.0GQ

ENEHGY We VI

Typical RBS profile showing the presenceof gold at laser irradiated zones on GaAssubstrate.

ob 2

LASER PLATED GOLOCN N-GoAi (100)DOPED TELLURIUM

LASER ENERGY DENSITY (JOULE/cm2)

Fig. 5: Typical correlation between amount ofgold deposit and laser energy density.

314

Typical I-V curvemeasurements ofSchottky barriercontacts and ohmicresistance of golddeposits on GaAssubstrate.

LASER PLATED GOLD N-GaAl: T« (100)WAVE LENGTH = 0.S3 |im: 'LASER ENERGY =0.28 JOULE/cm2

DEPOSIT RESISTANCE AND CONTACTMEASUREMENT (# 9)

TYPICAL SCHOTTKV HARRIER CONTACT.TWO POINT PROBIIH MEASURE.

TYPICAL OHMIC RESISTANCE DETERMINEDFROM I-V CURVE MEASURED ON THESAME LINE.

Dependence of golddeposit resistanceon laser energydensity.

-,20q

a ,6

LASER PLATED GOLDON N-GoAiDOPED TELLURIUM

3

0.1 0.2 03LASER ENERGY DENSITY (JOULE/cm2)

315

VAPOR PHASE SOLDERING OF SURFACE MOUNTEDELECTRONIC ASSEMBLES

E. Falkenstein and I. Fainaro

filbit Computers Ltd.Materials Engineering Dept.

INTRODUCTION

The ever increasing need for more compact electronic circuitryDrought up new electronic packaging solutions, such as theleadless chip carriers and the surface mounting technology ingeneral. One of the principal steps in the implementation ofthis technology is the interconnecting stage of surface mountedcomponents to the wired substrate boards. This usually consistsof a soldering process which can be realized by a variety of.tieans, a highly effective one being vapor phase reflow soldering.

VAPOR PIW.SE APARATOS

The condensation soldering process was developed and introducedinto production in 1973 at the Western Electric Company. It hasbeen proven to be an effective mass soldering method (fig. 1) •The condensation soldering process, or, as more commonly knownas the Vapor Phase Soldering V.P.S. has been used successfullyto reflow many varieties of electronic and other assemblies.The basic system consists of a stainless steel chamber, immersionheating elements, and vapor condensing coils. A processing liquidis placed in the chamber and heated to boiling, thus producing azone of saturated vapor above the liquid. Parts to be reflowed areintroduced into this vapor which condenses, giving up its latentheat. Systems are unpressurized thus the system is always at theboiling point of the fluid. This affords absolute control of thetemperature, regardless of the length of the reflow cycle.Operating temperature is determined by the fluid selected. Fluidscurrently used are chemically inert, and create an oxygen freeheating environment. Since vapor condenses on all surfaces ofthe assemblies, heating is extremely uniform, independent of thegeometry of the product. Recent advances in V.P.S. equipment havegreatly improved the adaptability of this technology to theproduction environment.

PRIMARY FLUIDS

The primary fluid for V.P.S. Must have a boiling point sufficientlyhigh to consistently produce high quality metallurgical bondsbetween solder and metals (about 20-30°C above the liquidus of thesolder in use is generally acceptable).

316

Conversely the temperature must be sufficiently low to minimize thermaldamage to any part of the product. The vapor as well as the liquid mustbe nonflammable, inert, chemically and thermally stable, low in toxicityin the operating temperature range, and of course to have the right heattransfer properties at an affordable price.

SECONDARY FLUID

The principal purpose of the secondary vapor blanket is to act as astable barrier between the primary vapor and the surrounding air.Therefore it must have a density between that of the primary vapor andair. Because of the continuous boiling-condensing cycle it must havea boiling point lower than the primary fluid. Complete miscibility ofthe primary and secondary fluids are desirable.

THE ELECTRONIC ASSEMBLY

Since the whole purpose is to create a sound metallurgical joint betweenelectronic components as the Leadless Chip Carrier LCC and the PrintedWiring Board PWB, Let us take a brief look at those parts. The ceramicLCC may be thought of as the die cavity portion of the ceramic dualin line package (DIP) having conductor patterns extending out to theexterior portion of the package. The external conductor patterns arereflow soldered to the appropriate PWB surface termination pads. ThePWB that was chosed is a copper clad invar board coated with layersof polyimide which embeds the micro wire interconnection between theappropriate pads and the connector which provides the electricalconnection to the rest of the system. The LCC's and connector arereflow soldered to the PWB to form a module.

APPLICATION OF SOLDER CREAMS

Solder creams are used to provide the joining between the componentleads and PWB termination area. The creams can be applied usingmainly three basic methods: syringe application, stenciling andscreen-printing, we use the two latter ones. The screen basicallyi~ prepared in the following manner: A Photosensitive material isplaced on a taut screen in a metal frame, it is photo exposed anddeveloped to obtain the desired pattern. In case of stencil thepattern is etched in a sheet of stainless steel or brass ofsuitable thickness. The PWB is placed under the screen or stencilin excact registration with the screen image. The solder cream isdrawn across and through the screen by a squeegee with a sharprubber edge.

ASSEM3LY REFLOW 5 CLEANINGLCC components are placed at the appropriate pads in such a mannerthat there will be contact by means of the solder cream between thecorresponding pads on the PWB and LCC. The whole assembly is bakedto drive out the volatiles from the cream, then it is run throughthe reflow cycle to obtain a sound metallurgical bond. After solderingthe assembly is cleaned and inspected as required.

317

SOLDER CREAMS

Solder cream consist of flux binder and solder alloy. The rheologyof the flux binder is such that the dispersion of heavy solderparticles (85-90w/o) can be maintained under normal room temperatureand pressure conditions. The flux binder contains a minimum of 60%pure rosin and 40% other ingredients including activators, solvents,lubricants and thickeners. The alloy is dispersed in the cream inform of tiny particles.

REQUIREMENTS FROM A SOLDERED JOINT

1. Good metallurgical bond.2. Complete solder coverage ot all metalized areas.3. Adequate space between the PWB and LCC (a minimum of 0.025 mm of

solder)4. Good alignment of the termination pads.5. Nice appearance, no cracks.

SECONDARYVAPOR

(VAPOR BLANKET)

PRIMARYSATURATED

VAPOR

Fig. 1 "Batch type vapor phase reflow system

31S

EXPERIMENTAL

Various aspects of the soldering and related processes were studiedand their impact on the quality of the resulting solder joint wasinvestigated. Main issues of concern included activation and pre-tinning of components and substrates, solder paste materials andproperties, printing of solder pastes, placing of components on solderprinted boards, drying of solder pastes, vapor phase soldering,clean-ing of soldered assemblies, and environmental testing.

INITIAL CONDITIONING

Wired substrate boards are usually received with pre-tinned footprintpads. Their degree of solderability largely depends on the extent ofoxidation on their surfaces. In order to assure their fitness forconsequent solderinc operations a tinning action is commonly carriedout. Unless greasing or serious tarnishing occured, this processyields a fresh uniform coating of a highly solderable character.Preliminary cleaning of the boards in a vapor degreaser enhancesthe solderability and is common practice.Tinning of boards is conveniently accomplished by dipping, by wave-soldering, or by solder-paste printing followed by vapor phasereflow. We prefer the latter method since it performs more selectivelywith no need for masking, and also affords higher thickness of coating.The quantity of solder applied in this stage contributes substantiallyto the total amount of solder in the eventual joint. If solderabilityproblems are evident, a sequence of cleaning, masking and wave-or dip-tinning can be followed up and, if necessary, repeated again. Ceramicleadless chip carrier components, of the type we use, usually comewith gold plated castellations. Carefully handled, those componentsmaintain their good solderability properties for extended periods oftime. Prior to being placed on the substrate boards they aretinned by dipping or in a soiaer wave. This action dissolves theinitial gold plating. After tinning, flux residue contamination isremoved- in a vapor degreaser cleaning plant.

PRINTING THE SOLDER PASTE

Solder paste printing is successfully accomplished either byscreening or by stencilling. Screens are elasti-. and spring back betterthan stencils. On the other hand, stencils are more resistant andeasier to clean while permitting higher print thickness. We triedboth methods anc? -.and to adopt the latter.

The viscosity of the solder paste should comply with the printingmethod. We found a viscosity of 500000 - 600000 cps appropriate forstencilling ,and somewhat lower values for screening. Our experimentswith various thicknesses of stencils indicate that a thickness of0.2-0.3 mm is satisfactory with regard to uniformity and linedefinition.After the solder paste being printed, the components are placed onthe sucstrate boards using vacuum pincers. The process of placingthe components is critical to the extent that such defects as slighttwists, translations or excessive pressure on the components may

319

cause bridging and solder balling in the soldered assembly. From apoint of view of accuracy as well as speed, automation of thecomponent placing process is highly recommended.The substrate boards, bearing the components placed on the wet solderpaste, are transferred to a drying oven, heated to 80-100°C for15-30 minutes. The drying drives out volatiles which may impair thereflow jointing process. Care should be taken not to overheat theassemblies in the oven since this may result in excessive oxidationas well as spreading of the solder paste.

VAPOR PHASE SOLDERING

The dried assemblies are now ready for "soldering. We use a batch typevapor phase reflow soldering unit, manufactured by Hybride TechnologyCorp. (Mass. USA). It operates with Fluorinert FC-70 (3M brand)liquid,producing a primary vapor of 215°C. A secondary blanket zone isproduced of Freon Ti? vapors. Assemblies to be soldered are loaded ontoan elevator basket which is then lowered at a speed of about lm/mininto the vapor zone. It resides about 1/2 min in the primary vapor andis raised into the secondary vapor zone.- for a cooling period, priorto exiting the tank. The controlled and effective nature of heatingby this method prevents over-heating and permits perfect soldering evenfor very densely populated boards.

CLEANINGPost-solder cleaning of LCC assemblies is considerably more difficultthan that of the equivalent DIP, due to the fact that the contaminatedareas are hard to reach. In addition to being difficult to clean,contamination of surface mounted assemblies is also difficult to detect.The conventional conductance test may yield erroneous results sincecontaminants confined by the closely overlying components may not beeasily leached off. We have performed some conductance tests on LCCassemblies and found that the cleaning time to reach the allowed levelof ionic contamination is as much as three times longer than with DIPs.

CONCLUSION

Vapor phase soldering combined with solder paste printing prove toform an extremely effective process of interconnecting surface mountedelectronic assemblies. Still, its successful implementation depends onoptimizing all stages of the process.

REFERENCES

1. Carmen Capillo "The Assembly of LCC to PWB"Presented at IPC Fall meeting, October 1982,San Diego, California.

2. G.M. Wenger; L.L. Mahajan "Condensation Soldering"IPC Assembly - Joining Handbook.

3. D.J. Peck, D.J. Spigarelli "Development of Continuous Vapor PhaseSystems" HTC Pond Lane, Concord Mass.

4. F.J. Dance "Mounting LCC on PCBs'",Electronic Prod. June 1982.

5. C.A. Mackay "Solder Creams and How to Use them"Electronic Packaging and Production , Feb. 1981 .

320

LATTICE, GRAIN BOUNDARY AND SHORT-CIRCUITSDIFFUSION OF PHOSPHORUS IN TaSi2 THIN FILMS

J. Pelleg*

Mater ia l s Engineering Department,Ben Gurion Univers i ty of the Negev,

Beer-Sheva, I s r a e l

[NTRODUCTION

TaSi2 /n p o l y c r y s t a l l i n e Si(Poly-Si) was found [1] to be an a t t r a c t i v e gatemetallization in device fabrication. The reduction in the gate level sheetresistance to about 2n/d, while simultaneously maintaining other benefi-cial properties such as patterning, stability during device fabricationand the simultaneous retrofit to the conventional Si-gate very largescale ingegrated (VLSI) processes were the prime incentive to explore thepossibility of using TaSi2/n poly-Si for application in VLSI technology.

Information on dopant migration and i ts distribution in composite layergate materials is lacking. As the packing density of discrete units on3i chips increases the diffusion of dopant inpurities acquires increasingimportance. Thus, in a recent work P was found to diffuse rapidly throughthe WSi2 [2], causing threshold voltage changes in the MOS (metal-oxide-si liconj capacitors.

Fhe objective of this work was to identify the mechanisms contributing tothe overall transport, discriminate between them and derive diffusiondata for them. The possibility of using cosputtered silicides directlyas the gate metal on gate oxide cannot principally be ruled out [3] andtherefore the diffusion specimens used in the current experiments con-sisted of TaSi2/poly-Si and TaSi /SiO_ configurations.

EXPERIMENTAL

3i wafers were used as s u b s t r a t e s . A thermal oxide of 600 S thickness wasgrown on them in dry oxygen a t 1000°C. A 3000 ft undoped po ly-S i was thendeposi ted by low pres su re chemical vapor deposi t ion (LPCVD). On anothers e t of Si wafers only thermal oxide was grown to a 1000 8 th i ckness . Taand Si in a nominal r a t i o of 1:2 were cosput tered on t h e s e wafers to anominal thickness of ^> 2500 S. The wafers were then s i n t e r e d a t 900°C/30min in a flowing Ar ambient t o form the c r y s t a l l i n e TaSi_. The re su l t i ngc r y s t a l l i n e TaSi2 had a very fine grained s t r u c t u r e of aBout 0 .05-0.1 urn.The sheet r e s i s t a nc e of the TaSi7 /Si0~ was 2.75n/D and t h a t of TaSi ? /poly-Si was 2 . 2 n / D -

A l a y e r of 0.2 uC c a r r i e r free P rad io i so tope was deposi ted on each ofthe wafers by evaporat ion from a tungsten filament and i t was used as an

* Work performed in Bell Labora tor ies , Murray H i l l , NJ 07974, U.S.A.

321

instantaneous source of tracer in the diffusion measurements. In order toeliminate loss of P by evaporation, the surface of each wafer was cappedby a thin layer of about 60 X Ta.

Quadrant wafers of the TaSi2/poly-Si and TaSi/SiO2 specimens werg diffu-sion annealed simultaneously in the temperature range of 500-900 C forshort times in a flowing Ar aaibienf.. This temperature is below 0.5Tm,and therefore contribution from short circuits may be expected. Thespecimens were submicron sectioned [4] and the activity of P wasmeasured by liquid scintilation (LS) and also by the residual activity(RA) method.

RESULTS

For short-circuits (SC) and grain bounday diffusion (GBD) analysis of theexperimental data the solution of Whipple [4] - Suzuoka [5,6] was usedin the form [7,8] given by:

Here D, ., refers to the diffusivity either in the SCs (D,) or in GBsD is the lattice diffusivity, t is the diffusion annealing time, 5 isthe GB width, C is the concentration of the diffusant, y is the distanceand K is the segragation factor given by K=C,/C,. C, and C., are theconcentrations of the diffusant impurity in theuBs(or SCs) and/the lattice.K is usually unknown. According to this expression logC vs. Y shouldbe linear if D and D, ,, are independent of composition and position.From the slopes of such curves and known values of D, only the productK6D, ,, can be determined. Figs. 1 and 2 summarize the results of themeasurements by RA technique. Each profile of Fig. 1 can be dividedinto three parts: the near surface region (associated with the Ta cappingand no physical significance will be attached to this at present), thelinear tail and the intermediate region. There is a striking differencebetween the slopes of these curves and those shown in Fig. 2, where anadditional contributing mechanism seems to have been observed. This isindicated by the arrows. In order to evaluate KflD, ., from the profiles,D is needed and can be evaluated by the technique given by Pelleg [5].The temperature dependence of D, K<5D, and K'AD, (here K1 has the samemeaning as K and A is the width of the SC) forTaSi2/poly-Si andTaSi2/Si02 are given by:

TaSi2/poly-Si: D = 4.21xl0"12exp(-0.67/kT)cm2/sec

K<SDb = 1.34xl0"15exp(-0.52/kT)cm /sec

TaSi2/SiOs: D = 5.05xlO~12exp(-0.63/kT)cm2/sec

K<SDb = l.llxl0"17exp(-0.2S/kT)cm3/sec

IC'AD, = 2.68xl0'17exp(0.58/kT)cm3/sec

322

DISCUSSION

The predominant feature of all the penetration profiles of TaSi^/SiO- isthe existence of a central region having a larger slope, a shorter tailregion and an ill-defined near surface region if at all. It had to bepostulated that an additional SC diffusion (SCD) was operating. It isbelieved that this is a result of segregation effects. When a bufferpoly-Si exists between TaSi and the gate oxide, virtually an infinitesource of Si is available for the diffusion of Si into TaSi- during the900°C/30 min sintering, which can segregate into the open sites in GBsand SCs. The degree of plugging the high diffusivity paths in SCs andGBs depends on temperature and time. In TaSi^/poly-Si Sc channels wereplugged to such an extent that the effective paths for diffusion becomepractically the same as the regular paths for diffusion in the lattice.GB channels, however, were affected to such a degree only, whichpermitted a distinct observation of GB diffusion (GBD).

In the case of TaSi- no supply of Si was available to plug SCs and GBswhich remained open for rapid transport along their paths. Both, SCs andGBs could exert their effect on the overall diffusion as indicated bythe penetration profiles. Indirect evidence supporting this concept canbe inferred from experimental observations [1,3,6-7]. Also, recent workindicates that a solid solubility of Si in TaSi- exists [8]. In additionit has already been observed that the most stable TaSi. is not thestoichiometric but a Si rich silic:de [6].

SUMMARY

1) Si segregation has been suggested as the cause for the difference inthe penetration profiles characterizing diffusion in TaSi_/poly-Si andTSi^/SiC^ 2

2) Diffusion in TaSi 2 /poly-Si occurs v ia GBs and the l a t t i c e . The pre-sence of the poly-Si buffer layer a c t s as a source of Si for segregationin TaSi- which completely plugs SCs.

3. Diffusion in TaSi /SiO_ occurs in the l a t t i c e , SCs and GBs. Channelsfor rapid t r a n s f e r of P along them remained open due t o the absenceof Si segregat ion.

4 . Complete s a tu r a t i on of the open channels in TaSi /po ly -S i by Si orby a more e f fec t ive so lu te should slow down diffusion r a t e s of dopantsin TaSi

REFERENCES

1. S.P. Murarka, D.B. Fraser , A.K. Sinha and H.J. Levins te in , IEEE Trans.Electron Devices ED-27, 1409 (1980) .

2. M.Y. Tsai , H.H. Chack, L.M. Ephrath, B.L. Crowder, A. Cramer,R. S. Bennett, C.J . Lucchese and M.R. Wordeman, J . Electrochem. S o c ,128, 2207 (1981) .

3 . S.P. Murarka, J . Vac. Sci . Technol . , 17, 775 (1980).4. J. Pelleg and S.P. Murarka, J. Appl. Phys., 54, 1337 (1983).5. J. Pelleg, Thin Solid Films, 110, 115 (1983).

323

6. S.P. Murarka, D.B. Fraser, W.S. Lindenberger and A.K. Sinha, J . Appl.Phys., 51, 3241 (1980).

7. S.P. Murarka, M.H. Read, C.J. Doherty and D.B. Fraser, J . Electrochem.S o c , 129, 293 (1982).

8. J . Pelleg and L. Zevin, unpublished work, 1984.

Fig- 1.

P penetration pro-files in TaSi2/poly-Si specimens. Datapoints beyond thearrows are associa-ted with GBD.

00 10 20PENETRATION DISTANCE 6 ' 5 MO"6cme'5)

Fig. 2.

P penetration pro-f i les in TaSi2/Si02specimens. Threeregions are seen.The segment betweenthe arrows is rela-ted to SCD.

0.0 10 2 0 3 0 4.0PENETRATION DISTANCE6'5 (IO"6cm6"l

324

THE LIMM TECHNIQUE FOR DETERMINATION OF THE SPATIAL DISTRIBUTION

OF POLARIZATION AND SPACE CHARGE IN POLYMER ELECTRETS

Sidney Bo Lang

Ben-Gurion University of the Negev, Beer Sheva, Israel

A recently developed technique for determining the spatial

distribution of polarization in polymer electrets has been extended to

the case of combined polarization and space charge.

Recently, the Laser Intensity Modulation Method (LIMM) for the accurate

determination of the spatial distribution of polarization in polymer

electrets was developed by Lang and Das-Gupta (1). This paper presents

an extension of LIMM. to the case of combined polarization and space

charge. The experimental system is illustrated in Fig. 1. Its essential

features are the following,, A thin, electroded polymer electret film,

which possesses combined polarization and space charge, is mounted in an

evacuated sample chamber containing windows through which radiant energy

can be admitted. Each surface of the film, in turn, is heated by means

of a laser beam which is intensity-modulated at frequencies varying from

100 Hz to 100 kHz. The heat is absorbed on the electrodes and produces

temperature waves which are propagated into the film and are attenuated

more strongly as the frequency of modulation is increased as shown in

Fig. 1. Schematic diagram of experimental apparatus used in LIMM.

325

0-4 0-6 0-8

POSITION (x/L)

1-0

Fig. 2. Maximum amplitudes of thermal waves in a 25.4 y-m PVDF film.

0.5 1

POSITION (X/L)

Fig. 3. Assumed polarization distribution and calculated values.

Fig. 2. The nonuniform heating is convoluted with the polarization dis-

tribution to produce a pyroelectric current which is a function of the

frequency, and the thermal diffusivity and thickness of the film. An

analysis was developed for deconvoluting the current-frequency data to

give the polarization distribution in the form of a Fourier cosine

series (1).

The method is extended by using the previously-developed electrostatic

and thermal analysis to find a relationship between the pyroelectric

current and an integral which is a function of the space charge distri-

326

1§ o

ICO

EXAMPLE

\

\

IA

\

0 0.5

POSITION (X/L)

Fig. 4. Assumed space charge distribution and calculated values.

bution, the modulation frequency, and several thermal, optical, and

geometrical constants. The unknown space charge is represented as

a modified Fourier sine series. Finally, a linear relation is develop-

ed between the component of the current in phase with the laser beam

modulation and the Fourier coefficients. The combination of this re-

lation and the corresponding one developed for polarization are used

in a least-squares analysis to give both the polarization and the space

charge distributions.

The correctness of the theory and the stability and the accuracy of the

computational technique were verified by a simulation procedure.

Figures 3 and 4 illustrate an example of assumed polarization and charge

distributions, respectively, as shown by the solid lines. These values

were used to simulate the current-frequency curves which would result

from the combination of the assumed polarization and charge. The simu-

lated data were used in the least-squares analysis to recalculate the

polarization and charge, the results of which are ^hown as super impos-

ed points in Figs. 3 and 4. The agreement with the assumed distribu-

tions is very good except near discontinuities, thus demonstrating the

correctness of the theory*

Experimental measurements were made on a 25.4-ym film of polyvinylidene

fluoride (PVDF) which had been poled with a voltage of 1000 V for 15

min. at 110°C, and then cooled with the field applied. Calculated

50

2 5 "

NCE

< ooQ_

-25

(b)

0

327

- K = 0.001 l = 0.002540.8K OR 1.1L1.ZK OR 0.9L

0.5

POSITION (X/L)

Fig. 5. Experimentally measured polarization distribution.

?

bJ

:E

CHAR

G

<m

0

-1

-2

(c)

- J

7/-/

/ '/f

• K = 0 . 0 0 1• 0 . 8 K- 1.2K

OROR

,

l=0.00254UL0.9L

liIkft

-

O 0.5 1

POSITION (X/L)

Fig. 6. Experimentally measured space charge distribution.

polarization and space charge distributions are shown in Figs. 5 and 6.

The effects of errors in the thermal diffusivity and film thickness are

illustrated and can be seen to be negligible.

k detailed mathematical analysis of the method and an interpretation

of the experimental data will be presented in a forthcoming publi-

cation.

Reference

1. S.E. Lang and D.K. Das-Gupta, Ferroelectrics (in press).

328

CHALCOGENIDE INFRARED GLASS FIBERS

A. Bornstein, N. Croitoru

Department of Electron Devices, School of EngineeringTel-Aviv un ive rs i ty , Tel-Aviv 69978, I s rae l

In t roduc t ion The h i s t o r i c a l evolution of o p t i c a l fiber communicationsystems as an a l t e r n a t i v e to radio and copper-based t ransmiss ion l i n k shas r e c e n t l y been marked by the search for ma te r i a l s which exh ib i t lowDptical l o s s at progress ively longer wavelengths. Since 1980 the pro-gress in mate r ia l s development has been spurred by the need both infiber o p t i c s and in high power IR lase r systems for t r ansparen t mate r i -a l s with absorption l e v e l s in the range of = 10~ cm" and down,four orders of magnitude below the l e v e l s required in conventional op-t i c s . These goals have now been achieved by chemical p u r i f i c a t i o n ofNad KC1, KB", CaF^, SiF .

The lo s ses inÎR fi t ter opt ics can be divided in to two groups:intrinsic losses and extrinsic losses. The extrinsic losses can alsobe divided into two groups: mechanical causes such as air bubbles,structural inhomogenities', imprefection in the core pattern and clad-ding surface which can be removed by precisely controlling and improv-ing the drawing technique. The second cause for extrinsic losses isprimarily associated with impurity absorption, such as oxide absorp-tion. For example, the absorption of As 0, at 81C, 934, 980, 1100,1165 cm" , and Si-0 band in 900 c m , or the aBsorption of H Se at 2200c m , and at 3600 cm" by OH. The absorption due to Ee - Se bondingthat usually appears in the 800 cm" area occurs because of a deviationin stoichiometry. The impurities of carbon and arsenic cause broad ab-sorption throughout almost all of the IR region. These extrinsic ab-sorptions can bs removed by using oxide catchers (such as urea)', byworking with very pure reagents and drawing fibers within a pure inertatmosphere. The most important factor i s a very pure reagent and inertatmosphere.

The intrinsic causes can be divided into three loss mechanisms.Two of them, absorption from electronic band edge ta i l and Rayleighscattering, are responsible for a decrease in optical absorption withincreasing wavelength. The third mechanism is multiphonon absorption,which causes increases in absorption with wavelength. This mechanismis the most important in the infrared region. The multiphonon absorp-tion decreases with increased atomic weight of the elements involved inthe absorption , thus justifying the usage of heavy elements.Therefore the multiphonon absorption moves to lower energy as the glasscomposition changes from Oxide to Sulphur, to Selenide or to Tellurite.

The recent interest in developing infrared optical waveguides hasbeen stimulated by advances in the use of C0? laser in surgery, cut-ting, welding, and heat treatment. The availability of such a wavegu-ide will play an important role in infrared image relay, remote meas-urements^ remote heating etc. Its theoretical high transparency mayalso serve for generating long distance communication l inks , althoughthis application is more a thing of the future.

329

From a practical point of view glasses rather than crystalline ma-terials are highly desirable for long optical fibers. Low toxic flour-ide glasses can be used in the near infrared region, while chalcogenideglasses are the most promising materials for a wider range of wavel-engths. Although much research on chalcogenide glass has been,, d,oge,only limited results are available concerning fiber fabrication.c>°'

The fibers we have been working with, are multimode fibers and arethick in relation to the wavelength. The main uses intended for thefibers are in power transmission, up to few watts in a single fiber,maintaining high energy density at the output port.

This paper describes techniques for preparation of low toxic chal-cogenide glasses and pulling fibers with precautions for eliminatingoxide impurities formation. We evaluate optical perfection of theseglasses and fibers, determine the maximum power that the existingfibers can withstand without being damaged by constant illumination(CW) and short pulses, and the angular scattering of the radiation atthe output port. We include a search for a possible dependence of theoutput angular scattering on the focal length of the lens focusing thelaser1, illumination at the fiber's end, and the fibers length.

Glass and Fiber Preparation; The chalcogenide glasses were preparedusing 99.9992 pure As Se or As, Se, elements. The reagents wereweighed and placed togetner in silica glass ampoules.

The ampoules were sealed after being evacuated for several hours.The sealed ampoules were heated to 900 C for 5-24 hours in a resistancefurnace. After the heating, the ampoules were withdrawn from the fur-nace and cooled in a i r . Glass rods 7mm in diameter and 30-150mm inlength were obtained depending on the reagent quantity.

We examined different methods for purifying samples from undesir-able oxygen. One of the methods that had been tried consisted of ad-ding CO(NH ) ? (urea) to the raw material in the evacuated ampoule. Theurea decomposes into CO, NH , N_ and H at 900 C. These gases removedall non-metallic oxides, sucn as Aso0o, Se0o, Sbo0, by reactions suchasSe

ASp0,+3C0 2As+3C0_; Fig. 1 compares the lnirared spectra of Asglass samples before and after treatment with urea. It i s very

clear from this figure that the urea removed the absorption peaks at„„„ __- __ J .^r-n __.-i r-_.. c o i a s e r nne the most important improve-

; lOcrn" peak.800 cm and 1050 cm . Forment was the removing of the

Two systems wereused for fibersdrawing;Crucible Method; Thissystem is intended forpulling an optical Cofiber out of a melt ~glass in a crucible. A ^glass ampoule contain- £ing the chalcogenite graw material was put inthe crucible in thefiber drawing machinewhich includes: A fur-nace with temperaturecontrol and a drun thatcollects the fibers.

MICROMETERS6

80

6 0 -

40z<£[E 20

IP II I2

C02 LASER

(70mm) WITH UREA NH4

(40mm)_ Y

As 40 6

1800 1600WAVENUMBERfcnrf1)

1400 1200 1000 800

330

The drum rotates on an axis and moves in a direction perpendicular tothe rotation, as i t s axis is threaded and is spun by a stable motor,whose speed is adjusted according to the desired fiber thickness.

The fiber is drawn from a nozzle which is at the bottom of thefused silica crucible. The crucible is set in the center of a tempera-ture controlled resistance furnace. The fiber diameter depends on theglass viscosity and drawing speed. The glass viscosity has a strongdependence on temperature and composition, so the furnace temperature•nust be very precisely controlled.The Rod Method; This system is intended for pulling an optical fiberDUt of a glass rod (preform). The system consists of a rod mount, aheating coil, collecting drum, and controller. The glass rod 6mm iniiameter and 2 cm long i s mounted in a clamp that can be lowered or ra-ised by a dc motor. The heating coil is fixed in i ts place and can ma-intain a fairly steady predetermined temperature which i s measured bythermocouple, and automatically controlled. The rod is lowered so that2mm of i t is inserted into the heating coi l .

Measurements and Comments; The optical losses through the glass fiberswere measured by the cut back technique in which the power transmittedby a fiber is recorded before and after a length of the fiber is re-noved.

A C0_ laser was used as the light source for the 10.6 m wavel-ength region while a pyroelectric IR detector was used for detectingthe light immediately behind the fiber. Precise positioning of thefiber with respect to the focused laser spot was accomplished by meansof an X-Y-Z translation stage, which allowed movement of the fiber enduntil a maximum power throughout was detected.

The fiber was exposed to a predetermined C0_ laser illuminationdensity. The i.iput power was calculated and the output power was meas-ured .

The laser 's power was raised until the fiber was damaged. Thiswas done with CW as well as pulsed radiation. The output scatteringpattern was measureda) Near field: qualitative impression could be achieved by placing aliquid crystal coated plastic sheet near a fiber end. The liquid crys-tal changes i ts color when its temperature changes, and a pinhole thindark spot appeared on the sheet.b) Far field scattering: The output end of the fiber was mounted at thecenter of a rotating table, on which an IR detector was installed, sothat the detector could 'move1 in an arc centered at the fiber 's outputend. The detector's window was 2.5 cm away from the fiber. Since nosharp angular dependence of radiation intensity was detected a largewindow of 10 (4 mm width) was used for measurement's convenience sake.

We have got extremely different results for different samples.Some fibers were not damaged at input power density even up to 10kW/cm . We estimate that the great differences are mainly due todifferences in raw material quality.

The angular dependence of the field intensity at the output ap-peared to be constant (within the limits of measurements accuracy),when the focal length of the lens or the fiber length were changed.These results were expected since:a) If the laser beam was a pure Gaussian, the beam's waist was for anylens in use a few mm. long, so that the input field was always planewave l ike .

331

b) For unclad fiber of high refractive index, no excessive attenuationof the higher order modes have been expected. These modes are respon-sible for the radiation in the wider angles. Most of the energy is ra-diated at angles 20°-i)0° measured from the f iber 's optical axis.

High energy density is detected right at the end of the fiber. Apiece of paper was burnt by radiation transmitted through a fiber, whenplaced close enough to i t s end. The highest output power density was0.8 kW cm .

Two types of experiments were performed in order to measure theattenuation due to bend in the fiber. The experimental arrangement wasset so that the only condition changed during the experiments, was thebend radius. All other conditions such as the inlet and the outlet ofthe fiber and the power of the laser were kept fixed. In the f irs t ex-perimental arrangement the fiber was bent around a cylinder to form a90° angle. Cylinders of various diameters were used down to 2 cm. Inthe second experimental arrangement, the fiber had one full winding.By increasing the distance between the inlet and outlet, the windingdiameter was changed down to 2 cm.

No significant influence of the bending on output power was deter-mined in the measurements resolution abili ty.

Conclusion: In this paper we have described the preparation of chalco-genide fiber in two different methods crucible method and rod method.Various measurements of infrared absorption power transmission outputangular dependence, and bend dependence were described. Attenuation inchalcogenide glasses, and some of their implications were discussed.The results indicate that the losses in the fiber tested are 0.1 dB/cm.The attenuation stays unchanged even when the fiber is bent to circularradius of 2 cm. We did not find angular dependence of the field inten-sity at the output when the focal length of the lens or the fiberlength were changed. We obtained fibers that were not damaged at powerdensity up to 10 kW/cm and which transmit power density of 0.8 kW/cm .

We did not detect any relation between the output power orscattering pattern and the fibers' radius of bending. Such relationwas not expected since any radius of curvature obtainable with glassfibers i s very large compared to the wavelength.

We have examined the existing fibers quality, by the various meas-urements described here. Great improvements are needed for makingthese fibers commercial.Acknowledgement This research was supported by a grant from the Nation-al Council for Research and Developnent, Israel.

References

1. A.L. Gentile et al in "Optical Properties of Highly TransparentSolids", Edit. S. S. Mitra and E. Bendow, (Plenum Press, N.Y.1975), p. 105.

2. J.A. Savage, P.J. Webber and A.M. Pi t t , IR Phys. 2Q_, 313(1980).

3. A. Bornstein, R. Reisfield, Non Crystalline Solids, 50, 23(1982).

4. A. Bornstein', N. Croitorii, E. Marom, Advances in IR fibers; LosAngeles. Technical Symposium by Society of Photo-Optical Instru-mental Engineering (SPIE) January 1982.

332

PROGRAMMING OF CRYSTAL DIAMETER IN CZOCHRALSKI GROWTHBY A COOLING PLOT: APPLICATION TO InSb

M. Azoulay and Z. Burshtein

Nuclear Research Centre - NegevP.O.B. 9001, 84 190 Beer-Sheva, Israel

The growth of single crystals by the Czochralski method is very com-mon for the growth of many types of crystals (1). Its main advantage isin preventing stresses associated with some other growth methods, such asthe Stockbarger (2) or Bridgman (3) methods. In the latters, the interac-tion of the crystal with the ampoule may result in enhanced striae andfracture. In the Czochralski method, an oriented single crystal seed isdipped in the melt, which is contained in a crucible. The seed is thengradually pulled out, usually also being rotated around its axis. The goalof rotation is mainly to improve the radial symmetry of the thermal pro-file.

In the following we assume radial symmetry to the furnace and pull-ing system. Consider a crystal of diameter rc, being pulled out of a cru-cible of diameter ro>rc. The pulling distance will be denoted by z, andthe melt height by hs. The values z=0 and hs=hi define the situation asthe pulling starts. The crucible bottom is at hs=0. In Fig. 1 we illust-rate a momentary situation during the pulling.

Let T c be the temperature at a point rc in the crystal circumference,which is in contact with the melt. Then, Tc = Tc(hs,rc,t). A full diffe-rentiation of Tc gives

CD

When the temperature in the furnace is controlled, say by a thermocouple,at a fixed point, the temperature at that very point, To, is a function oftime only. Therefore, by adding and subtracting 3TQ/"3t = dT0/dt, Eq. 1 maybe written as

dT dT 3T dh 3T drc _ o c s c c 3 , . , ..

dt~ " d ^ + a T dt~ 3r^ dt~ + 3t CTc " V * C2}

We assume now that thermal "stiffness" conditions may be realized in thegrowing cell, i.e. that temperature differences in the crucible vicinityare independent of the stage of the process. In other words, "stiffness"means, that any change in T o would change all the temperatures in thecrucible vicinity by exactly the same amount. Else, it means that thecrystal growth does not change the thermal profile. Later on we shallpoint out some measures that should help realize such conditions. Withthe above assumption, the fourth term on the right handside of Eq. 2 van-ishes.

dTc

dt

3Tc" 3h

dh

dt

3Tc3rc

drcdt

3Tc3t

333

From the balance of solidifying mass one easily obtains

dh* - dz

dF~ • • ~ i r 2 dt • l J

ro p l / p c - rc

where p^ and pc are the melt and crystal densities, respectively. For theisoment we assume uhat the crystal radius rc is kept constant, i.e.drc/dt=0. Under the above conditions, and inserting Eq. 3. into Eq. 2.,the latter reduces to

dT dT 3T r 2 (» ,c _ _ £ _ c . c . dz_ ,„

dt " dt " 3h 2 . 2 dt ' v J

s ro pl/pc " rc M

where rcC<») is the steady-state crystal radius.

In the course of growth of pure crystals from a stoichiometric meltthe temperature T c must remain constant. Therefore, dTc/dt=0. Then, weget

dT , 3T r 2 O 0o dz c c (dt " dt 8h 2 . 2, • ^}

5 ro pl/pc - rc M

The physical meaning of Eq. 5. is clear. The temperature T Q must fall ata rate that would compensate the increase in the melt surface temperatureas it drops towards the hotter regions in the furnace. Eq. 5. presents astraightforward recipe to present the crystal radius rr by a cooling plot.It requires a fair measurement of the longitudinal thermal gradient at themelt surface level 3Tc/3h. This may be done by measuring simultaneouslythe temperatures of two different points near the crucible, of differentheights, but of the same distance from the symmetry axis. It is interest-ing to note, that the radial thermal gradients play no role in Eq. 5. Theonly pertinent material properties are the melt-to-solid densities ratiopl / pc-

We have examined the foregoing idea in the growth of pure InSb singlecrystals. The crucible was made of silica, with three thermocouples atta-ched to its side, to measure the longitudinal temperature gradient. Theradial gradient has not been measured. The growing cell was a silica tube,with a resistivity heating element surrounding its central zone. A con-stant 30cc/min flow of hydrogen at normal pressure through the cell servedto eliminate oxides from the melt surface, as well as to improve the heatexchange. The seed was a square shaped rod oriented along the (211) crys-tallographic direction. In Fig. 2 we show a photograph of a crystal grownin our system following a cooling plot derived from Eq. 5. The growingconditions were: Liquid to solid densities ratio p^p = 1.14; Longi-tudinal temperature gradient 3Tc/3hs = -12.5°C/cm; Pulling rate dz/dt =

334

2.2cm/hr; Crucible radius rQ = 2.0cm. The cooling plot set in order toget rc0°) = 1.1cm was dT0/dt = 8.3°C/hr. No program was set for the broad-ening of the radius. As seen, the crystal radius grew gradually, reachinga final constant value of 1.0cm ± 5% (actually, (area/ir)'5, since thecross section is approximately trapezoidal). This value was kept constantover a length of 5 cm.

The above result is an example of the good fit that has beenobtained between the predicted radiuses and the experimental results. Par-ticularly, it proves that the temperatures, measured by the thermocouplesattached to the crucible side, represent the actual longitudinal tempera-ture profile in the melt. Still, some error might be involved by our simp-lifying assumptions.

FIG. 2

Photograph of InSb singlecrystal grown with a cooling

plot following Eq. 5.

FIG. 1

Illustration of momentary si-tuation of crystal and cruci-ble during pulling. The ther-mocouple at height h Q is usedto control the heating elementpower. Other thermocouples maybe distributed around the cru-cible to measure the thermalgradients in the course of

crystal growth.

335

Thus, has been shown that under conditions where the thermal profilein the crucible vicinity may be assumed to be rigid over a certain tempe-rature range, and independent of the growing stage, the radius of a crys-tal grown by the Czochralski method may be programmed by a simple coolingplot. The pertinent parameters are the longitudinal and radial thermalgradients, the crucible radius, and the melt-to-solid densities ratio. Acorrection, to account for the change in freezing temperature, is requiredfor the growing of highly doped crystals.

The following measures should help realize the required conditions:Reducing the thermal masses of the hot parts, increasing the thermal con-ductivities, generating high uniform temperature gradients near the cru-cible - negative in the longitudinal upward direction, and positive inthe radial direction, and pulling sufficiently slow. We have shown, thatour principle is practical in growing InSb crystals of diameters around2cm. It should presumably be the same for many other materials. The prin-ciple put forward might save the need for complicated diameter controlsystems, particularly in small-scale crystal growing systems for labora-tory use. We have not considered here some other important effects, suchas self and forced convections in the melt, entailing the thermal gra-dients and seed rotation speed (4,5). These might influence the crystalperfection, and should be considered independently.

REFERENCES

1. J.C.Brice, The Growth of Crystals from Liquids, Selected Topics inSolid-State Physics, Vol. 7, E.P. Wohlfarth, editor (North-Ho11and,1973) Ch. 7.

2. D.C.Stockbarger, Rev. Sci. Instr. 7_, 133(1936)3. P.W.Bridgman, Proc. Am. Acad. Arts Sci. 6£, 305(1925)4. J.R.Carruthers, J. Cryst. Growth, 36_, 212(1976)5. Y.Miyazawa, Y.Mori, S.Homma and K.Kitamura, Mat. Res. Bull., 13,

675(1978) ~

336

TOE DEPENDENCE OF THE HIGH TEMPERATURE,HIGH SOLAR FLUX STABILITY OF MATERIALSON SURFACE STRUCTURE AND COMPOSITION

A. Ignatiev

Department of Physics and ChemistryUniversity of HoustonHouston, Texas 77004

INTRODUCTIONThe stability of materials at elevated temperatures has long been a

point of great interest in materials research.(l>2) Until recently,however, such interest has centered principally on the stability of amaterial under high temperatures generated by the absorption of eitherinfrared or particle radiation. Such are not the conditions encounteredby materials used at elevated temperatures in solar apparatus. In thiscase the materials are exposed to concentrated solar radiation of wave-lengths from -0.3 ]im to - 2iim. Radiation from this solar spectrum is byand large absorbed over a much shallower depth (up to a factor of 100 to1000 shallower!-*)) than infrared radiation and hence this surface loca-lized absorption can lead to a surface temperature much higher than themeasured bulk temperature. Such an effect may prove deleterious onseveral fronts. In addition, the ultraviolet and near ultraviolet com-ponents of the solar radiation may indues various chemical changes at thesurface which may prove to be either beneficial or deliterious in termsof affecting the long term stability of the material in a solar enviro-ment.

It is clearly an enhancement of long term stability under elevatedtemperatures in solar environments that is sought after in solarmaterials. The work described below will show that such an enhancementcan become a possibility when the basic mechanisms responsible for thenoted photoeffects are identified and characterized. It is witn thatknowledge as an input that efforts can be made to generate new solarstable materials.

PHOTOEFFECTS IN SOLAR MATERIAL ABSORBER COATINGSWe have discussed in the past specific photoi iduced effects active in

the black chromev4-

5)

a nd black cobalt(

6) and wish to note here some of

the details of the studies as background for our current investigation utphotoeffects in metals and ceramics. For both electroplated black chromeand black cobalt (formed by oxidation of plated cobalt metal(7)) it wasobserved that the reflectance of the samples changed less from the "asprepared" case upon solar heating than upon infrared (oven) neating. Tnedifferences were in fact quite significant (Figure 1) e.g., for blackcobalt under oven heat at 460°C in air for 50 hours there was a resultantdecrease in solar absorptance Aα » -0.11, whereas for an identical solarheated sample under equivalent (bulk) temperature conditions, the solarabsorptance decreased only uy Aα « -0.S.

337

100

»S PfSEP.

OVEN HEM

03 u U IB SIWtVElEMTH (pa)

Figure 1. Spectral hemispherical reflectance for black cobalt samplesthat were solar simulator irradiated in air at 6UQ fcH/m? and 46QUC for50 hours, oven heated in air at 160"C fcr 50 hours ana for and "asprepared" sample.

The basis for such a photoeffect was investigated by applying thesurface sensitive techniques of Auger electron spectroscopy (AES), x-rayphotoelectron spectroscopy (XPS) and mass spectrometry. It was found inQcth cases that the effect was jjhotodesorption of oxygen bearing species(CO2) from the surfaces of the absorber coatings as a result of the nearultraviolet component of the solar radiation. This photodesorptionreduced the nunber of oxygen species at a surface available for oxidationand hence reduced the rate of oxidation - oxidation being the principaldegradation mode in black cnrome and black cobalt. It is seen here thatthe photoeffect in these two materials has been quite beneficial and shouldtherefore be utilized to its fullest extent when designing solar absorberappa.-atus. Other materials are also expected to be used in solar tech-nology, for reflectors, insulators and structural materials and it is ofimportance that active photoeffects be characterized in the systems.

METALSA large variety of metals are expected to be utilized in the con-

centrated solar enviroment. These include iron, chromium, aluminum,copper, nickel and various steels. It is of importance, therefore, todefine and characterize any photoeffects that significantly affect thestability in these systems. This nas been done to date for chromiun,iron and aluminum samples.

Chromium samples exposed to ~700 kW/m2 solar simulated radiationexhibited two modes of behavior depending on cleaning history of thesample surface. Samples cleaned in ultrahigh vacuum by argon ion bom-

338

bardment such that the surface was atomically clean showed no photoef-fects under solar i r rad ia t ion . However , samples not cleaned, i .e .retaining the native oxide and other surface impurities., showed markedphotodesorption of CO2 upon solar i r rad ia t ion (Figure 2). The desorptioneff ic iency was measured by mass spectrometry to he K3 X 10-6molecules/photon under i r radiat ion at 3I100A and varied l inear ly with pho-ton f lux . The wavelength threshold for the photodesorption was near5000A (~2.5eV) and thus the effect did not correspond to any known photo-induced effect in metals. The threshold behavior d id , however, nearlycorrespond to that expected for photodesorption from C r ? ^ . There arepresently discussions underway which address the discrepancy in themeasured and expected photodesorsption thresholds for l>2^3. (2.5 eV vs3.4 ev)(a>9) with the basis for the discrepancy lying in the mostappropriate description of the basic mechanism responsible for thedesorption (10,11)^ j ^ -jS f a i r to say, however, that a photo desurptionroechamism(s) is (are) active in Cr2(J3 and similar to that previouslynoted for the black chrome absorber coatings i t reduces s igni f icant ly theoxidation rate of the chromium.

SOUR HEAT

I OVEN HEAT

A . .Ji - 12 II 21 44

MASS NUMBER

Figure 2 . Hass spectra of solar-simulator irradiated, air exposedchromium (660 kW/m2 at 410°C in 5x10-'° torr) and infrared heated airexposed cnromium (410°C). Note the large increase in the 44 ainu (CO2) peakunder solar simulated irradiation.

Iron samples exposed to solar simulated radiation showed no photode-sorption effects when the surface was atonically clean. They did,however, exhibit photodesorption of CU^ <*t an efficiency almost twoorders ot magnitude lower than that of"chromium for a native oxide coatadsurface. On the other hand, iron samples exposed in air to solar simu-lated radiation at elevated temperatures showed a new photoefect. Thesamples exposed at 6U0 to 700 kW/rn- fluxes at 410°C, 5O5°C all showedincreased rates of oxidation under solar irradiation as compared to ovenheating in air. Figure 3 shows AES depth profiles (obtained by bom-barding with argon ions) of a solar simulator irradiated iron sample andequivalent sample heated in an oven. It is clear that the solar irradiatedsample has an oxide coating that is "-30^ tnicker than that of trie oven

339

Heated one, i .e. i t has undergone enhanced oxidation under the influence ofthe solar simulated radiat ion. The basis for this is believed to beenhanced dissociation of molecular oxygen at the iron sur-fdce by the solarphotons thus generating more highly reactive oxygen atoms for part icipationin the oxidation process. We have then, in the case of iron a detrimentalef fect due to solar i r rad iat ion of a material and must c lear ly include thisp o s s i b i l i t y in the u t i l i z a t i o n of iron in concentrated solar environments.

- — 5 = - ^ -

Fl

- . C _

OVEN HEAT

S O U R MEJITwD I M

HUTltlK tlit (all)

rtgure 3. Auyer electron spectroscopy/ion bubardment depth profilesfor oven heated and solar simulator heated (in dir) iron samples (110DC. 30minutes, 660 kH/m2). The Sputtering Tune can be converted to depth intotne materials by the factor of -50 ft/minutes. Note the thicker oxide layer(oxygen signal extending to -23 minutes) for the solar heated sample ascompared to the oven heated sample.

I rradiat ion of aluminum by concentrated solar simulated radiationresults in dif ferent effects in two specif ic temperature regimes, below40n°C and above 400°C with e flux dependence also observed above 400°C.The AES depth prof i le of an aluminum sample solar i rradiated for 2 hoursat 520°C in air at ~ 1.3 MU/m2 is shown in Figure 4a. For comparison, qdepth prof i le of a sample solar i rradiated at 520°C in a i r at ~ 25OkW/m2 isshown in Figure 4b and d depth prof i le for a sample heated in an air ovenat 520°^ is shown in Figure 4c. Several points can be made in the com-par is icn.

1) The high f lux solar heated sample has a thinner oxide (- 250A) asdetermined from the ion etch rate during p r o f i l i n g , than the lowflux solar heated sample (-3S0A) or the oven heated sample (-350A).

2) The stoichioinetry of the high f lux sample exhibits a higher oxygento dluuinun r a t i o in the surface region than does the low fluxsample. However, the stoichioTiutry is quite similar to that ofthe oven heated sample. Tiie stoicniometries of the oven heatedand hign f lux heated samples as obtained from the figures areapproximately: oven - Mβ^; high f lux - A J ^ O S ^ . Tiiesestoichiometrics are obtained from the atomic concentrations in tnedepth prof i les as determined from the ALS peak-to-peak heights andAES sensi t iv i ty f a c t o r s ! 1 2 ) . The sens i t i v i ty factors have ~20%uncertainty in them and are dependent on apparatus used.Therefore, i t is appropriate to denote (as is expected) the sur-face region of the oven heated alonintin sample as f u l l y oxidizied

340

The surface region of the high flux heated sample is thenalso AA2O3, however the low flux heated sample is noticably dif-ferent at - AJIO, i.e. denoting a not fully oxidized surfaceregion.

3) The top few 10's of Angstroms of the high flux heatad sample aremuch reduced in oxygen concentration with most of the aluminumreduced to the metallic state.

The above differences, although detailed only for samples treated at~520°C are consistently observed for samples treated at 470°C, 4()0°C and430°C.

: L

M

V/\

n

Al

0

Spuitf

—

HIGH FLUX- 1.3 mM

52Q"C

—

Law Fmx-2a) KX/M'520"C

5yut(er Tine (

g

fl)

\ /\ /

/ \7 V

^y \fl "^

/

riVLN

520"C

Figure 4. AE5 depth profiles giving the atonic concentration ofoxygen and aluminum for aluminum samples heated under different conditions.

I rradiat ion of aluminum samples below ~4(J0°C results in very l i t t l edifference between o.'en heated, high f lux solar heated and low f lux solarheated samples. The stoichiometrics are a l l near AJI2U3 and thethickness are a l l comparable and quite th in (<50A).

The mechanisms responsible for the noted behavior of aluminum exposedto concentrated solar simulated radiation are not yet f u l l y understood,however, several points can be made. i)The observed thinner oxide underthe high f lux i r rad ia t ion probably indicates a surface temperature m.iciihigher than that measured by the thermocouple. Less oxide growth his

341

been observed in aluminum above 600°C.U3) ii)The reduced amount of oxy-gen within the very top layers of the high flux sample indicates the pre-sence of an active photo-induced effect which reduces the oxide at thesurface of the sample. Such an effect may be photodesorption of oxygenbearing species, and this can be and is currently being tested for by tneirradiation of an oxidized sample in vacuum while monitoring desorbedspecies by mass spectrometry.

The lower oxygen to aluminum ratio in the low flux heated sample ascompared to the high flux or oven heated samples may have its basis inlower thermal gradients in the sample due to the low flux condition, and areduced, but still active photodesorption effect. The lower thermal gra-dient generating sample temperature uniformity approaching ttiat of tneoven neated sample, result in oxidation of the sample to a depth equiva-lent to the observed in the oven heated sample. The reduced oxygen toaluminum ratio in the surface region.however, as well as some ennancedoxide reduction in the very top surface of the sample is probably aresult of photodesorption of oxygen beaming species under the low fluxirradiation conditions. The decrease in reduction of the oxide at thevery top surface in the low flux sample as compared to the high flux sampledirectly indicates the effect of the magnitude of the solar flux on thesamples. Quantitative measurements are now underway to describe the effectmore fully.

CERAMICSIn working with the aluminum system we nave begun to grow thick

(3-5ym) thermal /U2U3 layers on aluminum for the initial evaluation ofspecific photoeffects in ceramics. Preliminary measurements show twoeffects: i) photodesorption of oxygen bearing species for aluminum oxidesolar irradiated at a flux of 1.5MU/m2 anc| -d^ a temperature of approxima-tely 400°C in vacuum; ii) discoloration of tiie coating - darkening, under~1.5 flW/m2 in air at ~400± 1UU°C. Optical measurements to define changesin reflectance are to be undertaken, as ar^ high resolution (lym spatialresolution) AES measurements to define the basis for the discoloration.The discoloration (darkening) clearly increases the solar absorptance ofthe aluminum oxide sample and could result in locally high energy absorsp-tion leading to the cracking and melting previously observed at Gl1 andSANDIA.ll4).

CONCLUSIONIt has now been clearly shown that a nunber of solars relevant

iriaterials exhibit strong photo-induced effects which may be deleteriousor beneficial. It is such effects that nust be identified, charac-terized, and understood with respect to the basic mechanisms responsiblefor them. With this knowledge it is then possible to design solar systemcomponents that are highly stable under concentrated solar environments.

ACKNOWLEDGEMENTSThe assistance of A. Mesawri and A. Zomorrodian is greatly

acknowledged. Support for this work has been provided by the Departmentof Energy - STAKC and by the University of Houston Energy Laboratory.

342

REFERENCES

1 High Temperature OxidatJ_on_- j tes i stajnt _Coati_ngs_, e d . , National•Res"e¥rcH"Co"uncfl" T l a t . " Head""of S c i . , (T97tT).

2. I .E . Campbell and E.M. Sherwood, High Temperature Mater ia ls andTechnology, (Wi ley, New York, 19677". "" '

3. J .P . Jackson, Class ica l Electrodynamics, p. 225 (Wi ley, New York 1%54. G.B. SMith, G. Zajac and" A." Ignat ieV, Solar Energy 2$_, 279 (1982).5. A. Igna t iev , G. Zajac and G.b. Smith, Proc. SPIt 3?A, 170 (1982).6. A. Igna t iev , Yearly Prog. Report, DOE - STARC, Task"? (Houston, 1982)7. G.B. Smith, A. I yna t iev and G. Zajac, J . Appl. Phys. _5J., 4136 (1980).8. L. Korenbl i t and A. Igna t iev , Surf . .Sci (1984).9. 6.W. Fabel, S.M. Cox and 0. Lichtman, Sur f . S c i . , £ 0 , 571 (1973).10. P. Mark, R.C.A. Review 2^ , 461 (1965).11 . 0. Lichtman and Y. Shapird, in Diemistry and_ Phys_i_cs of So_l_ids_-_H,

e d . , R.Vanselow p. 397 (CRC PreTsTT9T8)T12. Handbook_ôf_ Auger_^p_ectros_copy_, ed. L.E. Oavis e t . a l . (Physical

Tnect romcsr~f f inn . ,1976)."13. P.Doherty and R. Davis, J . Appl. Phys. 34, 619 (1963).14. High Tenp. Mater ia ls and Ther.nal Science Coord. Meeting, (SEKI,

Golden 1983).

343

TRANSPARENT CONDUCTOR FILMS AS A MATERIALFOR PHOTOVGLTAIC JUNCTIONS WITH POLYCRYSTALLINE SILICON

Z. Harzion, M. Zafrir, J. Rishpon*, S. "ottesfeld* and N. Croitoru

Tel-Aviv University, Faculty of Engineering, Dept. of Electron DevicesTel-Aviv, 69978, Israel

*Dept. of Chemistry

INTRODUCTION

In recent years, a great effort has been expended to find a materialwhich might replace single crystal Si for the manufacture of terres-trial solar cells. Polycrystalline silicon [poly-Si) is one of thecandidates which might yield inexpensive solar cells. Such cells aremade of p type poly-Si as the base converting semiconductor, and atransparent conductive oxide (e.g. ITO) as the other element of thejunction. Tcansparent conductors are degenerated wide bandgapmaterials, which are used as an antireflecting coating and window aswell as active layer and upper contact in surface barrier devices. Thepoly-Si consists of small crystallites, and therefore, grain boundaries(GB) appear to be the main reason for the low conversion efficiency andthe relatively short term stability of these cells compared to thesingle crystal silicon solar cells. Much work has been done on thetheoretical modeling of GB effects in poly-Si (1,2). It has beendemonstrated that cell performance is mainly controlled by intragraindefect density and less by the grain size. GB are expected to containhigh density of interfacial states which act as minority carrierrecombinations sites, and produce a potential barrier against thecarriers transport. Increase of the efficiency will only be possibleif the limitation of GB recombination can be overcome (e.g. by passiv-ation) . The efficiency of solar cells depends largely on the life-timeof the carriers generated by the incident light. One of the methods todetermine life-time is to study the response of the solar cells tolight pulses. The method is usually applied to a p-n junction in asingle crystal device. We have recently extended this method totransient measurements of photovoltage decay (PVD), photocurrent andphotoluminescence in photoelectrochemical (PEC) solar cells (3,4). Wehave now applied this method also to poly-Si solar cells. The PVD aswell as the impedance in a wide range of frequencies (0.1 Hz - 1 MHz)were measured and compared with those obtained for single crystal solarcells.

EXPERIMENTAL

The preparation of ITO solar cells proceeded as follows: a) etch-ing for 8 hours in KOH-methanol solution; b) deposition of Al film of2000A and diffusion for back contact; c) sputter etching and ITOdeposition by sputtering in MRC 820 system. Some solar cells werepassivated with iodine saturated methanol solution before the ITOsputtering. I - V characteristic of the cells was recorded. Thestructure of the solar cell is shown in Fig. 1.

PVD measurements were done by applying a short light pulse (10 nsec)on the cell. High impedance probe was connected across the cell in the

344

open circuit photovoltage mode.Impedance measurements of thecells were done in the range of0.1 - 500 Hz (5). More detaileddescription of the PVD experi-mental system is given else-where (4) .

RESULTSFrom the I-V characteris-

tics of ITO/poly-Si solar cellsunder illumination, Voc wasdetermined as 0.35-0.38 V,I *> 2mA/cm and the FF was luw(5?35). Passivation v.lth iodineincreased Voc by50 mV and doubled I

The PVD measure-ments of poly-Si andsingle crystal ~,olarcells can be summar-ized as follows:1) the V at the

ochighest pulse, inten-sity (2pJ/cmA) forthe poly-Si solarcells was between300 and 410 mV. Forthe single crystalsolar cell, highervalues were observed.(500 mV); 2) thePVD curve has a non-exponential shapeand lasts for 1-30msec; 3) PVD ofpoly-Si solar cellreduces to half ofits initial valuewithin 10 usec,while for singlecrystal, it hardlydecays within thesame time window(Fig. 2b); 4) Atthe highest intens-ity of the incidentlight, 20% of theinitial photovoltageof the poly-Si solarcell decays at arate of 18-60 mV/ysec, while forsingle crystal solar

I.

.3 -

xoE .4

.2

I.-AI diffused contact2L-Silso polysilicon3.-2TO film (Schottky contact)4.-AI diffused, top contact

Fig. 1ITO/poly-Si solar cell structure

(a)

SINGLE CRYSTAL

POLYCRYSTALLINEf- —i 1- 1 J i -

0.5

Tmsec

1.0

(b)

1.0

.8

.6

.4

.2

s.

POLYCRYSTALLINE

T IO6secFig. 2a. PVD curve of ITO/Si junctionb . In i t i a l 10 psec region of the

8

PVD

10

100

1

—

ri •

55,3 .

V

34

1 ,

1 ' 1 ' 1.^25.6

1 , 1 , 1

1 1 '

—

% -

345

cell, the rate is6 mV/ysec. To gcharacterize the fshunting of the ITO/ 5poly-Si junction, theimpedance of the solarcell was measured. A jjjtypical Cole-Colecurve at zero exter-

Tig.T Theiwpo^y o" " 50 ' 100 150 " 200" 250 " 300-Si solar cell has RE Z Kohmgood shunting proper- F- 3

ties. The different- Cole/Cole plot of ITO/poly-Si solar cell;xal resistance is ,~ , , 7, ___ ,,_ , , some frequencies are marked,about250 Kfi and does ^not limit cell performance.The cell exhibited parallel R-C behaviour (almost ideal semi-circle inthe range 0.1-500 Hz)-The series resistance was found in the range5O-2OOS2.-

DISCUSSION

Some analytical solutions of PVD curves for single crystal junctionhave been published (6-9). For a poly-Si based junction, grain boundaryeffects should be considered. Even for a junction in the singlecrystal, the analytical expression of PVD does not take into account allthe boundary conditions due to the complexity of the continuity equation.By assuming a finite thickness of the diode (6-7) , the effect ofrecombination at the back contact and at the emitter is important. AnRC influence on the time constant of the cell was also found and shouldbe considered (8,9). In our measur.e-.:i3Pts the slow region ('vL-lO msec)of the PVD curve is controlled by the RC time constant of the cell,in accordance with the literature (9). The effect of light intensityon the effective diffusion length (Le£f) must also be considered.In poly-Si, Leff increases with light intensity either by collapseof the space charge region caused by existence of intergrain clusters(10)or by trap saturation of deep trap levol (It). In spite of the highlight intensity used in our measurements, the inixial PVD rate in theITO/poly-Si solar cell is very fast (18-60 mV/ysec) compared to asingle crystal Si/ITO junction (6 mV/ysec) and Si p-n junction(0.2-2 mV/ysec) (9). This means that the effective life-time (xeff)and Lef£ in poly-Si are shorter than in single crystal Si, thusreducing the collection efficiency of the photogenerated carriers, asseen from the photocurrent. The relatively high series resistance(50-200J2) also contributes to the low Isc. Part of the series resist-ance is due to the resistance of the ITO film itself, which has to beimproved.

CONCLUSION

The PVD measurements give information about the decay rate of thephotocarriers while the Cole-Cole curve measures the impedance of thecells. These are two of the main factors which determine the solar cellefficiency. Passivation with iodine improves the efficiency. PVD andCole-Cole measurements on the passivated solar cells will be performedand compared with those obtained for non-passivated solar cells.

346

ACKNOWLEDGEMENT

This research was supported by a grant from the National Councilfor Research and Development, Israel and K.F.A. Julich, Germany. Theauthors wish also to thank M. Evenor for helping in the computerprogramming.

REFERENCES

1. H.C. Card, J.G. Show, G.C. McGonigl, D.J. Thomson, A.W. deGroot andK.C. Kao; Proc. of the 16th IEEE Photovolt. Spec. Conf; San Diego,Calif.(1982) 633.

2. N.C. Haider; Proc. of the 16th IEEE Photovolt. Spec. Conf; SanDiego, Calif.(1982) 640.

3. 1. Harzion, D. Huppert, S. Gottesfeld and N. Croitoru; J. ofElectroanal. Chem. 15£(1983) 571.

4. Z. Harzion; Ph.D. Thesis, Tel-Aviv University(1983).

5. J. Rishpon and S. Gottesfeld; J. of Electrochem. Soc; Submitted.

6. 0. Von Ross; J. Appl. Phys. 52_(1981) 5833.

7. S.C. Jain and V.C. Ray; J. Appl. Phys. 54Q983) 2079.

8. A.R. Moore; RCA Rev. 40_(1980) 549.

9. J.E. Mahan and D.L. Barnes; Solid State Electron. 24_(1981) 989.

10. S.K. Agarwal, Harsh and S.C. Jain, P. De Pauw, R. Mertens andR. Van Overstraeten; Proc. of the 16th IEEE Photovolt. Spec. Conf;San Diego, Calif;(1982) 366.

11. C.T. Ho, R.O. Bell and F.V. Wald; Appl. Phys. Lett. 31(1977) 463.

347

PHOTOELECTROCHEMICAL CHARACTERIZATION OF•J <J

Geula Dagan, Gary Hodes- Saburo Endo,* and David Cahen

The Weizmann Institute of Science, Rehovot 76100, Israel*Dept. of Electr. Engin., Science Univ. of Tokyo, Tokyo 162, Japan

Ternary Cu-ln-chalcogenides are of considerable current interest assemiconductor materials because of the impressive performance of photovol-taic devices containing CuInSe2 (1). As part of a program to investigatethe series (Cu,X) (In2X3)j_x [X=S,Se; 0<x<l], we studied CuIn^Sg (X=S;x=0.17), the only true compound in the sulfide series, besides CulnS- (2).Because of the respectable photovoltaic performance of both CuInS^ andCuInSe2 as photoanodes in a photoelectrochemical cell, where aqueous poly-sulf ide is the redox electrolyte (3), we used CuInrSg as photoanode in sucha configuration. In this way we hoped to pbtain information on the effectsof 1) decreased Cu-content and 2) non-tetr'ahedral coordination (CuIncSg hasa spinel structure, rather than the chalcopyrite one) on the photovoltaicperformance and stability of Cu-ln-chalcogenides.

The outstanding stability of the dichalcogenides can be ascribed to theCu d-orbital participation in the valence band, something that can decreaseCu-X bond weakening upon excitation, and/or to the presence of a thinindium oxide layer on the surface, which might act as a protective barrierto photocorrosion. A further impetus for this study was the fact that noanalogous "CuIn^Seg" compound exists, but only a mixed cubic/hexagonalpolytype with a stoichiometry close to it. (4)

RESULTSThe material we studied was invariably n-type, both as single crystal

and as slurry-painted film. (This latter type of electrode is not furtherconsidered here.)Doping: Some photovoltaic response was obtained after heating the ini-tially highly resistive crystals in a stream of H 2 at 500-550°C for ca. 7hrs, or in low vacuum at ca. 200°C. Optimal results were found for crystalsthat were heated in evacuated sealed ampoules for ca. 2 days at ca. 400°C.In this way resistivities dropped by 4 orders of magnitude to ca. 0.25ohm-cm. Li-doping, using LiCl, also gave reasonably photoactive material.Surface treatments: Acid(1:3 aqua regia:H2O) etch, followed by H20 aqueousKCN H2O rinses, improved the photovoltaic response of as-doped crystals.In photoelectrochemical cells, using aqueous polysulfide electrolyte, thesecrystals gave short circuit currents up to 4 mA/cm and photovoltages ofca. 500 mV. Significantly better results were obtained after photoelectro-chsmical etching in tenfold diluted aqua regia at 2 V (reverse) bias vs.the etching solution potential. In this way short circuit currents above10 mA/cm were obtained, but the photovol^c,as were slightly lower (ca. 400mV).

348

Photoelectrochemical Characterization: When aqueous ferro/ferricyanide,rather than polysulfide was used as electrolyte, improved fill factors(0.5, rather than 0.3) and photovoltages up to 450 mV (for a photoetchedelectrode) were obtained, but the photoanode was not stable in this elec-trolyte.

In polysulfide, flat-band potentials from -440 to -560 mV vs. the poly-sulfide solution potential (ca -750 mV vs. SCE) were measured by capaci-tance voltage (Moct-Schottky) plots. Freshly etched crystals gave highervalues (-1.5 V vs. SCE) a? did measurements of the onset of photocurrent(-1.35 V vs. SCE) (3). These values are not too different from thoseobtained for n-CuInS-p) (-1.5 V vs. SCE). This similarity may indicatesimilar, solution-induced, shifts of the band edges that could be due tospecific adsorption of oligosulfides. The improved photo-I-V characteris-tics in ferro/ferricyanide point to kinetic limitations in polysulfide. Itis possible that these are connected with such surface adsorption.

Contrary to what is found for n-CuInS^ in polysulfide, the n-CuIncSsystem shows a slightly negative temperature dependence. This probaBlmeans that the (solid state) photovoltaic temperature dependence here isstronger than that of tha solution (including the ad- and desorption equi-libria) .

The poor long wavelength response of single crystal electrodes (whr-ncompared with that of polycrystalline thin films), can indicate that thephotovoltaic quality of the crystals is still far from optimal. From spec-tral response measurements (in aqueous sulfide) an indirect allowed tran-sition of 1.35 eV is derived. No clear evidence for a direct allowed tran-sition was found, contrary to solid state measurements on similar crystals(5).

The output stability of the CuIncSg/polysulfide system (at maximumpower) was found to be inferior to that of the analogous CuInS^ system.After an initial two-fold increase, until some 5 KC/cm photocnar . waspassed, a gradual decrease set in. In the best case this decrease was soslow that even after 17 KC/cm photocharge passed the final output wasstill above the initial one. After passage of this amount of photocurrentthe photo-I-V showed a decrease in open circuit voltage but an increase inshort-circuit current. The (poor) fill factor hardly changed. It istherefore reasonable to assume that the initial power increase is due todecreased recombination losses, possibly due to improved transfer kineticsacross the semiconductor/electrolyte interface. This can be brought aboutby "auto-photoetching".

Surface analyses show that the near-surface becomes Cu-depleted (as isobserved for the analogous CuInSe2 system). This is especially pronouncedafter photoelectrochemical etching. Then the surface seems to becomemainly indium sulfide. After a stability test sulfur/oxygen exchange isseen very near to the surface, indicating the presence of some indiumoxide. The above-mentioned results from capacitance-voltage measurementsalso indicate the occurence of some surface changes, in this case, veryshortly after imnersion in aqueous sulfide. These changes are also evidentin the large increase in spectral response, during use in this electrolyte,and in the decreased hysteresis of the capacitance-voltage plots, aftersome use in this electrolyte. The very high "donor-density", as calculatedfrom all these plots, indicate a near-degenerate surface. This could bodue to a highly doped oxide or oxysulfide top layer.

349

It is quite likely that, in contrast to CuInX-, the already Cu-poor spi-nel cannot tolerate further decrease of its Cu-content, without deleteriousconsequences. The decreased open-circuit voltage indicates that the poten-tial barrier near the interface decreases* The fact that the presence ofindium oxide does not prevent the decrease in output stability, suggeststhat the Cu-content of the material is to blame. It is thus likely thatthe Cu-d-orbital participation in the valence-band of the dichalcogenidesis the factor that is mainly responsible for their excellent output stabil-ity, although it should be borne in mind that the oxide film in the spinelis considerably thinner than on the disulfide. Thus, while the presentresults make this a plausible hypothesis, more experiments are needed toprove it.

REFERENCES1. R,A. Mickelsen and W.S. Chen, Proc. 16th IEEE Photovolt. Spec. Conf.

(IEEE NY 1983) pp. 781-785.2. J.J.M. Binsma, "Crystal Growth and Defect Chemistry of CuInS2", Ph.D.

thesis, Nymegen, The Netherlands (1981).3. Y. Mirovsky, Ph.D. thesis, Feinberg Grad. School, The Weizmann Institute

of Science, Rehovot, Israel (1983); Chs. 4,5; Y. Mirovsky and D. Cahen,Appl. Phys. Lett. 40, 727 (1982).

4. J.C.W. Folmer et al. J. Electrochem. Soc. 130, 442 (1983), RNP511 and tobe published.

5. A. Usujima, S. Takeuchi, S. Endo, T. Irie, Jpn. J. Appl. Phys. 20, L505(1981).

ACKNOWLE DGEMENTSWe thank the German Federal Ministry for Research and Technology (BMFT) forfinancial support through the Nuclear Research Centre KFA, JUlich and theIsrael National Council for Research and Development, and John A. Turner(Solar Energy Research Institute, Golden, CO, USA) for capacitance-voltagemeasurements and their interpretation.

350

MNi Fe - HYDRIDE COMPACTS FOR HYDROGEN HEAT PUMPS4.15 0.85

Y. Josephy, Y. Eisenberg and M. RonDepartment of Materials Engineering

Technion, Israel Institute of TechnologyHaifa, Israel 32000

Hydrogen heat pumps (h.h.p.)» an alternative energy conversion means oreven energy source, are an important and developing part of the futurehydrogen technology.

The main characteristics required from hydrides suitable for h.h.p. arelarge hydrogen flow rates and concomitant thermal effects. The thermaleffects are produced by the enthalpy of the hydrogen absorption/desorption reaction.

As a result of the hydriding reaction, the thermal conductivity of mostuseful hydrides decreases by orders of magnitude. Rechargeable hydridesare known to undergo comminution upon repeated cycling, ultimatelyturning into micron-sized particles. The fine particle size of thepowder facilitates the chemical reaction rate by providing a largesurface area, but it further decreases the thermal conductivity of thehydride. The resulting powder mass has a thermal conductivity of theorder of 1 w/m°C or less, which, in the materials meant to be utilizedin h.'i.p., does not provide a sufficient heat transfer rate. Inparticular, the point of lowest temperature of the h.h.p. represents abottle-neck for the heat transfer.

Presented here are hydride compacts that have a high thermal conductivityand provide large heat transfer rates. Such compacts, called p.m.h.(abbreviation of porous metal-matrix hydrides) are being developed inour laboratory. P.m.h. compacts have a complex microstructure consistingof a porous metallic matrix in which fine hydride particles are embedded.The metallic matrix material does not react with hydrogen but serves toconduct heat to and from the fine hydride particles (1,2,3) ., The micro-structure of a p.m.h. compact is seen in fig.l.

Fast hydriding/dehydriding kinetics are a necessary condition for ahydride suitable to be used in a hydrogen heat pump. The kinetics ofthe RNi5 compounds (where R is a rare earth), are known to be relativelyfast (4). In our experiments (unpublished), we found that the compoundMNi4.i5FeQ_85 is also endowed with relatively fast kinetics. Thecommercially available hydride MNi4.15Feo.85Hx was converted into p.m.h.compact by means of an aluminum matrix. A specially developed sinteringprocess, that makes the p.m.h. compacts stable during repeated cycling,has been described elsewhere (5,6.7).

The thermal yield of the p.m.h. compacts was evaluated under conditionssimilar, as far as possible, to the ones prevailing in the heat

Research sponsored by U.S.A. Israel Binational Science Foundation,Jerusalem, Israel.

351

Fig.l. Optical micrograph of a p.m.h.-compact ofMNi4#i5Feg.85 with 18 wt/o aluminum;Magnification: 1050.

exchanger of a hydrogen heat pump. A tube simulating a modular elementof a heat exchanger was filled with the above compacts and incorporatedinto a high pressure measuring system. A jacket, through which waterflows as heat transfer medium, was formed by a concentric tube. Thefollowing parameters were measured.

• hydrogen flow and hydrogen flow integral

• hydrogen pressure within the modular element and withinthe hydrogen line

• temperature of inlet and outlet water as well as twotemperatures at a known radial distance of the hydridecompacts.

Data were collected, recorded and processed by a microcomputer-basedsystem. Plots of the simultaneously recorded parameters in the process-es of hydrogen absorption and desorption are shown in Fig.2a, b andFig.3a, b.

Fig.2a and b show the instantaneous and the integral hydrogen flows andthe hydrogen pressures. The hydrogen flow is transient in character.Several values for the integral hydrogen flow are shown in Table 1, for60, 120 and 180 seconds, for both absorption and desorptxon.

352

so ,

40

30

10

a

90

BO

70

60

50

40

a: 30

20

la

0

me mss.IHfflESS.

IHIEOn. F19I

120 ISO 240 300 560 420

TIME. SECONDS

340 600 660

Fig.2. Hydrogen flow and pressure plotted vs. time, for amodular element filled with MNi^#isFeg_35HX p.m.h.-compact with a 18 wt/o aluminum matrix; the waterflow rate was 2 lit/min.

a. during the process of hydrogen absorptionb. during the process of hydrogen desorption

Table 1.

Process

Absorption

Desorption

Fraction of hydrogen sorbed

60 sec

0.86

0.57

120 sec

0.97

0.86

180 sec

0.99

0.96

Remarks

Fig. 2a

Fig. 2b

The mean water temperature excursion upon absorption and desorption didnot exceed 2°C, at a water flow rate of 2 lit/min. During absorption,the temperature of the hydride increases by 17°C within 15 seconds,while during desorption the temperature drops by 6.5°C, also within 15seconds.

353

SO j

in

so

20

10

a

a

Q] DEEP TDf* n>1

^ MIER IK IDT nti

Q wiot ajT l o r nti

Icr

TIME. SECONDS

Fig.3. Two temperatures, designated shallow and deep at aconstant radial distance of the MIH.4misFeo.85HX p.m.h.-compact with 18 wt/o aluminum matrix, plotted vs. time.Water inlet and outlet temperatures vs. time.The water flow rate was 2 lit/min.

The temperature of the hydride as a function of time during the desorp-tion process was used for the estimation of the thermal conductivity(see Fig.3b). At the point of the lowest temperature of 10°C, wheredT/dt - 0, quasi steady-state conditions were assumed. A heat transferequation including an internal heat generation term, qv, was implied (2).The term qv was derived from the measured hydrogen flow and thetemperature difference, AT, from the measured temperatures at a knownradial distance. As a result, a value of k = 21 W/m°C was obtained.

The fractions of desorbed hydrogen vs. time, for three types ofMNi4#i5FeQ#85^x hydrides are shown in Fig.4. The three types ofhydrides were: MNi4#i5Fe0#85

Hx hydride in powder form, p.m.h. with a27 wt/o aluminum matrix and p.m.h. with a 18 wt/o aluminum matrix.The p.m.h. compacts with a 18 wt/o aluminum matrix show the greatestfraction of desorbed hydrogen after given times of desorption. Forthese particular p.m.h. compacts, the thermal yields per kg of hydridewere calculated and marked on the upper curve of Fig.4. The thermalyields were calculated using the hydrogen flow, multiplied by theenthalpy of the reaction. The enthalpy of MNi/, p^5Feo. 85 hydrogenation isknown to be AH 'v. 1050 joules/liter H2 (2) . The maximum uptake was takenas 1.2 wt/o of hydrogen. The densities, p, and the calculated fractions

354

MNJ4B F»a85 H^

PMH-COMPACT WITH ISol/o ALUMINUM

PMH-CCMPACT WITH 27.1/0 ALUMINUM

• —POWDER

TIME(MIN)

Fig.4. Fraction of desorbed hydrogen (integral flow) vs. timefor MNi^j^FeQ 85^x ~ powder and p.m.h.-compacts with18 wt/o and 27 wt/o of aluminum matrix. Water at roomtemperature (between 15 and 20cC) and a flow of 2 lit/min.

of porosity of the two compacts were: 5 gr/cnP and 4.7 gr/cm3, and0.20 and 0.12 respectively, for the 18 and the 27 wt/o aluminum matrices.The porosity includes open-interconnected as well as isolated pores.Structural studies for establishing the distinction between open andclosed porosity are under way.

DISCUSSION

For a particular hydride, the thermal conductivity of p.m.h.-compactsincreases with increasing percentage of aluminum matrix. Similarresults were found in this work. However, the hydrogen yield and,consequently, the thermal power are higher for the 18 wt/o than for the27 wt/o aluminum matrix compact. This is so in spite of the fact thatthe thermal conductivity of tha 27 wt/o aluminum compact is higher thanthat of the 18 wt/o aluminum one.

The rationale explaining the apparent contradiction, that increasingthermal conductivity does not cause an increase in the hydrogen yieldis seen in the following.

The porosity, and in particular, the open-interconnected porosity, isan important factor in determining the permeability (or passability) of

355

hydrogen through the porous structure of the p.m.h.-compacts. Thevolume fraction, and the shape and size distribution of the open-interconnected pores, strongly influence the hydrogen flow kinetics.

The porosity was found to be higher for the 18 wt/o, than for the 27 wt/oaluminum compact and it may have the effect of facilitating the hydrogenflow, swamping the effect of the higher thermal conductivity of the27 wt/o aluminum. In order to evaluate these effects quantitatively, thevolume fraction, and the shape and size distributions of the open,interconnected porosity, are currently under intensive investigation.

In general, the higher hydrogen and thermal yields of the p.m.h. compacts,compared with powders are evident from the results discussed here. Thebeneficial effect on the hydrogen and thermal yields, of the increasedthermal conductivity of the p.m.h. compacts, is capable of further improve-ment by better understanding of the influence of various factors on thehydrogen flow.

REFERENCES

1. Ron, M., 11th IECEC, p.954 (1976).

2. Ron, M., Navon, U. and Levitas, I., Proc. Int. Symp. on Metal HydrogenSystems, Miami Beach, April 1981, Pergamon, Oxford, 1982, p.701.

3. a) Groll, A. and Nonemacher, A., 17th IECEC, 1982, p.1185.b) Rudman, P.S., Sandrock, G.D. and Goodell, P.D., J. Less-Common Metals,89_, 437 (1983).

4. Goodell, P.D. and Rudman, P.S., J. Less Comm. Met. 89_, 117-125 (1983).

5. Ron, M., Gruen, D., Mendelsohn, M. and Sheft, I., J. Less-CommonMetals, 74_» 445 (1980).

6. U.S. Patent No. 113873.

7. Israeli Patent application No. 66552. Patents applied for in sixother European countries as well.

356

CROSSLINK DENSITY OF POLYMERS - CAN IT BE DETERMINED

BY SOLVENT SWELLING?

Moshe Gottlieb

Chemical Engineering Department, Ben Gurion LI., Beer Sheva

ABSTRACT

The swelling of polymer networks is analyzed with emphasis on themain assumptions inherent to the classical development. Flory's newtheory of rubber elasticity is included. Comparison of experimental dataobtained by swelling the same network by two different solvents indicatelarge discrepancy between theory and experiment. The entire method isshown to be unreliable for exact determination of crosslink density.

INTRODUCTION

Determination of the degree of crosslinking of a polymer network bymeans of the amount of solvent uptaken at saturation has long been usedas a standard method (1, 2), the experimental simplicity of this tech-nique accounting for its widespread use. In order to obtain the concen-tration of crosslinks from the saturation polymer volume fraction threeassumptions are made: 1. a model for rubber elasticity is assumedusually "phantom" (3] or affine (4); 2. the mixing of polymer and solventis described by a thermodynamic model - customarily by the Flory-Hugginslattice model; 3. the chemical potential of the solvent is assumed to bedescribed by the sum of two contributions, one due to elastic deformationof the network and the other due to the mixing of the two species (thisis the so called Flory-Rehner [5) additive Gibbs free energy changeassumption).

Some doubts have been raised recently regarding the last assumption(6, 7). Also, it has long been known that both rubber elasticity theoriesmentioned above fail to describe experimentally observed stress-strainbehavior (8, 9). Hence, it seems appropriate to determine the validityof the swelling method for the determination of crosslink density.

THEORY

According to the Flory-Rehner assumption the chemical potential ofthe solvent relative to its reference value is given by

(yl " Pl 3 = (P1 " elasticity + ^1 " mixing W

If we assume that the Flory-Huggins lattice model is applicable

^1 ~ Vmixing = - V + V2 + XV2 <2>

where V2 is the volume fraction of polymer in the solution and X is

357

Flory's interaction parameter which is a function of V2. The elasticcontribution to eq. (1) is obtained from a rubber elasticity model by

where AAe£ is the free energy change due to elastic deformation of thenetwork, ni the number of moles of solvent and N the Avogadro number. Atequilibrium eq. (1) is set equal to zero and hence

N(8AAeJl/3n1) = -[lnCl - v2) + v2 + xv]] W

For phantom network (2, 8)

p h - v2) • v2 • xVjKvl/Vjll/SCl/Vj) (6)

where £ is the number of network cycles per unit volume, \\ is the molarvolume of solvent, v§ is the polymer volume fraction upon networkformation (crosslinking stage). The swelling equilibrium value of V2should be used in 2q. 6. The cycle rank ? is related to the concentra-tion of junctions p and junction functionality $ by y/£ = 2/<|> - 2. Therelationship between ? and V2 for a phantom network is indicated inFig. 1 for the case of a Polydimethyl siloxane (PDMS) network swollen inbenzene. The value of x used here is taken from Brotzinan and Eidlinger(10). A similar expression may be obtained for an affine network model

Unlike the phantom network, the last equation depends on junctionfunctionality. For a given swollen network [y-f) £af is always smallerthan Cph approaching it as <j> •+ °°. It should also be pointed out thatalthough fph and £af depend on the solvent (vj) and the polymer (x) theratio 5af/fph is independent of the polymer solvent system. The depen-dence of £af on V2 for PDMS/benzene is depicted in Fig. 1 for 3 and 4functional networks.

As already mentioned, the phantom and affine network theories areincapable of describing stress-strain behavior of real networks due tothe omission of intermolecular interactions. A new theory (8, 9) basedon the concept of suppression of junction fluctuation has been proposedand found capable of describing networks under different modes ofdeformation. The model has two parameters K and t, both of which dependon network topology only. According to theory polymer networks show abehavior which is intermediate to the phantom and affine models. Phantomnetwork corresponds to K = 0 and affine network toK + » also, phantomnetwork behavior is resumed at high deformations (or swelling). Thecycle rank for the Flory-Erman model is given by

eFE = ?ph[i + ( v / s W 1 (8)

358

where K (v ) is a function that depends on K and ? (cf. ref. 9). Eq. (8)is plotted in fig. 1 for several typical parameter values (9,10). Itshould be stressed again that the ratio ?FE/?ph is independent of thepolymer/solvent system for a given set of parameters,.


In order to test the swelling method a model network of known degreeof crosslinking has to be prepared. Despite the progress in the area ofmodel network formation, structure determination is still questionable.But even if this problem is resolved the question of network elasticitymodel exists. It is possible to resolve both problems simultaneously byswelling the same network by two different solvents. Since the networkin both cases is identical the ratio of Ejl and EjH the cycle ranksdetermined by solvent I and II respectively, should he equal to one. Andsince the ratio of 5af/5ph an(i ?FE/?ph a r e independent of the system usedit is possible to assess the swelling method without a commitment to aspecific network elasticity model.

In Fig. 2 the ratio of EJ values obtained for model networks eachswollen by two different solvents (11, 12) are plotted as function of EJas computed from the crosslinking chemical reaction and stoichiometry.The solvents used for swelling are benzene, toluene, cyclohexane and PDMSoligomers. For the latter x = 0 has been assumed. The data indicate al-arge discrepancy well beyond experimental error, for the majority ofsystems implying the invalidity of the method for the determination of £..It is also evident from Fig. 2 that at high degrees of crosslinking whichwill correspond to relatively low swelling, EJ1/?11 -»• 1. This may be arindication that the non-additive contribution to the free energy becomesless important at low solvent concentrations.

In Fig. 3 EjI/EjH for PDMS model networks swollen by PDMS oligomersof increasing molecular weight (13) are shown as function of the solventmolecular weight ratio. From this figure it is clear that large molecularweight species are less reliable in determination of degree of cross-linking resulting in a 60% error.

The main conclusion from this work is that 20-40% error is expectedin the determination of crosslink density by solvent swelling. Only atvery high degree of crosslinking the method is of any reliability. Thedata presented strongly indicate the incompleteness of swelling theory atits present level of development.

REFERENCES

1. Collins, E.A., Bares, J., Billmeyer Jr., F.W., Experiments inPolymer Science, John Wiley, N.Y. 1973; Orwoll, R.A., Rubber Chom.Tech. 1977, 50, 451.

2. Flory, P.J., Principles of Polymer Chemistry, Cornell UniversityPress, Ithaca 1953.

3. James, H.M. and Guth, E., J. Chem. Phys. 1947, 15, 669.4. Flory, P.J., J. Chem. Phys.,1950, 18, 108.5. Flory, P.J. and Rehner, Jr., J., J. Chem. Phys. 1943, 11, 521; ibid

1950, 18, 112.6. Brotzman, R.W. and Eichinger, B.E., Macromolecules 1983, 16, 1131.7. Gottlieb, M. and Gaylord, R.J., Macromolecules 1984, 17, 000.

359

8. Flory, P.J., J. Chem. Phys. 1977, 66, 5720.9. Flory, P.J. and Erman, B., Macromolecules 1982, 15, 800.10. Brotzman, R.W. and Eichinger, B.E., Macromolecules 1981, 14, 1445;

ibid 1982, 15, 531.11. Gottlieb, M. and Macosko, C.W., unpublished data.12. Meyers, K.O., Bye, M.L., and Merrill, E.W., Macroraolecu]es 1980,

13, 1045.13. Gent, A.N. and Tobias, R.H., J. Polym. Sci. Polym. Phys. Ed. 1982,

20, 2317.

10

Fig. 1. The cycle rank of PDMS networks swollen in benzene as predictedfrom the polymer volume fraction by the different theories.

360

I.U

Oβ

0.6

0.4

/

/

1

• BENZENE - TOLUENE

• BENZENE - POMS

* BENZENE - CYCLOHEXflNE

6 8

x I05 mole/cm3

10 12

Fig. 2. Cycle ranks of the same network computed from swelling results intwo different solvents.

0 4 -

Fig. 3. Cycle ranks obtained from swelling a network in a series ofoligomers of increasing molecular weight.

361

PHASE SEPARATION IN RUBBER MODIFIED FLAME RETARDANT HIGHTg EPOXY SYSTEMS

Hemi N. Nae

Department of P l a s t i c s ResearchThe Weizmann I n s t i t u t e of Science, Rehovot, I s r a e l 76100

INTRODUCTION

P r o p e r t i e s of epoxy systems are determined by t h e i rth ree dimensional network formed during po lymer i za t ion .Most epoxy r e s i n s are based on d i f u n c t i o n a l epoxides of thed i g l y c i d y l type (DGEBA) or t e t r a f u n c t i o n a l such ast e t r a g l y c i d y l ether of diamino diphenyl methane (TGDDM).High g l a s s t r a n s i t i o n t empera tu re , Tg, a re a r e s u l t of usinghighly c ross l inked epoxy systems cured with aromaticdiamines such as diamino diphenyl sul fone (DDS). In t h i swork, a th ree funct ional r e s i n , t h r e e g lyc idy l e ther of t r i shydroxy phenyl methane (TEN) i s used . Due to i t s hit,'.!o r o s s l i n k i n g d e n s i t y , TEN/DDS has a Tg of 334 C ( 1 ) .Epoxy r e s i n s are b r i t t l e and have r e l a t i v e l y low impacts t reng th . Addition of small amounts of rubber topolyfunctional epoxy systems improves crack res i s t ance andimpact strength of the cured systems. This improvement hasbeen a t t r ibuted to the i n - s i t u formation of rubbery domainsof a def in i te size and shape during the cure cycle (2 ) .Phase separation occurs when the rubber and epoxy becomeincompatible. Rubber separates well before the l i q u i d - t o -rubber (gelat ion) t r a n s i t i o n (3 ) . The size and shape of therubbery domains depend on the time and temperature of cure.Brominated polymeric addi t ives (BPA) of the diglycidyl typewere introduced to flame retard graphi te-reinforced epoxycomposites ( 4 ) .The object ive of t h i s work is to study the effect of thaintroduction of reac t ive brominated addi t ives and rubber onthe cure behavior and phase separation of high Tg epoxysystems with implications to fiber reinforced composites.

EXPERIMENTAL

M a t e r i a l s ; T r i g l y c i d y l e t h e r of t r i s (hydroxy pheny l )me thane , TEN (XD 7 2 4 2 . 0 0 L ) , Dow Chemical Company and TGDDM(MY 7 2 0 ) , Ciba-Geigy were cured wi th DDS, C iba -Ge igy . A BPAc o n t a i n i n g 50% Br (F2001P) , Makhteshim Chemical Works, wasp r e - r e a c t e d with a ca rboxyl t e r m i n a t e d b u t a d i e n e -a c r y l o n i t r i l e CTBN rubbe r (Hycar 1300x13) , B.F. GoodrichChemical Company. All f o r m u l a t i o n s were s t o i c h i o m e t r i cc o m p o s i t i o n s (1 mole epoxy with 1 mole amine hydrogen)d i s s o l v e d in methyl e t h y l ketone ( 1 : 1 . 5 w / v ) . The c u r e c y c l e

362

was: i|5 min . a t 30 C/vacuu*. 8 min a t 125 C, 2 hou r s a t 177C. Specimens were pos tcu red at 192 C for 4 h o u r s . Graphi tec l o t h impregnated with the r e s i n s o l u t i o n was curedaccord ing to the same c u r e c y c l e .Dynamic Mechanical A n a l y s i s : An automated t o r s i o n a l bra idana lyze r (TBA), P l a s t i c s Analys is I n s t r u m e n t s I n c . , was usedto moni tor modulus and l o g a r i t h m i c decrement a t about 1 Hzunder Helium flow.Thermogravimetr ic A n a l y s i s (TGA); DuPont 951 TGA wi th 1090Thermoanalyzer was used a t 20 C/min in Nitrogen a t m o s p h e r e .Morphology: Ze i ss O p t i c a l Microscope, Jeol JSM 35C E lec t ronMicroscope and P h i l i p s Scann ing /Transmiss ion E lec t ronMicroscope were used to study morphology. SEM micrographswere of gold coated f r a c t u r e d s u r f a c e s . STEM mic rographswere of samples cured on copper g r i d s . The cured sampleswere s t a i n e d with osmium t e t r o x i d e .A l a s e r l i g h t s c a t t e r i n g ana lyzer was developed to i d e n t i f ythe s i z e of the rubbe ry domains . Ths average s i z e of thephase sepa ra t ed rubber p a r t i c l e s was c a l c u l a t e d from theimage of the s c a t t e r e d l i g h t using Bragg r e l a t i o n .


Three e v e n t s are observed in TBA s p e c t r a of i s o t h e r m a l cureof TEN/DDS or TGDDM/DDS with 19% bromine and 2% rubber( F i g u r e 1 ) . These a r e des igna ted as t he p r e - g e l , g e l a t i o nand v i t r i f i c a t i o n t r a n s i t i o n s ( 5 ) . The system c o n t a i n i n gTEN/DDS g e l s and v i t r i f i e s sooner than the system c o n t a i n i n gTGDDM/DDS. This i s due to d i f f e r e n c e s in the c r o s s l i n kd e n s i t y of the cured macromolecu les . Addit ion of BPA andrubber i n h i b i t s g e l a t i o n and v i t r i f i c a t i o n somewhat but nots i g n i f i c a n t l y even when the amount of a d d i t i v e s i s 30? ofthe sys t em. This i n d i c a t e s t h a t the cure r e a c t i o n dependsmainly on the r e a c t i v i t y of the epoxy r e s i n and the cur inga g e n t . Apparent a c t i v a t i o n e n e r g i e s , c a l c u l a t e d from logtime v s . 1/T are a l s o s i m i l a r to t h o s e of the nea t s y s t e m s .The d i f f e r e n c e in the epoxy r e s i n l e a d s a l so to d i f f e r e n c e sin the g l a s s t r a n s i t i o n t e m p e r a t u r e , Tg, of the s y s t e m s . Tgof t h e modified TEN/DDS i s 273 C and of TGDDM/DDS 220 Ccompared to 3311 and 220 of the nea t sys tems r e s p e c t i v e l y . As u b - z e r o r e l a x a t i o n co r re spond ing to the Tg of t he rubber i so b s e r v e d . After h e a t i n g above 250 C t h e m a t e r i a l s t a r t s todegrade and the Tg of t he rubber becomes more d i s t i n c t .TGA of a l l systems show a c a t a s t r o p h i c weight l o s s a t310-350 C. T h e r e f o r e , the l i m i t i n g f a c t o r for s t r u c t u r a lpurposes i s the Tg of the m a t e r i a l and not i t s thermals t a b i l i t y . However, prolonged exposure above 250 C mayr e s u l t in p a r t i a l d e g r a d a t i o n . The char y ie ld a t 300 C i s35-455, i n d i c a t i n g the flame r e t a d r d a n c y of t he modifieds y s t e m s . The c o n t e n t of bromine was chosen as 195 to impartflame r e t a r d a n c y to the modified s y s t e m s .

Epoxy/DDS with BPA but without rubber form an homogeneousm a t r i x . Upon the a d d i t i o n of r e a c t i v e rubber phases e p a r a t i o n , which accompanies p o l y m e r i z a t i o n , o c c u r s . I t s

363

extent depends on the temperature of cu re , the nature of thereact ing monomers and the rubber concen t ra t ion . Themorphology i s apparently arrested before gela t ion and therubber fores spheres homogeneously dispersed in the epoxymatr ix. The rubber is chemically bonded to the epoxy matrixbut separa tes due to i t s incompat ib i l i ty with the ma t r ix .The rubbery domains are of an average ; lameter of 2-4 m witha core diameter of 1-2 m. The size and shape of the rubberspheres are the same in 2,4 and 6? rubber (Figure 2 ) . Crackpropagation i s mainly un id i rec t iona l and stops at the rubberp a r t i c l e s . Energy d iss ipa ted by th i s mechanism is apparent lyresponsible to the increased toughness of such systems. Insystems containing 8? rubber there are agglomerates in theform of c e l l s of an average diameter of 20-180 m containingmany rubber p a r t i c l e s . The c e l l s have an envelope of anaverage thickness of 10-20 m (Figure 3 ) . The envelope i sprobably a r e s u l t of phase inversion of the rubbery moiety.An optimum in the toughness of such systems about 6J rubberis probably due to the formation of c e l l s which acce le ra t efa i lure of these systems. Phase separat ion is much mored i s t i n c t in TEN/DDS systems than in systems containingTGDDM/DDS due to the d i f ferences in c ross l ink ing densi ty andthe nature of the epoxy r e s i n . Laser d i f f r ac t ion pa t t e rnsshow a short range packing of the rubber p a r t i c l e s with anaverage p a r t i c l e size of 2.6 m which c o r r e l a t e s well withthe size observed in the SEM. Composites prepared fromBPA/rubber containing systems do not show similar phaseseparation since the f ibe r s are c lose ly packed, l imi t ingspace for rubber p a r t i c l e s development. However, rubberydomains are formed and contr ibute to the increased impactstrength of such systems (i<). The d i s t r i b u t i o n of BPA can bemapped by EDAX l ine scanning of bromine which shows anhomogeneous d i s t r i b u t i o n of the BPA in the epoxy mat r ix .Acknowledgement: The help of Dr. Z. Nir, Makhteshim ChemicalWorks, i s g rea t ly apprec ia ted .

REFERENCES

1 . H . N . Nae a n d J . K . G i l l h a m , ACS, D i v . Org . C o a t . P l a s t .C h e m . , P r e p . , 4 8 , 5 6 6 ( 1 9 8 3 ) .

2 . C . K . R i e w , E . H . Rowe a n d A. R. S i e b e r t , i n " T o u g h n e s s a n db r i t t l e m e n t o f p l a s t i c s " , ACS, A d v a n c e s i n C h e m i s t r yS e r i e s N o . 1 5 4 , 1 9 7 6 , p . 3 2 6 .

3 . L . T . M a n z i o n e , J . K . G i l l h a m and C . A . M c P h e r s o n , J . A p p l .P o l y m . S c i . , 2 6 , 0 8 9 ( 1 9 3 1 ) ; I b i d , 2 6 , 9 0 7 ( 1 9 8 1 ) .

4 . Z . N i r , U . J . G i l w e e , D . A . K o u r t i d e s a n d J . A . P a r k e r ,SAMPE Q u a r t e r l y 14 ( 3 ) ( 1 9 3 3 ) .

5 . J . K . G i l l h a m , i n " D e v e l o p m e n t s i n P o l y m e rC h a r a c t e r i z a t i o n - 3 " , J . V . D a w k i n s , E d . , C h a p t e r 5 ,A p p l i e d S c i e n c e P u b l i s h e r s , E n g l a n d ( 1 9 3 2 ) .

364

3 0 10LOO TIMEISEC.I

30 10 50LOG TIME [SEC I

Figure 1: TBA spectra for isothermal cure of TGDDM ( ) and TEN( 1 svstem. (a) relative rigidity (b) log decrement.

Figure 2: SEM micrograph of TEN/DDS,19% Br and 6% rubber.

Figure 4: SEM micrograph ofgraphite fiber rein-forced compositecontaining TEN/DDSwith 19% Br and 6%rubber.

Figure 3; Optical micrographs of (a) TI-N/DDS and (b) TGDDM/DDS with19% Hr and 8%' rubber.

365

RECENT ADVANCES IN THE STRENGTH AND TIME DEPENDENT FAILURE PROCESSOF KEVLAR MONOFILAMENTS AND COMPOSITES

H. Daniel Wagner, S. Leigh Phoenix, P. Schwartz

Cornell University, Ithaca, N.Y. 14853

ABSTPACT

In this paper we discuss recent theoretical efforts to characterize thestatistical strength and lifetime of Kevlar monofilaments andcomposites. We present some new experimental data on the variability ofthe mechanical strength of single aramid filaments. Significantspool-to-spool variability is discovered within a single production lot.The strength ~f single filaments is dependent upon both the flawdistribution and the filament diameter variability across the yarn. Itis found that the Weibull shape parameter for strength is probably not amaterial constant for aramid filaments.

INTRODUCTION

Most brittle fibers used in modern composite materials exhibit largevariability in tensile strength, with measured coefficients of variation(cv) typically ranging froir. 0.05 to 0.25. Kevlar is no exception andshows considerable variability in strength Hue to the presence of flawson the surface and within the interior oF the filament. Simple compositematerials, in which fibers are arranged parallel to each other andimpregnated with a binding matrix, exhibit considerably less variabilitythan the corresponding fibers (cv = 0.03 to 0.07). In creep-ruptureapplications, the variability in lifetime is generally much higher thanthat for the strength, for both single fibers and composites. Clearly,statistical considerations are important when studying the strength andfatigue of fibrous materials. In the present paper we discuss thesignificance of recent experimental data generated in our laboratory withsimple Kevlar fibers in the light of the results from a lnicromechanicalmodel of the statistical failure process developed recently by Tierney[1] and Phoenix and Tierney [2]. This model predicts that a simplerelationship exists between the Weibull shape parameters of the Lifetime(b*) and the strength (b) distributions, namely, that b* = (l+r)b where ris a positive constant which can be determined independently usingresults from stress-rupture experiments. Phoenix [3] discusses thetheoretical justification of the above fiber model in terms of thnkinetic failure c- idealized molecular crystals of the form found instiff polymeric filaments such as Kevlar. In particular, the keyconstant r involves an approximation to the potential function for chainscission and is shown to follow r = If/kT, where T is absolutetemperature, k is Boltzman constant and II has the units of activationenergy.

566

EXPERIMENTAL RESULTS

Five spools of Kevlar (49 and 29) aramid yarns, taken from differentproduction lots (Table 1) were used, and 5 cm long specimens were testedin tension with an Instron Model TM machine at a rate of 2.54 cm/min,under standard conditions (21°C, 65% RH).

Table 1; Specifications for Kevlar aramid yarns

spool type linear density filaments/yarn lot(tex)

ABCUK

K-49K-49K-49K-49K-29

21.621.621.642.244.4

134134134267267

7404874048740487043874043

breakingshape8.84.18.17.87.0

loadscale (N)0.3970.4290.4660.3930.384

(shape10.410.210.49.010.5

tenacity)scale(N/tex)

2.272.452.442.302.50

1. For each spool studied, the filaments were carefully removed and theirlinear density (an indirect measurement of diameter) measured with avibroscope, prior to being tested.

2. A Weibull analysis was performed for breaking loads and tenacities(breaking load/linear density), and the Weibull shape and scaleparameters were estimated using the method of maximum likelihood(MLE). Table 2 presents the results of this analysis.

Table 2: NLE of Weibull shape and scale parameters

spool

ABCUK

note: 50 specimens per spool were used (gauge length = 5cm)

In all cases the shape parameter for tenacity is greater than that forfailure load, that is, the cv for tenacities is smaller than that forbreaking loads, an indication of the elimination of the fiber diameteras a variable. We are not aware of any previous theoretical studybased on weakest link models which makes reference to a possiblevariation in the diameter of individual fibres within the links.Rather, a common assumption is that the fibers in the bundle areidentical. A rather large diameter variability is found among thefibers within spool H (see Figure 1).

3. In order to test the spool-to-spool variability among Kevlar 49 spoolsa formal test of equality of the Weibull parameters, devised by Thonanand Bain [4] was used. The conclusion reached from this analysis(which is discussed in greater depth in [5]) is that the filaments inspools A and I! are significantly weaker than those from spools R andC; that there is no significant difference between spool1; R and C; andthaf there is no significant difference between spools A and U.

367

Fig.l: Filament diameter variability in Spool B (Kevlar 49)

4. In an earlier study, Runsell [6] claimed that the scatterobserved in the failure loads of specimens selected along asimple filament was the same as that observed if the fibers hadbeen sampled randomly from the bundle, showing that the scatter wasdue to the distribution of faults on the fiber rather than due tovariations among the fibers themselves. We performed a replica ofBunsell's experiment with spools A and B and we found that only iffiber variability across a yarn is within limits is Bunsell'sargument valid. In the case of spopl B, for instance, the highdegree of size variability outweighs the effects of flaw distributionand we find significantly greater scatter taking place across theyarn (cv = 0.26) than along it (cv = 0.11).

5. Finally, the effect of filament gauge length was studied for eachspool. The linear behavior predicted by the Weibull theory for log(strength) vs. log (length) is not observed. Moreover, a regressionanalysis yields shape parameters of~18, a value which is abouttwice the expected theoretical value based on the data obtained frorasimple gauge length tests. Interestingly, Phoenix and Wu [7]obtained a Weibull shape parameter of 18.2 by backcalculating fromlifetime data for single filaments using b* = (l+r)b from themicromechanical model alluded to praviously. This is remarkably sim-ilar to the present results which were obtained in a differentmanner. The reason for the discrepancy is unclear. Our data showstronger fibers at shorter gauge lengths, but not as strong aspredicted by the Weibull model; weaker fibers are obtained at longergauge lengths, but not as weak as would be predicted. This behaviorhas been observed in the past in carbon, glass and silicon carbidefibers. Since a progressive change in the slope of the log-logdependence of strength on length is observed, a change in the Weibullshape parameter is probable, therefore suggesting that this parameteris not a material constant- in aramid monofilaments.

368

REFERENCES

1. L.J. Tierney, Adv. Appl. Prob. U_ (1982), 95.2. S.L. Phoenix, L.J. Tierney, Eng. Fract. Mech. (1984, to appear).3. S.L. Phoenix, Proc. 9th U.S. Nat. Congress of Appl. Mech., Cornell

Univ., Ithaca, N.Y. (1982).4. D.R. Thoman, L.J. Bain, Technom. H_ (1969), 805.5. H.D. Wagner, S.L. Phoenix, P. Schwartz, submitted paper.6. A.R. Bunsell, J. Mater. Sci, _10_ (1975), 1300.7« S.L. Phoenix, E.M. Wu, UCRL-53365, LLNL, Univ. California (1983).

369

DEFORMATION PROCESSES IN IMPACT MODIFIED PVC

A. Hadas and A. Siegmann


Technion-Israel Institute of Technology

The impact strength of PVC is usually enhanced by the addition of non

compatible rubbery polymers, which form a second phase distributed in the

PVC matrix. The dependence of impact strength on blends composition is

not linear but has an S shape. As can be seen in Fig.l, there exists a

critical rubber content (7-15 phr) at which the impact strength abruptly

increases and then levels off. The effectiveness of the rubber depends

mainly on its own mechanical behavior, compatibility with PVC and phase

morphology. The mechanical behavior of PVC and modifier are schemat-

ically described in Fig.2. The much lower elastic modulus and much

higher elongation at break of the rubber should be noticed. The rela-

tionship between the impact behavior, the blend structure and the nature

of the energy absorbing processes are the subject of the present work.

Low tensile strain rate behavior of PVC blends are reported. A compar-

ison between the effects of a rubbery modifier and a plasticizer forming

two and one phase systems, respectively, is made.

PVC blends with two different mofifiers (EVA and acrylic rubber) and a

plasticizer (DOP) were prepared by dry mixing followed by roll milling

at elevated temperature. The blankets were molded to form sheets

(0.7 mm in thickness) in a preheated press at 190°C foi J min. Tensile

specimens were cut and tested using an Instron machine a. strain rates

of 50 and 130%/min. The microdeformation processes were followed using

optical microscopy.

Representative stress strain curves of PVC and several of its blends

with DOP are depicted in Fig.3. The major effects are seen in the yield

zone and the elongation at break. The elastic modulus and yield stress

(Figs.4 and 5) gradually decrease with increasing rubber content whereas

first increasing followed by a steep decrease with increasing

u

I

MODIFIER

CONTENTFig.l: Impact St. as a function of

modifier content (schematic).

STRAIN

Fig.2: Schematic α-ecurves.

370

PVC

Ql 0.2 0.3 0.4 0.5 0.6 0.7 O B 0.9 ID Q

Fig.3: Stress-Strain Curves at PVC/DOP blends. (E=130%).

plasticizer content. Increasing modifier content results in a moreductile behavior, a more uniform deformation changing gradually fromlocalized deformation (necking) into a uniform one. Simultaneously, theinstability intensity at the yield zone is depressed namely, thedifference between yield and flow stresses declines.

The microdeformation mechanism in the various systems was studied viaoptical microscopy. In stretched PVC, adjacent to the neck, crazes wereformed, the density of which decreases with distance from the neck (Fig.6).Crazing is known to be the deformation mechanism in brittle polymers.Upon further deformation the neck continues to propagate and the crazeswiden (Fig.7) and some even develop into cracks - sites of failureinitiation. Upon the incorporation of DOP the craze size significantlydecreases. At high DOP content shear bands are formed (Fig.8) whereasat intermediate loadings both mechanisms are active and their interactionis also observed (Fig.9). Upon the incorporation of impact modifiers, inaddtion to the macroscopic changes during deformation, as mentioned above,the microdeformation mechanism is predominantly shear banding as seen inFig.8. It should be mentioned that the deformation mechanism is affectednot only by the blend composition but also by the deformation rate.

In summary, impact modifiers impart large changes also in the low ratetensile behavior of PVC. Differently from impact behavior, the tensilebehavior does not exhibit any abrupt changes at any modifier loading.With increasing modifier content the systems become more ductile, thedeformation is delocalized and the deformation mechanism changes fromcrazing to shear banding.

371

150"M

E

111 100

DOP

5 10 15CONTENT (phr)

20

Fig.4. Elastic Modulus of PVC/Modifier Blends (e = 1.3 min"1),

5 10 15CONTENT (phr)

Fig.5. Yield Stress of PVC/Modifier Blends (e = 1.3 min"1).

Fig.6. Craze System developed inStrained Rigid PVC.

Fig.7. Craze widening in strainedRigid PVC.

372

Fig.8. Shear Banding in deformed PVC Containing 10 Fhr EVA.

Fig.9. Shear Bands arresting crazes developed in deformed PVCcontaining 5 Phr DOP.

373

FATIGUE CRACK PROPAGATION MECHANISMS IN POLYMERS

A. Bussiba, Y. Katz and H. Mathias

Nuclear Research Centre Negev, P.O.Box 9001, Beer-Sheva, Israel.

INTRODUCTION

In recent years more research attention has been given to fatigue crackpropagation rate (FCPR) behaviour in polymeric materials (1,2). In fact,paralleling similar studies in metallic systems, fracture mechanics con-cepts have been adopted to polymers by examining the relationship betweenthe FCPR and the applied cyclic amplitude controlled by the stress inten-sity factor range AK (3). In polymers too, internal and external variablesinfluence the fatigue behaviour. In addition, there are specific factorswhich are related to the time dependent nature of polymers, their uniquestructural changes and fracture modes. For example, particular emphasishas been given to the proper description of the craze-plastic zone forma-tion at the vicinity of the sharp crack tip (4) .

The issue of equilibrium crack - craze - plastic zone in glassy polymersseems to be very important to the static and cyclic loading behaviour,with and without environmental effects. As known, polymers are well recog-nized and used as structural materials. Thus, fatigue studies aimed toexplore fatigue resistance, are necessary in order to improve fatigue lifeevaluations and to refine the available guideness for material selection.

The current study investigates two kinds of brittle polymers, with theintention to meet several objectives; (i) To determine the fatigue crackpropagation curve from the near threshold values, AKtn, up to the criticalfatigue upper value, Kjfc, including frequency influences, (ii) To gatherexperimental results in order to enable comparative analyses, includingBCC and HCP metallic alloys below the ductile - brittle transition temper-ature, (iii) To examine the polymer fatigue macro- and micromechanisms.Here, the transparency of selected polymers is well considered, clearlybeneficial for better observations of the crack front during initiationand propagation stages.

Actually, the problem of fracture mechanics assessment has been addresedin earlier studies (5), but seems relevant and attractive regarding sev-eral other fundamental problems. For example, crack front profiles (6) orlead interaction phenomena in fatigue crack extension (7) have been studiedby realizing the unique advantages offered by polymeric systems.

EXPERIMENTAL PROCEDURE

Two typical brittle polymers, polymethyl-methacrylate (PMMA) and polysty-rene (PS) were selected. Standard mechanical testing was performed byusing uniform tensile specimens. In addition, fracture toughness parame-ters were determined on pre-cracked single edge notched (SEN) specimensfrom room temperature up to Rn°C. REN specimens of IS mm in thickness were

374

also utilized for cyclic loading tests. All initially machined notches wereextended by a razor-blade sharp cut. Fatigue tests were performed by anelectro-hydraulic closed loop system with an amplitude-controller device.Cyclic load amplitudes were controlled by the stress intensity range ofsinusoidal waveform at frequencies of 15 and 80 Hz. Tests were carriedout at room temperature with a load ratio R = P . /P - 0.

Fatigue crack growth extension was optically tracked, and the FCPR curvesda/dN vs. AK were established from the near threshold, AK-j.n, up to thecritical upper bound, Kj£c. The cyclic AK was calculated according to thefollowing relationship:

K = Y- ^ (1)BW

Where AP is the range of the cyclic applied load, a is the crack length,B is the specimen thicknes, W is the specimen width and Y is the geometriccorrection factor given by:

Y = 1.99 - 0.41(a/w) + 18.7fa/w)2 - 38.48(a/w)~ + 53.85(a/w)4 (2)

Crack front and fracture modes were observed by light and scanning electronmicroscopy, with particular attention to microcrack extension rates, bandwidth and mode transitions along the different stages of the FCPR curve.

EXPERIMENTAL RESULTS AND DISCUSSION

Table I summarizes static and dynamic properties for the tested materials.The fatigue properties were obtained from the relevant experimental FCPRcurves for the two applied frequencies. As expected, both polymers behavedin a brittle manner and indicated significant sensitivity to the tempera-ture and to the strain rate.

Table I : Mechanical and dynamics properties of the tested polymers.

Polymer

PMMA

PS

of

MPa

48

35

0

0

eo'o

.4

.3

KIcMPa»m

296K

1.5

2.0

k353K

1.2

1.4

AKthMpa-

15Hz

0.38

0.5

m2

80Hz

0,

0,

.58

.62

KlfcMPa-m

15Hz

1.4

1.8

k80Hz

1.2

1.3

Fig. 1 illustrates the FCPR curve for the PMMA. This typical behaviourwas also obtained for the PS, with similar tendencies regarding the fre-quency effects on the FCPR, on A K ^ and on Kjfc. Moreover, the sigmoidalshape of the da/dN vs. AK curve, as known for metallic systems (8), wasactually preserved. In addition, the whole FCPR curve approached towardsa discrete critical value as reflected by the typical narrow range betweenAK-j. and Kj£c. This fact is not surprizing, since the brittle behaviourof the PMMA and the PS resulted in FCPR curves highly dependent on theapplied AK, as obtained in the case of metallic alloys, below the ductilebrittle transition temperature (8). The role of the frequency confirmsthe results by Skibo et al. (9) and Jilken et al. (10), namely, increasingthe cyclic frequency causes a significant decrease of the FCPR in bothtested nolvmers.

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Figure 2 shows the dependency of the fatigue band size and the number ofcycles per band, N*, on AK for the PMMA and PS at 80 Hz. Figure 3a demon-strates the initiation stage in PMMA as obtained by fracture surface obser-vations. As shown, this stage is associated by extremely low energy fracturesurface. Figure 3b shows microcracking and microcavities at the earlierstage II. Figure 3c illustrates fatigue striations at stage II,while later,at stage III,secondary cracking occurred which enhances the fatigue crackextension rate.

Probably the most significant result which differ completely from metallicfatigue crack propagation behaviour, is connected to the micromechanismsin terms of the discontinuous crack extension. This issue is emphasizedby N and has been addressed earlier by Hertzberg et al (11). As shown, theband size or striations in polymers are nor associated with incrementalcrack extension caused by one cycle only. The latter is well recognized inmetallic materials with an excellent fit between the macroscopic crackextension rate and the microscopic FCPR based on striation spacings (12).

Hertzberg et al (11) have attempted to explain this unique discontinuousgrowth bands in terms of a more general phenomenon, namely, craze formationin polymers which affects also the fatigue crack extension processes.However, it seems that more has to be done in order to extend and to refinethe mentioned idea. Firstly, the exact description of the craze profile isneeded in more realistic terms of a crack-craze-plastic zone configuration(4). Secondly, the craze, as such, is associated with a threshold valuewhich should probably be incorporated in a proposed model.

As concluded by Israel et al.(13), the Dugdale plastic zone model was notfully adequate for describing craze geometries in PMMA. Therefore, a modi-fication of the Dugdale - Rarenblatt model has been proposed(13),describingthe crack tip configuration in terms of crack-craze-plastic zone, and ex-pressed at equilibrium by the following relationship:

•nc - 2(0 - a ) cos" b/a - 2a cos" c/a = 0 (3)

where a is the applied stress, CTC and OyS are the stress within the crazeand the yield strength within the plastic zone, respectively. c,b,a arethe crack-craze-plastic zone geometrical values.

Actually, this model at equilibrium describes the elastic stress intensityterm by means of the fictitious crack length, which includes the cohesiveplastic zone contribution and the semi cohesive term associated with thelocalized micro-cracked and void-filled craze region. Thus, at equilibriumthe craze size is related to the applied stress intensity factor, and amonotonic increasing function is obtained for the craze size dependency onthe applied AK. Furthermore, this modified description indicates the roleof the crack-craze-plastic zone profile on the effective localized stressintensity factor. In fact, the effective stress intensity factor is affectedby craze formation, assuming the retractive stresses exerted by the crazeat the crack tip vicinity.

Referring back to the current fatigue findings in PMMA and PS, clearly,the micromorphology in the tested polymers indicates the role of the crazeformation. On one hand, dynamic craze-plastic zone causes a reduction inthe nominal stress intensity. On the other hand, subsequent incrementalcrack extension might be activated only after an additional cumulative

damage, mainly applied to the microfibrils, and finally resulting in adiscontinuous band extension. In fact, the band size increases with higherAK. Higher values of AK result in higher values of the elastic term, thatmust be offset by the corresponding increase of the dynamic craze-plasticzone size. This tendency was readily observed in the fatigue band size, asshown in Fig. 2. In contrast to the band size, N* decreased with the in-crease of AK, which can be explained by realizing the effective subsequentdamaging potential, available for further crack extension for relativelyhigh AK values. Consequently, the increase of the cyclic range, AK, in-creases the dynamic cra?.e~plastic zone size, but not the total fatigueresistance capacity, which ends up with the enhancement of the fatiguecrack extension rates.

ACKNOWLEDGEMENT

The authors wish to express their appreciation to Mr. M. Kupiec and Mr.Y. Fachima for experimental assistance.

REFERENCES

1. J.C. Radon, Int. Journ. of Fracture, Ij3, 1980, 533-552.2. J.A. Sauer and G.C. Richardson, Int. Journ. of Fracture, j^, 1980, 499.3. S. Arad, J.C. Radon and L.E. Culver, Int. Cong, on Fracture, ICF3,

Munich 1973, Part VII, paper VI - 323.4. R.W. Hertzberg and J.A. Manson, in Materials Experimentation and Design

in Fatigue, F. Sherratt, J.B. Sturgeon and R.A.F. Farnborough Eds.,Westbury House, pp. 185-198, 1981.

5. Y. Katz, P.L. Key and E.R. Parker, Trans. ASME, £0_, Ser. D, 1968, 622.6. Y. Katz, A. Bussiba and H. Mathias, in Mechanical Behaviour of Materials,

ICM3, K.I. Miller and R.F. Smith Eds. Pergamon Press, 1979, 383-391.7. Y.W. Mai, Int. Journ. of Fracture, 15_, 1979, 103-106.8. A. Bussiba, H. Mathias and Y, Katz, Proc. Materials Engineering Conf.

Haifa, Israel 1981, 123-125, I. Minckof Ed=, Freund Publ. House.9. M.D. Skibo, R.W. Hertzberg and J.A. Manson, Proc. Int. Conf. on Fracture,

ICF4, D.M.R. Taplin Ed., University of Waterloo Press, 1977, 1127-1133.L0. L. Jilken and G.G. Gustafson, in Fatigue Thresholds, J. Backlund, A.F.

Blom and C.J. Beevers Eds., Chamelion Press V.2, 1981, 715-733.LI. CM. Rimnac, R.W. Hertzberg and J.A. Manson, in Fractography and

Materials Science, L.N. Gilbertson and R.D. Zipp Eds. ASIM STP 733,1981, 291-313.

12. Y. Katz, A. Bussiba and H. Mathias, Proc. European Conf. on Fractureand the Role of Microstructures, ECF4, K.L. Maurer and F.E. MatzerEds., EMAS Press, 1982, 503-511.

13. S.I. Israel, C.S. Kantamneni and W.W. Gerberich, in Mechanical Behaviourof Metals, ICM3, K.J. Miller and R.F. Smith Eds., Pergamon Press, 1979,393-402.

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378

AUTHORS' INDEX

Aboelfotoh M.

Addess S.

Adler L.

Admon U.

Agronov D.

Arigur P.

Avni R.

Azoulay M.

Bamberger M.

Bar-Ziv S.

Borns te in A.

Botstein 0.

Brandon D.G.

Brat T.

Breyer N,

Buckley D.

Burstein Z.

Bussiba A.

Cahen D.

Carmi U.

Chaim R.

Croitori N.

Cyterman C.

Dagan G.

Dariel M.P.

Dayan D.

291

146

185

82

119

137

128

262

267

272

283

332

35

209

225

328

71

55

299

133

283

332

373

347

267

272

51

55

328

343

299

347

30

82

100

Dirnfeld S.F.Edmonds D.V.

Eisenberg M.

Eisenberg Y.

Elkabir G.

Eliezer D.

Endo S.

Fainaro 1.

Falkenstein E,

Finberg I .

Fink J.L.

Fishman M.

Freund A.

Gafri 0.

Gal-Or L.

Gilad I .

Goldstein J .

Gottesfeld S.

Gottlieb M.

Gri l l A.

Grinbaum Y.

Grossman E.

Grunbaum E.

Grunze M.

Grushko B.

Hadas A.

Hall D.

39

108

299

350

232

152

158

167

347

315

315

283

67

95

119

137

258

95

163

63

77

343

356

249

262

285

253

249

262

82

235

71

88

369

245

379

Harzion Z.Herrmann B.

Hershitz R.

Heuer A.H.

Hodes G.

Huang H.C.W.

Ignatiev A.

Inspektor A.

Iram A.

Jacobsen K.M.

Josephy Y.

Katz Y.

Kazinetz M.

Kimmel G.

Kohn G.

Kornblit

Lahav A.

Laks C.

Lalman J .

Landau A.

Lang S.B.

Levin L.

Livne Z.

Livni T.

Lodder G.C.

Manory R.

Marcu V.

Mark-Markowitch M.

Mathias H.

McBreen P.H.

343

146

21

51

347

175

74

336

267

272

95

221

350

163

373

258

59

82

100

104

146

74

299

92

245

59

324

201

46

225

82

262

304

167

163

373

245

Minkoff I .Minkovitz E.

Moscovits M.

Munitz A.

Nae H.

Nir . A.

Nissenholz Z

Notis M.R.

Nowick A.S.

Pelleg J .

Phoenix S.L.

Polak M.

Prinz B.

Rabin B.

Rabinovitz E

Raveh A.

Rishpon J.

Ron M.

Rosen A.

Rozenak P.

Rosenthal Y.

Rotel M.

Rotem

Rubinstein I

Ruhle,M.

35

158

245

46

361

K2

142

63

77

11

67

92

100

104

320

365

249

253

209

258

146

272

267

272

343

350

119

137

225

232

152

167

175

225

214

51

380

Sar ie l J . 59

104

Schwartz P. 365

Seidman D.N. 21

Sharon A. 216

Shikmanter L. 30

Siegman A. 369

Spalvins T. 283

Stern A. 146

167

201

Talianker M. 30

59

Tamir S. 277

Tarby S.K. 77

Tenne R. 304

Totta P.A. 175

Tu K.N. 288

291

Turnbull D. 1

Zafrir M. 343

Zahavi J. 175

277

308

Zevin L. 92

258

Zuta Y. 39

Wagner H.D. 365

Weinberg F. 128

Weiss B.Z. 71

88

Williams D.B. 63

second israel materials engineering conference

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