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Reprint fromMaterials Science and Technology
A Comprehensive Treatment
Edited by R. W Cahn, P. Haasen and E. 1. Kramer
Chapt.er 7
Zirconium Alloys in Nuclear Applicationsby C. Lemaignan and A.
T. Motta
from Volume 10 B
Nuclear Materials, Part 2edited by B.R.T. Frost
Published byVCR Verlagsgesellschaft mbHP. O. Box 101161, D-69451
WeinheimFederal Republic of Germany vc~
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7 Zirconium Alloys in Nuclear Applications
Clement Lemaignan
CEA/Centre d'Etudes Nucleaires de Grenoble/DTP/SECC, Grenoble,
France
Arthur T. Motta
Nuclear Engineering Department, The Pennsylvania State
University,University Park, PA, U.S.A.
List of Symbols and Abbreviations 27.1 History. . . . . . . . .
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
. . . . . . . . . . . . . . .. 47.1.1 High Temperature Water
Reactors. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
. . . . .. 47.1.2 Current Use. . . . . . . . . . . . . . . . . . .
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
.. 47.2 Fabrication and Products 57.2.1 Processing. . . . . . . . .
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
. . . . . . . . . . . .. 57.2.2 Microstructure...... . . . . . . .
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
. . . . .. 77.2.2.1 Alloys and Alloying Elements " 107.2.2.2 Heat
Treatments and Resultant Microstructure 177.2.3 Properties
187.2.3.1 Mechanical Properties 187.2.3.2 Diffusion Data 227.3
In-Reactor Behavior 247.3.1 Irradiation Damage and Irradiation
Effects 247.3.1.1 Displacement Calculations 247.3.1.2 Irradiation
Effects in the Zr Matrix 257.3.1.3 Irradiation Effects on Second
Phases 287.3.1.4 Irradiation Growth 297.3.1.5 Irradiation Creep
327.3.1.6 Changes in Mechanical Behavior 337.3.1.7 Charged-Particle
Irradiation 347.3.2 Corrosion Behavior 367.3.2.1 General Corrosion
Behavior 367.3.2.2 Oxidation of the Precipitates 377.3.2.3 Water
Radiolysis 397.3.2.4 Hydrogen Pickup 407.3.3 Pellet-Cladding
Interaction 417.4 Challenges. . . . . . . . . . . . . . . . . . . .
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
.. 467.5 Acknowledgements 477.6 References 47
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2 7 Zirconium Alloys in Nuclear Applications
List of Symbols and Abbreviations
a,cao, CoA
-
IGSCCI-SCCLHGRLWRMIBKPCIPKAppmPWRRBMKRXR&DSCCSEMSIPASTEMSOCAP
RTBSTEMTrexUTSVVERYS
List of Symbols and Abbreviations 3
iodine intergranular stress corrosion crackingiodine stress
corrosion crackinglinear heat generation ratelight water
reactormethyl-isobutyl-ketone (process)pellet-cladding
interactionprimary knock-on atompart per millionpressurized water
reactorRussian graphite-moderated boiling water reactorfully
recrystallizedresearch and developmentstress corrosion
crackingscanning electron microscopestress induced preferential
absorptionscanning transmission electron microscopesecond order
cumulative annealing parameterstress relievedto be
specifiedtransmission electron microscopytube-reduced
extrusionultimate tensile strengthVoda-Voda energy reactor, Russian
type PWRyield strength
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4 7 Zirconium Alloys in Nuclear Applications
7.1 History
7.1.1 High Temperature Water Reactors
Soon after the observation of the fissionof uranium 235, L.
Szilard and F. Joliot-Curie recognized the possibility of usingthe
chain reaction phenomenon as a sourceof energy. Initially, test
reactors were de-signed with no constraints on thermal effi-ciency.
The aim was then to understandneutron physics and to study the
behaviorof materials under irradiation. Low tem-perature, pool type
reactors were con-structed in which the structural materialused was
exposed to a comparatively mildenvironment. Aluminum and beryllium
al-loys were used for core components, due totheir low thermal
neutron capture crosssection and acceptable corrosion rate inwater
below 100°C.
Once nuclear power reactors for sub-marine propulsion and
production of elec-tricity were designed, thermal efficiencybecame
mandatory, and materials had tobe found that could withstand the
hightemperature of the coolant, usually water.Zirconium (Zr), with
its very low thermalneutron capture cross section, was a poten-tial
candidate, but had poor ductility andcorrosion resistance. The
first pressurizedwater reactors were loaded with fuel clad-dings
and other structural elements (guidetubes and grids) made of
stainless steel.
An improvement in neutron efficiencywas a driving force for the
development ofindustrial type Zr-based alloys. At the endof World
War II, the nuclear submarineprogram undertook a large effort in
thatfield. Systematic testing and research anddevelopment (R &
D) by the U.S. Navy re-sulted in the development of an
efficienthafnium (Hf) separation process and in-dustrial scale
ingot production procedures.During the test of a series of binary
and
ternary alloys, an accidental contamina-tion of a Zr-2.5 % Sn
(Zircaloy-1) melt bystainless steel caused the serendipitous
dis-covery of an alloy of good corrosion be-havior. Composition
variations aroundthis alloy led to Zircaloy-2. Zircaloy-3, avery
low tin variant, was soon abandonedin favor of the better
Zircaloy-4, aNi-freevariant designed to decrease hydrogenpickup.
Similar tinkering with composi-tional variations to improve
corrosion re-sistance and strength led to the develop-ment by the
Soviet Union of anotherfamily of alloys using the Zr-Nb
binarysystem, later used by Canada as well. TheZr-Nb system allowed
the possibility ofobtaining a fine two-phase structure thatleads to
higher strength.
The subject of Zr metallurgy has meritedbooks (Lustman and
Kerze, 1955) and re-views (Douglass, 1971; Cheadle, 1975) inthe
past. Detailed aspects of the metallurgyof the IV-A series (Ti, Hf,
and Zr) can befound in Vol. 8, Chap. 8, of this Series. It isthe
purpose of this work to review the useof Zr for nuclear
applications and to pre-sent some of the more recent developmentsin
the field.
7.1.2 Current Use
In today's nuclear power reactors, Zralloys are commonly used
for structuralcomponents and fuel cladding. For lightwater reactors
(LWR), the commonchoices are Zircaloy-4 in pressurized
waterreactors (PWR) and Zircaloy-2 in boilingwater reactors (BWR).
The heavy water-moderated natural uranium CANDU re-actor (Canadian
deuterium uranium), aswell as the Russian RBMK reactor, useZr-Nb
alloys.
In fuel assemblies and bundles, claddings are made out of
Zircaloy-2 orZircaloy-4. Those components are exposedto the fission
products at the inner surface
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at temperatures close to 400°C. At theouter surface they are in
contact with lightor heavy water at coolant temperatures(from 280
to 350°C). Typical heat fluxesacross the cladding are in the range
of30-50 W· cm- 2. Those tubes have differ-ent geometries, depending
on reactor de-sign (Fig. 7-1). In PWR's fuel rods clad-dings are 4
to 5 meters long and have adiameter of 9 to 12 mm for a thickness
of0.6 to 2.8 mm. BWR fuel rods are usuallyslightly larger. In
CANDU, fuel bundlesare short - 0.5 m - to allow on-line
refu-elling. The cladding is very thin - 0.4 mm- and is designed to
collapse around theV0 2 pellets early during irradiation. In
theRussian VVER's the fuel rod geometry is3irnilar to PWR's but the
usual claddingalloy is Zr-1 % Nb.
Structural components of the fuel as-semblies are guide tubes
and grids thatcompose the skeleton. They have to with-stand
mechanical stresses during normalor accidental operation as well as
the oxi-dizing hot water. In BWR's each assemblyis surrounded by a
Zircaloy-2 channel boxthat avoids cross-flow instabilities of
thetwo-phase coolant. Geometrical stabilityof those components is a
critical aspect ofcore design as it affects fuel loading
capa-bility, cooling efficiency and neutronphysics behavior of the
core.
In the case of CANDV's and RBMK'sthe coolant is separated from
the modera~tor and flows around the fuel bundlesin pressure tubes,
usually made of Zr-Nballoys. Those large components (10 m x20 cm x
5 mm) are considered as a struc-tural part of the reactor with a
design lifeof tens of years. They are thus exposedto a high
irradiation fluence (up to3 X 1026 n m - 2) in contact with
thecoolant on the inner surface. Mechanicalstability of those large
components affectsthe overall geometry of the reactor.
7.2 Fabrication and Products 5
7.2 Fabrication and Products
7.2.1 Processing
Zirconium is commonly found in natureassociated with its lower
row counterpartin Mendeleev's table, hafnium. Most of thecommon Zr
ores contain between 1.5 and2.5 % Hf. Due to its high thermal
neutroncapture cross section, Hf needs to be re-moved from Zr for
nuclear applications.
The most frequently used ore is zircon(ZrSi04 ) with a worldwide
production ofabout one million metric tons per year.Most of the
zircon is used in its originalform or in the form of zirconia
(Zr02) asfoundry die sands, abrasive materials orhigh temperature
ceramics. Only 5 % isprocessed into Zr metal and alloys.
The processing of Zr alloy industrialcomponents is rather
complex due to thereactivity of the metal with oxygen. Thegeneral
scheme is presented in Fig. 7-2:Ore processing, zirconium/hafnium
sepa-ration, reduction to metal, alloy melting,hot and cold
deformation processing.
The first step is to convert the zirconinto ZrCI4 , though a
carbo-chlorinationprocess performed in a fluidized bed fur-nace at
1200 °C. The reaction scheme is thefollowing:
Zr02(+Si02+Hf02)+2C+2CI2~
~ ZrCl4 ( +SiCl4 +HfCI4) + 2 CO
After this step, Zr and Hf are separatedusing one of the two
following processes:
(i) Wet chemical: after reaction withammonium thiocyanate (SCN
NH4 ) asolution of hafnyl-zirconyl-thiocyanate(Zr/Hf)O(SCN)2 is
obtained. A liquid-liquid extraction is performed with
meth-yl-isobutyl-ketone (MIBK, name of theprocess). Hf-free Zr02 is
obtained afterseveral other chemical steps: hydrochlona-tion,
sulfation, neutralization with NH 3 ,
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6 7 Zirconium Alloys in Nuclear Applications
(b)
(a)
(c)
Figure 7-1. Fuel cladding and other components,made of Zr
alloys, used in different reactor types: (a)PWR fuel assembly
(courtesy FRAGEMA), (b) BWRfuel assembly and channel (courtesy
GEc), (c)CANDU fuel assembly and surrounding pressuretube.
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and calcination. ZrCl4 is the final resultof a second
carbo-chlorination process(Stephen, 1984).
(ii) Direct separation process: this is anextractive
distillation within a mixture ofKCI-AICI 3 as solvent at 350°C. The
vaporphase, generated by a boiler at the lowerpart of the
distillation column, is enrichedin Hf, while the liquid phase traps
the Zr(Moulin et aI., 1984b).
In either case, Zr metal is obtained by areduction of ZrCl4 in
gaseous form byliquid magnesium, at about 850°C in anoxygen-free
environment. Residual quan-tities of Mg and MgCl4 are removedfrom
the "sponge cake" by distillation at1000°C. After mechanical
fracturing, the)ieces of sponge are sorted, giving the ba-sic
product for alloy ingot preparation.
High purity Zr can be obtained by theVan Arkel process. This
consists of the re-action of Zr with iodine at moderate
tem-perature, gaseous phase transport as ZrI4and decomposition of
the iodide at hightemperature on an electrically heated fila-ment,
the iodine released being used for thelow temperature reaction in a
closed looptransport process, according to the follow-ing
scheme:
Zr+2Iz~ ZrI4 (g) ... ZrI 4 ~ Zr+2Iz(g)t I
250- 300°C 1300-1400°C
For industrial alloys, a compact ofsponge containing the
alloying elements -° (in the form of Zr0 2), Sn, Fe, Cr, Ni,and Nb
- in the desired composition, ismelted in a consumable electrode
vacuumfurnace, usually three times. These vacuummeltings reduce the
gas content and in-crease the homogeneity of the ingot. Typi-:al
ingot diameters range between 50 and
80 cm, for a mass of 3 to 8 metric tons.Industrial use of Zr
alloys requires either
tube- or plate-shaped material. The first
7.2 Fabrication and Products 7
step of mechanical processing is forging orhot rolling in the ~
phase, at a temperatureclose to 1050 °C. Hot extrusion is used
toobtain tube shells or Trex (tube-reducedextrusion), while hot
rolling is used for flatproducts. For Zircaloys, at that stage a
~quench is performed to increase the corro-sion resistance of the
final product. Thistreatment controls the distribution of sec-ond
phase particles, if no further process-ing is performed above 800°C
(Schemel,1977). Further reduction in size is obtainedby cold
rolling either on standard orpilger-rolling mills. Low temperature
re-crystallization is performed between thevarious size reduction
steps.
7.2.2 ~crostructure
Pure zirconium crystallizes at ambienttemperature as an
hexagonal close-packedmetal, with a cia ratio of 1.593 (i.e., a
slightcompression in the c-direction comparedto the ideal ratio of
1.633). Lattice parame-ters are ao=0.323 nm and co=O.515
nm(Douglass, 1971). The thermal expansion'coefficients have been
measured by Lloyd(1963) on single crystals. The difference
inthermal expansion coefficients between thea and c-directions (see
Table 7-1) impliesthat the cia ratio tends towards the idealratio
at higher temperatures - i.e., towardsa more isotropic
behavior.
At 865°C, Zr undergoes an allotropictransformation from the low
temperatureh.c.p. (J, phase to body centered cubic~ phase. On
cooling, the transformation iseither martensitic or bainitic,
dependingon the cooling rate, with a strong epitaxyof the (J,
platelets on the old ~ grains ac-cording to the scheme proposed by
Burg-ers (1934) :
(000l)a;11{110}13 and
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8 7 Zirconium Alloys in Nuclear Applications
Chloridedistillationprocess
Chlorination
Carbon
Zircon(Si02 • Zr02 .Hf02 )
Zr/HfSeparation
Hafnium-free ZrCI,T
Compacted pure ZrCI,
T
TY
Inspection ~
T
Figure 7-2. Processing andfabrication showing all thesteps in
the fabrication ofZr alloy components (cour-tesy CEZUS): (a) From
zir-con to sponge Zr, (b) fromsponge to alloy ingot, (c)from ingot
to final product.
Vacuum distillationZr.Mg _ Zr + Mg
Zirconiumsponge
Kroll reductionMgCI2 ZrCI, + Mg - Zr. Mg + MgCI2-+
Mg...
Magnesium-+
Crushing
Blending
(a)
metal. The main properties of the Zr and Zralloys are given in
Table 7-1. It should benoticed that the main reason for selectingZr
as a nuclear material is its low thermalneutron capture cross
section, which isabout 30 times less than that of iron, givinga
better thermal reactor neutron efficiency.
One should also note its strong aniso-tropic behavior. For
elastic properties,the differences in thermal expansion andYoung's
modulus along the main directionof the hexagonal lattice induce the
develop-ment of internal stresses after any heattreatment due to
grain-to-grain thermal
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7.2 Fabrication and Products 9
First SecondRaw melt (third)materials melt
Electrode Ingot
::::':b,.E~:.~i:9~ ·~r4,EI.::dj:~~~~l/'"'1scrap ~ ~ ~ +
WaterAlloy ~ Electrode Arc outadditives I!!!!!I Compacting welding,
Liquid
Water-cooled t poolcopper Water
(b) crucible in
Forgingpress Hollow billet
Wire in coil
--......~ Tubing
Coil
~o
Billet
Cold rolling
Sendzimir cold-rolling mill
Wire rolling mill
)dpen- or close-
dIe forging
! 1-. Forged partsBar
C~ •Machined partsWire drawing
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10 7 Zirconium Alloys in Nuclear Applications
strain incompatibilities; after annealing at500°C, the
(c)-planes are thus in tensionat stresses up to 85 MPa, depending
on theoriginal texture (MacEwen et al., 1983). Ina similar way,
after plastic strain, elas-tic recovery is orientation-dependent
andleads to compression along the (c) planesof almost 200 MPa (Holt
and Causey,1987). For an industrial material, the elas-tic and
thermal expansion coefficientshave to be computed from its texture,
de-scribed either by the pole figures (Rosen-baum and Lewis, 1977)
or better by usingthe complete orientation distribution func-tions
(Sayers, 1987).
The relative solubility of the various al-loying elements in the
II and ~ phases is oneof the bases for the choice of additions
aswell as heat treatments.
7.2.2.1 Alloys and Alloying Elements
The zirconium alloys in use today fornuclear applications are
limited in num-ber: besides pure Zr, only four alloys arecurrently
listed in the ASTM standards(ASTM, 1990). Those are shown in
Table7-2. The first three are used for claddingand structural
materials, like guide tubesin PWRs and BWRs, channel boxes inBWRs
and structural materials in CANDUreactors, while the last one,
grade R 60904,is used exclusively in pressure tubes forCANDU
reactors.
For cladding tubes, only Zircaloy-2 and4 are listed in ASTM B
811-90. Other alloyshave been developed during the history
ofnuclear power, but except for the Zr-1 0/0Nb alloy used for
cladding in RussianPWR's (VVER) (Tricot, 1990), none are incurrent
use anymore, except for specializedapplications such as the
Zr-Nb-Cu alloyused in garter springs for CANDU pres-sure tubes.
The needs for better performance of nu-clear fuel assemblies and
structural parts,mainly with regard to corrosion resistance,has led
metallurgists and fuel designers tointensive R&D efforts in
order to improvethe properties of those Zr alloys by ad-vanced
compositions and thermomechani-cal processing, and to optimize the
mi-crostructure within the current ASTMalloy specifications.
Indeed, although safety concerns restrictthe introduction of new
alloys because ofthe large amount of data that needs to
beaccumulated to verify the safe behavior offuel elements in case
of reactor accidents,the specifications for the alloys used
todayare broad enough for optimization ofproperties within the
specified compos.tion ranges. Moreover, the microstruc-tures may be
varied significantly becauseof the ll-B phase transformation of
zirco-nium, and because of the different solubil-ities of the
alloying elements in the differ-ent phases. The main alloying
elements arenow considered in turn:
Oxygen is to be considered as an alloy-ing element, and not an
impurity. It isadded to the compacts before melting assmall
additions of Zr0 2 powder. The usualoxygen content is in the range
of800-1600ppm and its purpose is to increase the yieldstrength by
solution strengthening. A 1000ppm oxygen addition increases the
yieldstrength by 150 MPa at room temperature,(Armand et al., 1965).
Oxygen is an II sta-bilizer, expanding the II region of the
phasediagram by formation of an interstitialsolid solution.
The Zr-O phase diagram is given inFig. 7-3: at high
concentration, oxygenstabilizes the II phase to liquid
temperatures: During high temperature oxidation,simulating a
reactor accident, a layer ofoxygen-stabilized ll-zirconium is found
be-
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7.2 Fabrication and Products 11
I Table 7-2. Composition range of standard Zr alloys.ASTM Ref. R
60802 R 60804 R 60901 R 60904
Common name ZircaIoy-2 Zircaloy-4 Zr-Nb Zr-Nb
Alloying elements (mass %)
Sn 1.2 -1.7 1.2 -1.7Fe 0.07-0.2 0.18-0.24Cr 0.05-0.15
0.07-0.13Ni 0.03-0.08Nb 2.4 -2.8 2.5-2.80 to be specified on order
usually 1000-1400 ppm 0.09-0.13 TBS
Impurities (max. ppm)
Al 75 75 75 75B 0.5 0.5 0.5 0.5Cd 0.5 0.5 0.5 0.5C 270 270 270
150Cr 200 100Co 20 20 20 20Su 50 50 50 50Hf 100 100 100 50H 25 25
25 25Fe 1500 650Mg 20 20 20 20Mn 50 50 50 50Mo 50 50 50 50Ni 70 70
35Ni 80 80 80 65Pb 50Si 120 120 120 120Sn 50 100Ta 100Ti 50 50 50
50U 3.5 3.5 3.5 3.5V 50W 100 100 100 100
tween the ~ quenched structure and thezirconia.
Tin is also an a, stabilizer. It forms in thea, and ~ phases a
substitutional solid solu-tion. Tin-based precipitates have been
re-ported in the literature (Bangaru, 1985)but they appear to be
artifacts of TEMsample preparation (Charquet and Alheri-tiere,
1985). At a concentration of 1.2-1.8%, it is used for an increase
in corrosionresistance especially by mitigating the dele-terious
effect of nitrogen in deteriorating
corrosion behavior. Due to a better con-trol of processing
parameters, and conse-quently of nitrogen content, the usage oftin
tends to be lower in the current alloys.Tin, however, also has a
limited impact onmechanical properties, by increasing thetensile
yield strength, and therefore itscomposition should not be
excessively re-duced.
Iron, chromium, and nickel are consid-ered as "~-eutectoids",
because, in theirphase diagrams, these elements have an
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12 7 Zirconium Alloys in Nuclear Applications
oWeight Percent Oxygen
5 10 15 20 25 30 35
Figure 7-3. Zr-O phasediagram (Abriata et aI.,1986).
70
.. ,I I
271 O°C :..~.t.G....... ,L
Q~f:-::2.~~~:~~:
2065°C---F=::::""4LO-...=.::..:.=.-=-----:6~2 :
• I• I
__ _:_1_~~~~~. ..~~~~ :__ :31.2 66.51
-12050 C BZr02_X._---. __ ._.----------- .. -.----------
-------
29.8-- . :: ~?P.o~ .. .. _---_. _...
(aZr)
25%,2130°C
2800
2600
2400
2200
0 2000
~ 1800
~ 1600~ 1400~ 1200
1000:' 29.1 66.7:
800 86.3°C (a' Zr) :' : aZr02_x
600 ( "Z) (C!:l" zr).~ ••••• __ • :-_qQQ~9 .1400 ~ r .~': .....
28.6 66.7:
(al"Zr),~....¥··\\ \7 (a4"Zr) :200
-+------T-'-'''''-----'-r-''-'--'-'--'-.------,-------,--__;---'--...,-----'o
10 20 30 40 50 60
Atomic Percent Oxygen
eutectoid decomposition of the ~ phase(Fig. 7-4 to 7-6). As
mentioned in Sec.7.1.1, they were added to the early Sn-based
alloys after an accidental pollutionby stainless steel of a melting
lot showed anenhancement in corrosion resistance, lead-ing to the
Zircaloys 2 and 4.
At common concentrations, these ele-ments are fully soluble in
the ~ phase. Thetemperature of dissolution of those ele-ments is in
the range of 835-845 °C - thatis, in the upper r:t + ~ range
(Miquet et al.,1982). In the r:t phase their solubility is verylow:
in the range of 120 ppm for Fe and
1009080
Weight Percent Iron
o 10 20 30 40 50 60
702000+---'--,--'----+---',--------'--,-----'--.----'-r------,L--.---L--,--------i
Figure 7-4. Zr- Fephase diagram (Ariasand Abriata, 1988).
100
99.92
(aFe)-
90
1538°C(oFe) ,
-92.9,.'.. 1394°C
8070
90.2 -99.3:, (yFe)~
1357°C
; 92_5_oC ---i912°C
: no°c~gneticTransformation:-ZrFe3..y 470°C
" 1482°C-r-------, ," 1337°C
60
~'I>~~~\'Ii: -2750C
,--------: Magnetic Transformation
50
55.1
403020
L
10
200
1855°C1800 ,
", '1600 ~ \, ,, ,, ,1400 , ,
o "';"(BZr)' ,
i :::: 6\928j~.jl> m ~':4:?------- --~ 863°C r: - - - -;.r.-
-,7- Zr2Fe cD~ 800 '- I' 87835;~ ;',. _-'~ ?C~O~_ _____ LL~ 4.0 ::
N
600 0.02 ::
-(a"Zr) :;400 : : _3000C
~::---------
N ::o+-_---,-----J'L..:-.-----,------,--,---'--..,.---'---
-
7.2 Fabrication and Products 13
Weight Percent Chromium
o 10 20 30 40 50 60 70 80 90 1001900 +-18-55-
o-C.----L--.-----+--,---L------,--L----,r-'----.-'--+--L-'1-86-"-3-oC
-I
Figure 7-5. Zr-Crphase diagram (Ariasand Abriata, 1986).
90 100Cr
80
yZrCr2 66.7%,1673°C 0
..... ----~.:: _\19?. ---. 15920 C 1
------.1-~~71t-1- _ -:~.~ 15320Ct-~~:~~BZrCr2 : ~
:~ (Cr)-.l[)
1332°C : ,....ir---(0
30 40 50 60 70Atomic Percent Chromium
L
20700
+-----.-----,---,-------r--.-----.------'-------',--..,--------,----j
o 10Zr
900
1100
1700
2: 1500~:::l
"§(1)
~ 1300
~
200 ppm for Cr at maximum solubilitytemperature (Charquet et
aI., 1989 a). Forthe Zr-Cr and Zr-Ni binary alloys, the sta-ble
forms of the second phase are Zr2Ni orZrCr2' These phases are
effectively theones observed in the Zircaloys, with Fesubstituting
for the corresponding transi-
tion metal. The Zr3Fe phase which ap-pears in the binary Zr-Fe
diagram is notfound in Zircaloy, probably because itsformation is
too sluggish (Bhanumurtyet aI., 1991).
Therefore the general formulae of theintermetallic compounds in
Zircaloy are
Weight Percent Zirconium
o 10 20 30 40 50 60 70 80 90 100
1800
1600 L
1455°C 1440°C0 1400e.....~:::l
~Q)
@- 1200~
1000 960°C
76(BZr)
845°C
800(aZr)
0 10 20 30 40 50 60 70 80 90 100Ni Atomic Percent Zirconium
Zr
Figure 7-6. Zr- iphase diagram ( ashand Jayanth, 1984).
-
-
14 7 Zirconium Alloys in Nuclear Applications
Zr2(Ni, Fe) and Zr(Cr, Fe)2' In Zircaloy-4,the Fe/Cr ratio of
those precipitates is thesame as the nominal composition of
thealloy. In Zircaloy-2 alloys, the partitioningof Fe between the
two types of intermetal-lic phases leads to a more complex
rela-tionship between nominal compositionand precipitate
composition, giving abroad range of Fe/Cr ratio in Zr (Cr, Fe)2'and
Fe/Ni in Zr2(Fe, Ni) (Charquet andAlheritiere, 1985; Yang et aI.,
1986).
The crystal structure of the Zr(Cr, Fe)2precipitates is f.c.c.
(C 15) or h.c.p. (C 14),depending on composition and heat
treat-ment, with characteristic stacking faults asseen in Fig. 7-7.
Both structures are Lavesphases. The equilibrium
crystallographicstructure is dependent upon the Fe/Cr ra-tio, cubic
below 0.1 and above 0.9, andhexagonal in the middle, following an
em-pirical rule proposed by Shaltiel et aI.(1976). In common
alloys, both types ofstructures are found, even in the samesample,
with random probabilities of oc-currences of each. The Zr2(Ni,Fe)
precipi-tates have a body-centered tetragonal C 16structure
(AI2Cu-type).
The size of these precipitates is of impor-tance for the
properties of the alloys, espe-cially the corrosion rate: while
better uni-
form corrosion resistance is obtained forZircaloys used in PWRs
if they containlarge precipitates, better resistance to lo-calized
forms of corrosion is seen in BWRsin materials that have finely
distributedsmall precipitates. It has been shown thatprecipitation
after ~ quenching is rapid(less than 10 min at 500°C) and that
thecoarsening rate of those precipitates con-trols their sizes in
the final microstructure.It is thus necessary to consider the
com-plete history of the various heat treatmentsfollowing the final
~ quenching to assessthe final precipitate size distribution.
Var-ious cumulative annealing parameters(CAP or LA) have been
proposed for thispurpose, based on recrystallization activa-tion
energy or corrosion behavior. Recently, the coarsening kinetics
have beenshown to be second order with an activa-tion energy of Q/R
= 18 700 K, and a betterdescription of the resultant precipitate
sizedistribution results from using the secondorder cumulative
annealing parameter(SOCAP) (Gros and Wadier, 1989).
In the Zr-2.5 % Nb alloy used inCANDU pressure tubes, very
little Fe isfound in the C1 phase, most of it being in theremanent
~ phase, in metastable solid solu-tion.
200 nm
Figure 7-7. Typical distri-bution of second phaseprecipitates
Zr(Cr, Fe)2 andZr2 i, Fe) in recrystallizedZircaloy.
-
-
7.2 Fabrication and Products 15
Figure 7-8. Zr-Nb phasediagram (Abriata andBolcich, 1982).
Weight Percent Niobium
10 20 30 40 50 60 70 80 90 100
10 20 30 40 50 60 70 80 90 100Atomic Percent Niobium Nb
I
2469°e-~
L /--------/1855°e~ ------~}7L ooe ,-' - -,---
,-'- -21.7
(BZr, BNb)
988°C
863°C -------- 60.6 ~r--..r-........~ V I ~620°C~(aZr) 1 .5
91.0
800
600
400o
Zr
~ 1600'E03 1400D-E~ 1200
1000
o2600
2400
2200
2000
01800~
Niobium (columbium) is a ~ stabilizer.From pure ~-Zr to pure Nb
there exists acomplete substitutional solid solution athigh
temperature (Fig. 7-8). A monotec-toid transformation occurs at
about 620°Cand around 18.5 at. % Nb. By waterquenching from the ~
or upper r:J. + ~ re-gions, the ~ Nb-rich grains transform
bymartensitic decomposition into an r:J.' su-persaturated h.c.p.
phase; subsequent heattreatment below the monotectoid tempera-ture
leads to the precipitation of WNbprecipitates at twin boundaries of
r:J.'needles (Williams and Gilbert, 1966). Inaddition a metastable
co phase can be ob-tained from the ~ by slow cooling or agingof a
quenched structure. A simple epitaxialrelationship is obtained
between the cophase and the parent ~ (Dawson and Sass,1970).
Hydrogen is not an alloying element byeesign, but its behavior
has to be assessed,
since during waterside corrosion, both thehydrogen produced by
the reduction of thewater for oxidation of the Zr matrix and
the hydrogen present for water chemistrycontrol, can be absorbed
to some extentinto the bulk of the alloy. Hydrogen atomsare located
at tetrahedral sites of the h.c.p.cell of the Zr matrix up to the
solubilitylimit (about 15 ppm at 200°C and 200 ppmat 400 DC). Above
the solubility limit hy-drogen precipitates as the equilibrium
()f.c.c. phase (ZrH1.66)' The correspondingphase diagram is given
in Fig. 7-9. Highcooling rates cause the precipitation of
themetastable body centered tetragonal yphase, ZrH (Weatherly,
1981).
Due to the volume expansion inducedby the precipitation of the
hydrides, thisnew phase tends to reduce its strain energyby
nucleating on low index crystallo-graphic planes. Habit planes are
{10tO} forpure Zr and {1 Ot7} for Zircaloys, inepitaxy with the
matrix according to therelationship (111)0 II (0001hr (Bradbrooket
aI., 1972). Further macroscopic growthof the hydride clusters
occurs in the planeof maximum tensile stresses or in the basalplane
if unstressed (Kearns and Woods,
-
16 7 Zirconium Alloys in Nuclear Applications
Hybride
8-E
I'" .'" ..'"oe.•
• + 0 •8 - Hybride 0
oo
Zr (B)
Zr(a)+
Zr(B) o........~-..........----".....'----.e---- -0- -', 0
Zr (E) + 8 - Hybride (Cubic) 0 \ 00
Zr(a)
950.---------------------------,
900850800
750~ 700~ 650::J
Ci5 600~ 550E~ 500
450400350300 l--- -----J
50 : 8: E-HybrideZr (E) + 8 - Hybride (Cubic) 8 :
~:(Tetragonal)
O~--L-----.---------.-____._-_____,--._-__.__-__.__----'--___+_'::""O"":"'___,__--=---_____r'_'
o 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0H/Zr
"'.ott.
Atomic Percent Hydrogen
0 10 20 30 40 501000
900
/863°C800 (BZr)0°i 700:J
~ 600 550°CQ) -'-n. 0.07 -0.659°C 1.43°CE 500~
(aZr)400
300
200
100
00 0.2 0.4 0.6 0.8 1 1.2 1.4Zr Weight Percent Hydrogen
60
I)
;""" E
": :, ,
Figure 7-9. Zr-Hphase diagram (Zuzeket aI., 1990).
1.6 1.8 2
1966). Thus the texture of the material aswell as the stress
state are critical parame-ters for the control of the
precipitationmorphology of the hydrides, a brittle con-stituent at
low temperatures.
Other minor constituents are often foundin the form of
precipitates. Among themare the carbide f.c.c. ZrC and silicides
cphosphides of various stoichiometries(Zr3Si, ZrSi2 , ZrP, Zr3P)
(Charquet andAlheritiere, 1985).
-
-
7.2.2.2 Heat Treatments and ResultantMicrostructure
At high temperature, the oxide layerthat develops during forging
is not protec-tive and therefore 0, N, and H can diffuseinto the
bulk of the ingot if no particularcare is taken.
Some of the precautions that have to betaken for
thermomechanical processingare reducing the time at high
temperaturein unprotected atmospheres and perform-ing intermediate
descalings and etchings ofoxidized forgings.
Thus, after ingot melting, the ther-momechanical processing
commonly usedfor industrial alloys is the following:• Hot forging
in the ~ range (1000 to
1050°C).• Water quenching from the homoge-
neous ~ phase (above 1000°C).• Intermediate temperature (upper
Ct)
forging and rolling, or extrusion fortubes.
• A series of cold temperature roilingsfollowed by intermediate
anneals invacuum furnaces.Homogenization in the ~ phase leads
to
the complete dissolution of all the secondphase particles, but
gives rise to significantgrain growth: after 30 min at 1050°C,grain
size may reach several millimeters.During the water quench, the ~
grainstransform into Ct needles by bainitic trans-formation due to
the slow cooling rate ofthe large pieces involved. These Ct
needlesnucleate at grain boundaries and each for-mer ~ grain leads
to a series of crystallo-graphic orientations corresponding to
the12 different permutations of the Burgersrelationships. This
leads to a typical "bas-ket-weave" microstructure as seen inFig.
7-10 a. The ~-eutectoid elements arerepelled by the transformation
front andprecipitate at the boundaries of those
7.2 Fabrication and Products 17
(o)
(b)
Figure 7-10. Typical microstructure of pquench ma-terial: (a)
Basket weave l1. grains from the same former~ grain. The contrast
obtained in polarized lightshows the 4 different orientations of
the l1. grains fromthe same ~ grain. (b) Second phase precipitation
at l1.grain boundaries. In bright field, the precipitates ap-pear
as small white dots at interplate boundaries.
needles (Fig. 7-10b). This ~ quench is usedfor an increase of
corrosion resistance ofthe final product, and is a reference
statefor further processing. The cold workingsteps increase the
homogeneity of the pre-cipitate distribution.
After each cold working step of plate ortube material, an
annealing treatment ismandatory to restore ductility. It is
usuallyperformed at 550-600°C to obtain thefully recrystallized
material (RX). The re-sultant microstructure is an equiaxed
ge-ometry of the Zr grains with the precipi-
-
18 7 Zirconium Alloys in Nuclear Applications
tates located at the grain boundaries andwithin the grains
(Figs. 7-7 and 7-11).
To improve the mechanical properties ofthe final product, the
temperature of thelast annealing treatment can be reduced toavoid
complete recrystallization. This isthe stress-relieved (SR) state,
characterizedby elongated grains and a high density
ofdislocations.
In the case of Zr-2.5 % Nb alloys, ~quenching in water of small
pieces leads tothe precipitation of at martensite supersat-urated
in Nb. Tempering at intermediatetemperature results in ~-Nb
precipita-tion at the lath boundaries and at twinboundaries within
the lath (Williams andGilbert, 1966), followed by transformationof
at into a. When quenching is performedfrom an a + ~ region, a
uniform distribu-tion of a and ~ grains is obtained, and theNb-rich
~ phase does not transform. Inthis later case, however, the texture
is lessuniform along the length of the tube thanin ~ quenching.
After rolling or extrusion,the Nb-rich ~ grains tend to align and
theresultant microstructure is shown in Fig.7-12. By aging, at
temperatures in therange of 500°C, the metastable Nb-rich ~phase
can be decomposed into an h.c.p. co
Figure 7-11. Microstructure of recrystallized Zircaloy-4:
equiaxed a. grains with homogeneous distributionof intermetallic
precipitates.
phase. This gives a sharp increase in me-chanical strength due
to the fine micro-structure obtained by the ~-HD transfor-mation
(Cheadle and Aldridge, 1973). Inthe usual form of the Zr-2.5 % Nb,
the coldwork condition after a + ~ extrusion andair cooling, the
microstructure consists ofZr grains with layers of Nb-rich ~
phase(close to eutectoid composition). Due tothe affinity of Fe for
the ~ phase, most ofthis element is found in the minor ~
grains.These ~ grains are metastable and de-compose upon aging to a
mixture of a-Zrand pure ~-Nb. The a-Zr phase itself ismetastable
and irradiation induced precip-itation of the supersaturated Nb
solid solu-tion can occur, which is believed to im-prove corrosion
resistance (Drbanic et al.,1989).
In the case of Zr-l % Nb used for VVERand RBMK, the
concentration of Nb islow enough to avoid microstructural
evo-lution during service.
7.2.3 Properties
7.2.3.1 Mechanical Properties
The mechanical properties of the Zr al-loys are strongly
dependent on severalparameters such as composition, textureand
metallurgical state. For practical pur-poses, the properties at
300-400°C aremost important and room temperature be-havior is used
mostly for comparison.
As reviewed by Tenckhoff (1988), thedeformation mechanism of the
hexagonalZr follows two main mechanisms, slip ortwinning, depending
on the relative orien-tation of the grain in the stress field.
Dislocation slip occurs mostly on prismplanes in the
a-direction. This is referredto as the {10IO}
-
(1120) RADIALDIRECTION
(1010)LONGITUDINAL ----
-
20 7 Zirconium Alloys in Nuclear Applications
Zircaloy-2Rolled Plate
(0002) POLE FIGURE
R
Figure 7-13. Pole figure of 1.5% cold-rolled Zirca-loy-2 sheet,
showing the distribution of basal planesat 35 DC from the normal
direction, and along theshort transverse direction in the sheet. R
indicates therolling direction and T the short transverse. The
nor-mal direction is perpendicular to the page. The thickcontour
line indicates a random concentration ofbasal poles, the dotted
line is half of random, and eachof the thin contour lines indicates
an increase of halfthe random value, so that the maximum
correspondsto about five times the random value. (Courtesy ofJohn
H. Root, AECL, Chalk River Laboratory.)
After cold processing the (1010)-direc-tion is parallel to the
rolling direction.During recrystallization heat treatment, a30
degree rotation occurs around the c-di-rection and the rolling
direction is thenaligned with the (1 I20)-direction for someof the
grains.
At room temperature, in the annealedstate, pure, oxygen-free Zr,
has a low yieldstrength of 150 MPa. This yield strengthcan be
enhanced by solution strengthen-ing, using alloying elements that
have sig-nificant solubilities in cx-Zr. Oxygen, tinand niobium are
candidate elements to
be considered. Nitrogen would be efficient,but degrades the
corrosion resistance.Tin causes only a small increase in ten-sile
strength (Isobe and Matsuo, 1991).By contrast, the addition of 800
ppm ofoxygen increases the yield strength to300 MPa. As a result of
this, Zircaloyshave minimal yield strengths in the rangeof 250-300
MPa and the Zr-2.5 % Nb al-loy, 300 MPa. As in other metals,
reduc-tion in grain size is also used to obtainhigher strength,
leading to the specifica-tion of a grain index of 7 or finer for
stan-dard products. For all those materials, theductility remains
high (above 200/0). Addi-tional strength is obtained by cold
work-ing, allowing the increase of the yieldstrength above 400-450
MPa. This is fo 'lowed by a final stress-relief heat treatmentto
restore ductility without drastic reduc-tion in strength.
Finally the texture itself can increase al-loy strength by
changing the Schmid factorfor slip or twinning. This can be
observedby the differences in strength between theaxial and
transverse directions. In addi-tion, due to the distorted shape of
the yieldlocus, and consequently of the orientationof the strain
vector, strain is also aniso-tropic.
The resistance to unstable crack growth(fracture toughness) is
relevant to largecomponents under tensile stresses. There-fore, the
fracture behavior of the Zr alloyshas been studied only in the
practical caseof the large pressure vessels for chemicalengineering
(Tricot, 1989) or pressuretubes in CANDU reactors. For
hydrogen-free Zr alloys there does not exist any brit-tle-ductile
transition with temperature asin ferritic steels and the rupture is
alwaysductile. The embrittlement of Zr alloys'related to the
behavior of hydrogen. Dur-ing operation, some hydrogen pick-up
oc-curs (see Sec. 7.3.2.4). This hydrogen pre-
-
cipitates in the form of platelets havinga shape which is a
function of textureand stress state. It can also diffuse alongthe
thermal gradient, giving locally muchhigher concentration of
hydrides. Thosehydrides are not ductile at low tempera-ture, and
the fracture behavior of an alloycontaining significant amount of
hydrogenis strongly dependent on temperature andhydrogen content.
The geometrical distri-bution of the hydrides has also a large
in-fluence: in the case of hydrides parallel tothe macroscopic
plane of the crack, crackgrowth is enhanced by the percolation
ofthe fractured platelets (Asada et aI., 1991).The values reported
for the fracture tough-ness K1c of different alloys lie in the
same_'ange of 120-150 MPa' m I/2 for hydro-gen free or for high
temperature data. Thetransition temperature is reported around250-
300°C for 200 ppm of hydrogen anddecreases with hydrogen content.
For hy-drogen content above 400 ppm, the roomtemperature fracture
toughness is reducedbelow 30 MPa . m I/2 (Simpson and Chow,1987;
Asada et aI., 1991). A reduction ofthe fracture toughness of
Zr-2.50/0 Nbpressure tubes with the neutron irradiationfIuence has
been observed and can belinked to the change in microstructure
(dis-location density and Nb-rich precipitates).
Thermal creep: The stress applied tostructural parts under
reactor conditionsis sufficient to induce creep. Although
in-reactor creep is always mixed with growth(Sec. 7.3.1.4), some
experiments were con-ducted to identify the basic mechanisms
ofcreep strain out-of-flux. Compared withmetals of similar melting
temperature, thethermal creep rates of zirconium alloys arehigh and
design limitations result from this)articular behavior.
Zirconium creeps at low temperature: atroom temperature, grain
boundary slidingis observed at low strain rates. Early creep
7.2 Fabrication and Products 21
experiments on pure Zr have shown a largevariation of the
activation energy withtemperature, indicating numerous mecha-nisms
involved in this deformation pro-cess, that have not been clarified
(Guibertet aI., 1969). In the range of reactor tem-peratures, the
activation energies of 2.7 eVfor standard alloys are close to Zr
self dif-fusion values (see Sec. 7.2.3.2). The stressexponents of
the creep rate are still underdiscussion, reported to be in the
range of 2at high stresses (Matsuo, 1987) or muchhigher at low
stresses (Murty and Adams,1985). The mechanisms considered
arecomplex with predominance of dislocationglide controlled by
local climb.
Due to the texture of the cylindricalcomponents and to the
different deforma-tion processes that could be considered,depending
of loading path, analyticalcreep experiments are based on
biaxialloading: closed end samples are tested intension or
compression under variousinternal pressures. Axial and
diametralcreep strains are measured in order to plotthe creep
stress locus. Testing stress-re-lieved (SR) and fully
recrystallized (RX)Zircaloys, Murty and Adams (1985) haveshown that
the results are best explainedif prism slip is considered as the
activemechanism for RX material and basal slipfor SR.
The response of Zr alloys to abruptchanges in loading stress
level during creeptesting has been discussed in detail and
aconsensus exists that a strain hardeningrule should be used for
moderate changesin applied stresses (Lucas and Pelloux,1981). In
the case of large changes of ap-plied stress for reversal loading
conditions(tensile/compression) some contributionof recovery has to
be considered (Matsuo,1989).
Most of the metallurgical parametersaffect the creep rate. For
example, while
-
22 7 Zirconium Alloys in Nuclear Applications
Zircaloy-2 and Zircaloy-4 have similarcreep strengths for the
same thermome-chanical processing, Zr-2.5 % Nb is muchmore
resistant to creep. The metallurgicalstate of the material also
influences thecreep mechanisms: although a SR materialhas higher
tensile strength, its high disloca-tion density allows it to creep
about 2 to 3times faster than the RX material.
The effect of alloying elements on creepproperties has been
found to be slightlydifferent from their effect on tensilestrength.
The effect of oxygen remains im-portant to improve creep
resistance, but itis smaller than the room temperature
yieldstrength improvement. Although it haslittle effect on tensile
strength, tin increasesthe creep resistance (McInteer et aI.,
1989).Carbon also reduces the creep rate, butdue to its low
solubility (150 ppm) it can-not be used efficiently for that
purpose.Irradiation creep is treated in Sec. 7.3.1.5.
7.2.3.2 Diffusion Data
The large amount of data currentlyavailable on solute and self
diffusion inCl-Zr was reviewed by Hood (1988). In gen-eral terms,
single crystal, pure materialsshould be used for the measurements
sothat there is some confidence that the re-sults represent
intrinsic diffusion data. Fora review of diffusion in solids see
Vol. 5,Chap. 2 of this Series.
The three main considerations in a studyof diffusion in Cl-Zr
and Cl-Zr-based alloysare: the atomic size effect, the anisotropyof
diffusion in the hexagonal lattice andthe effect of impurities.
Atom size: small solute atoms can dif-fuse through the
interstitial sublattice atrates which may be 10 to 20 orders of
mag-nitude faster than self or substitutional dif-fusion. Figure
7-14a (Hood, 1988) shows asummary of measurements of intrinsic
dif-
fusion in Cl-Zr (generally extrapolatedfrom high temperatures),
and made in sin-gle crystal specimens. The fast diffusingsmall
solutes (H, Fe, Ni, up to 0 and N)are characterized by activation
energiesbetween 0.6 and 2.5 eV and pre-exponen-tial factors of the
order of 10- 7 to10- 4 m2 . S-1. Those elements are thoughtto be
interstitial diffusers. Diffusion of thelarger solutes (Zr, Hf, Nb,
Sn) has thecharacteristics of self- or substitutional dif-fusion,
with activation energies of 2.8 to3.3 eV and preexponential factors
of theorder of 10- 4 m 2 . S-1. The best estimatefor the activation
energy of Zr self diffu-sion is 3.3 eV, of which 1.9 eV is the
va-cancy formation energy and 1.4 eV is themigration energy (Hood
et aI., 1992)When there is no experimental data, arough estimate of
the migration energyand preexponential can be obtained fromthe size
correlation curves shown in Fig.7-14 b.
Anisotropy: It has been recognized thatdiffusion along the a and
c-axes in Zr isgenerally different. For that reason, inprinciple,
diffusion measurements in Cl-Zrhave to be made in both directions.
Inpractice this can be done more easily forthe fast solutes (Fe,
Cr, Ni), for which ithas been determined that the diffusion
co-efficients are maximum along the c-direc-tion, by a factor of D
II /D -L = 3 (D II is thediffusion coefficient parallel to the
c-axisand D -L is perpendicular to the c-axis). Forself diffusion
the only measurement avail-able (Hood and Schultz, 1974) givesDII/D
-L = 1.3. The diffusion of other slowdiffusers is al 0 nearly
isotropic (Zhou,1992). The determination of such
diffusionanisotropies is important: a factor of twoin the diffusion
anisotropy has been show(Woo 1988) to produce significant effectsin
sink biases and microstructural evolu-tion through the so-called
DAD (diffusion
-
7.2 Fabrication and Products 23
T(K)1000 800 700 600 500
/3 aL.......,
10-10
10-14
(U
10-180 (r
~- NI
III
.510-22
a
10-26
/ SUB
RE //
10-3Figure 7-14. (a) Self and solute diffusion
/data in a.-Zr (Hood, 1988). (b) Atom sizecorrelation for
migration energy and pre-exponential factor for diffusion in
a.-Zr(Hood, 1988).
1.0 1.2 1.4 1.6 1.8 2.0
(a) 103IT IK-1)
I I I
Ni 0 _- Zr(0 qo Fe -10 I- At .Ta
\ 3.0 ~ ~\ /~(u I
2.5 f- /12 I- \ oCr - .0 I
H~ \ /
'.. \ I~ \ - ~
I14 I- (u / • (r
0 \Fn ./
\ l AU '06 .£:. Ig " Ag II
AlI~~16 - 1.0 f- I -"'- oSb
I. FeNi I
O.S~
,.-
Zr(?) I18 '-
I I I I I I
0.12 0.14 0.16 0.12 0.14 0.16
(b) ELEMENT AL RADIUS (nml ELEMENT AL RADIUS (nml
-
24 7 Zirconium Alloys in Nuclear Applications
anisotropy difference) effect, as mentionedin Sec. 7.3.1.4.
Impurities: The effect of impurities onZr self diffusion has
only recently begun tobe understood. It was determined that
sub-stitutional diffusion in Zr is dominated bythe effect of
residual Fe. Hf diffusion innominally pure Zr ('" 50 ppm of Fe)
wasshown to be higher by up to two orders ofmagnitude than Hf
diffusion in ultrapureZr « 1 ppm Fe). In a conclusive experi-ment,
when Fe was added to the ultrapurematerial, to the levels of the
nominallypure material, the Hf diffusion coefficientsjumped up to
the values observed in thenominally pure Zr (Hood et aI., 1992).
TheFe-enhanced migration energy is 0.7 eV,in contrast to the
intrinsic value of 1.4 eV.It is not known at present what kind
ofdefect configuration and migration mecha-nism creates such a
large diffusion en-hancement, e.g., a tightly bound vacancy-Fe
complex having been proposed (Kinget aI., 1991).
7.3 In-Reactor Behavior
In this section, we examine the behaviorof Zr components when
subjected to a nu-clear reactor environment.
In Sec. 7.3.1 the effects of the neutronflux in causing
irradiation damage to theZr components are discussed. The effectsof
this damage range from an increase inthe concentration of
dislocation loops,causing hardening and loss of ductility andthe
dimensional changes brought about byirradiation creep and growth,
to secondphase dissolution, decomposition and re-precipitation,
possibly influencing corro-sion resistance.
Waterside corrosion of Zr alloys is ex-amined in Sec. 7.3.2.
In-reactor, high tem-peratures, irradiation radiolysis, and,
pos-
sibly, deleterious changes in the materialcombine to accelerate
corrosion rates thatare normally slow outside irradiation.
Theeffects of alloying elements on corrosionrates are discussed and
the special prob-lems of nodular corrosion, and hydrogen/deuterium
pickup are considered.
In Sec. 7.3.3, the problem of stress cor-rosion cracking induced
by the pellet-cladding interaction is discussed. This is atypical
example of an in-reactor corrosionmechanism, since it requires
claddingstresses induced by interaction with the ex-panding V0 2
pellet, while the corrodingagent is iodine originating from
nuclearfissions.
7.3.1 Irradiation Damageand Irradiation Effects
7.3.1.1 Displacement Calculations
At the root of all irradiation effects dis-cussed later in this
section is the displace-ment of atoms from their normal
latticepositions by collisions with incident parti-cles. In the
case of in-reactor behavior, weare concerned with the collisions of
neu-trons with target atoms.
For the case of zirconium and 1 MeVneutrons, the recoil atom
from the originalneutron-atom collision (called primaryknock-on
atom, PKA), has a maximumenergy as high as 44 keV. This PKA
de-posits its energy in a relatively localizedregion, causing
secondary and higher or-der displacements. The localized nature
ofthe energy deposition creates a region ofhigh displacement
density, commonlycalled a collision cascade. The number
ofdisplacements in a collision cascade causedby a PKA of energy E
is v (E), typical val-ues being about 100 to 200.
The procedure for calculating v (E) andrelating it to the number
of displacementsinduced in a material by a given neutron
-
fluence is well explained in the literature(Olander, 1976) and
elsewhere in this Vol-ume (Chap. 9, Sec. 9.4.2), so it need
notdetain us here. In order to perform thecalculation, the
displacement energy Edneeds to be specified. This is the
minimumenergy required to displace an atom fromits lattice site.
According to measurementsin the high voltage electron
microscope(HVEM) (Griffiths, 1987) Ed for the pro-duction of
visible damage in pure Zr isabout 25 eV, varying by approximately1
eV with crystallographic orientation.This energy is much higher
than theFrenkel pair formation energy, reflectingthe fact that the
atom has to go through aDotential barrier in order to be
displaced.
After being displaced from their moststable lattice position,
the atoms have achoice of several interstitial configurationsto
assume, of which the lowest energy onesare more likely to be found.
These configu-rations have been studied by lattice calcu-lations
using interatomic potentials thatpredict macroscopic properties of
Zr withsome accuracy (Bacon, 1988). The possibledefect positions
for the h.c.p. lattice areshown in Fig. 7-15. Calculations
(Fuse,1985) show that the most stable interstitialis located at the
Es position, and has anenergy of 3.83 eV.
Although the calculation for the conver-sion of neutron fluence
to displacementsper atom has to be performed for eachspecific
reactor, with its particular fluxcharacteristics and power history,
a rea-sonable estimate of the conversion factor is2 displacements
per atom (dpa) for each1025 n' m- 2 (E>l MeV). This means
thatthe zirconium rods in a pressurized waterreactor (PWR) receive
about 20 dpa in the
years they stay in the reactor, corre-sponding to about 10- 7
dpa . s -1.
Each atom is therefore displaced on theaverage about 20 times
during the three
7.3 In-Reactor Behavior 25
(/)
x«I
()
1
Figure 7-15. Possible interstitial sites for h.c.p. zirco-nium.
Letters in the figure indicate positions andtypes of interstitials
(Fuse, 1985).
years it stays in reactor. Clearly, a lot ofannealing takes
place concurrently, sincethe Zr cladding maintains its structural
in-tegrity, and most of its good mechanicalproperties. Some of this
annealing takesplace within the cascade itself, the numberof stable
defects produced per displace-ment being much smaller than 1000/0
(Wooand Singh, 1992). However, the sheeramount of rearrangement
that is requiredcreates a potential for microstructuralchanges
under irradiation. The mecha-nisms by which this potential is
translatedinto actual material changes under irradia-tion is
reviewed in the following sections.
7.3.1.2 Irradiation Effects in the Zr Matrix
Under reactor operating temperatures,the point defects that
escape immediate re-combination can migrate to sinks, such asgrain
boundaries, free surfaces and dislo-cations. Due to the unequal
bias of thedifferent sinks for the defects, vacanciesand
interstitials can accumulate at differ-ent sinks, giving rise to
macroscopic ef-fects. For example, when network disloca-tions are
not present in large quantities(i.e., in recrystallized material),
the pointdefects can agglomerate into dislocation
-
26 7 Zirconium Alloys in Nuclear Applications
loops, causing a large increase in disloca-tion density.
Dislocation Structure
The as-fabricated material contains net-work dislocations,
resulting from coldwork. Those are mixed in type, both
-
dislocations of the type 1/2 [0001] and1/2[1123] are found,
which supports thegeneral idea of development of
(c)-typedislocations under irradiation.
Voids
Contrary to the behavior of stainlesssteel, Zircaloy does not
exhibit significantvoid formation under neutron
irradiation(Farrell, 1980). Under electron irradiationswelling is
observed in Zr which was previ-ously injected with helium (He)
(Faulknerand Woo 1980). Accordingly, TEM stud-ies have not found
many cavities and voidsin neutron or charged-particle irradiatedZr
alloys. This was confirmed by Baiget al. (1989) who could not find
any voidsin irradiated Zr by small angle neutronscattering. The
reason is thought to be thatthe gases that could stabilize voids
andcavities (O,H,N) are very soluble in the Zrmatrix, where their
equilibrium concentra-tion can be quite high (Figs. 7-4 and
7-5).
Some cavities have nevertheless been~ound in crystal bar Zr
after neutron irra-diation, in the interstitial denuded zonenear
grain boundaries (Griffiths et al.,1988). This is probably due to
the high
7.3 In-Reactor Behavior 27
Figure 7-17. Di tribution ofvacancy and interstitial loopsclose
to a grain boundaryin crystal-bar Zr irradiatedat 700 K to a
fluence of1.1 x10 25 n'm- 2 (E>l MeV).Loops exhibiting inside
con-trast with a 1212 diffractionvector are vacancy in nature
J (beam direction close to[1213]). Close to the grainboundary,
vacancy loops formpreferentially, in contrast tothe bulk, where
some inter-stitial loops also form (micro-graph courtesy of M.
Griffiths,AECL, Chalk River).
(a)
(b)
Figure 7-18. (a) Formation of
-
28 7 Zirconium Alloys in Nuclear Applications
vacancy concentration in those regions.For reasons unknown,
voids are alsofound near second phase particles. Theirtotal number
density is however quite lowand they are not thought to have
mucheffect in the irradiation behavior of zir-conium alloys.
7.3.1.3 Irradiation Effectson Second Phases
Crystalline to Amorphous Transformation(Amorphization) of
Precipitates
One of the most striking effects of irra-diation on Zr alloys is
the crystalline toamorphous transformation (amorphiza-tion)
observed in the intermetallic precipi-tates Zr(Cr, Fe)2 and Zr2(Ni,
Fe), com-monly found in Zircaloys. Those pre-cipitates, described
in Sec. 7.2.2, undergoamorphization under neutron irradiationas
reported by Gilbert et al. (1985) andconfirmed by Yang et al.
(1986). The trans-formation has been observed in
Zircaloysirradiated at 550 to 620 K (from bothLWR fuel cladding and
structural mate-rial) and at 350 K (calandria tubes fromCANDU
reactors). At 350 K, both typesof precipitates are completely
amorphousafter very low fluences (0.5 to 1 dpa). At thehigher
temperature range, the Zr2(Ni, Fe)precipitates are completely
crystalline,while the Zr(Cr, Fe)2 precipitates are par-tially
amorphous having developed a "du-plex" structure, consisting of an
amor-phous layer that starts at the precipitate-matrix interface,
and gradually moves intothe precipitate until the precipitate is
com-pletely amorphous. This is shown in Fig.7-19 A. A crystalline
core is present asevidenced by the stacking fault contrast,while an
amorphous layer has been formedthat will eventually envelop the
whole pre-cipitate. Amorphization is associated witha depletion of
iron from the amorphous
layer into the Zr matrix, while the Cr con-centration in the
precipitate remains con-stant.
It is thought amorphization occurs be-cause the irradiated
crystalline structure isdestabilized with respect to the
amorphousphase due to the accumulation of irradia-tion damage. A
review of the experimentalevidence and theoretical models for
amor-phization of those precipitates under neu-tron and charged
particle irradiation wasgiven by Motta et al. (1991) and Motta
andLemaignan (1992).
Precipitate Dissolutionand Reprecipitation
After amorphization, precipitate disso-lution is accelerated, as
illustrated in Figs.7-19 B, C, and D. A serrated interface isformed
as the precipitates dissolve, indi-cating either a dissolution of
the precipi-tate along preferential directions (Yang,1989), or a
minimization of interfacial en-ergy by faceting.
After precipitate dissolution, and postirradiation annealing, Fe
and Cr reprecipi-tate in the matrix, forming Cr-rich precipi-tates
close to the original particle, and ironrich precipitates further
away. This il-lustrates the difference in mobility ofthose two
elements in Zr, mentioned inSec. 7.2.3.2. It is possible that some
of theFe and Cr remains in solid solution.
Although precipitate dissolution may beaccelerated by the
amorphous transforma-tion, it also happens under neutron
irradi-ation at higher temperatures, where amor-phization does not
occur, so amorph-ization is not a precondition for
dissolu-tion.
There have been reports of irradia~tion induced precipitation
and dissolutionof other types of precipitates in zirconiumalloys.
ZrSnFe precipitates have been
-
7.3 In-Reactor Behavior 29
Figure 7-19. Precipitateamorphization and dissolu-tion under
neutron irradia-tion (courtesy of W. 1. S.Yang, GE). (A) The
duplexstructure in a Zr(Cr, FE)2precipitate in Zircaloy-4
ir-radiated to 5 x 1025 n' m- 2
at 561 K, showing a crys-talline core surrounded byan amorphous
layer. Ironis depleted in the amor-phous layer. (B), (C), and(D):
Zr(Cr, Fe)2 precipitatedissolution along preferen-tial
crystallographic direc-tions after irradiation to14.7x1025 n ·m- 2.
The in-terface faceting is arrowedin (B).
found in Excel alloys (Zr-3.5 % Sn-1.00/0Nb-1.0 0/0 Mo)
following neutron irradia-tion to 1.5 x 1026 n' m- 2 at 690 K
(Grif-fiths,1988). This was paralleled by the ob-servation by Woo
and Carpenter (1987) ofZrSSn 3 precipitates in Zircaloy-2
fol-lowing irradiation to 7.4 x 1024 n . m - 2(> 1 MeV) at 875 K
(total dose about3 dpa). These last precipitates were laterfound to
be associated with Fe. Also in2r-2.5 % Nb, the Fe-rich ~ phase
loses itsFe to the a. phase, and extensive precipita-tion of
fine-sized Nb rich precipitates inthe a. phase has been reported as
a result(Coleman et aI., 1981). This is shown inFig. 7-20.
Outside irradiation, by contrast, the ~phase decomposes under
thermal anneal-ing to a mixture of Zr-86 % Nb and theintermetallic
(Zr Nb)3(Fe Cr).
Second phase redissolution and redistri-Jution of alloying
elements, can have im-portant consequences for irradiationgrowth
and corrosion resistance, as ex-plained below and in Sec.
7.3.2.
7.3.1.4 Irradiation Growth
Irradiation growth refers to the dimen-sional changes at
constant volume of anunstressed material under irradiation
(Fid-leris, 1988). For Zr single crystals, irradia-tion growth
consists of an expansion alongthe a-direction, and corresponding
con-traction along the c-axis.
In polycrystalline materials the situationis more complex, since
grain boundariescan act as biased sinks for point defects, sothat
grain shape and orientation play alarge role. Usually, however,
growth be-havior of polycrystalline materials alsoconsists, of
expansion along the a-direc-tion and contraction along the c-axis.
Asnoted in Sec. 7.2.3.3, the fabrication pro-cess of Zr alloy
components induces a tex-ture. For Zircaloy cladding tubes,
prismplanes are preferentially aligned perpen-dicular to the axial
(longitudinal) direc-tion, which means that irradiation
growthcauses the axial length to increase and thecladding diameter
and thickness to dimin-
-
30 7 Zirconium Alloys in Nuclear Applications
Figure 7-20. ~-Nb precipi-tation in the a. phase ofExcel alloy
(micrographcourtesy of R. W. Gilbert,AECL, Chalk River).
0.4 ,------------------,
Figure 7-21. Growth strain versus fluence for differ-ent amounts
of cold work, at several irradiation tem-peratures from Rogerson
(1988). The dashed linesshow the growth strain for cold-worked (CW)
mate-rial: the strain is nearly linear with fluence and in-creases
with temperature. The solid lines show thegrowth strain for
annealed material: the growth ratesare noticeably smaller than in
cold-worked material,and also increase with temperature. After the
"break-away" fluence, growth rates in annealed materi .are
comparable to those of cold-worked materia."Breakaway" growth in
annealed materials has beenlinked to the development at that
fluence of
-
fluence of about 3 x 1025 n . m - 2, there oc-curs the
"breakaway phenomenon", thestrain rates jumping to approximately
thevalues observed in cold worked materials(Fidleris, 1988). The
observation of break-away growth has been linked to the
devel-opment of
-
32 7 Zirconium Alloys in Nuclear Applications
highly oriented grains. The parameters ad-justed to the behavior
of Zircaloy werethen used to successfully predict thegrowth of
Zr-2.5 % Nb (Holt and Fleck,1991).
in order to improve the knowledge of theparameters that control
creep. Some ofthis work has been performed in materialstest
reactors, and some during detailed ex-aminations of the behavior of
structuralmaterial in power reactors. The experi-mental procedures
used for those analyses,are based upon pressurized tube expan-sion
stress relaxation measurements ofthin'plates, tensile testing or
shear testing(i.e., springs loaded in tension). The lastmethod is a
specific experiment designedto analyze the pure creep behavior.
Forexample, Causey et al. (1984) were able toderive a contribution
of about 30% (c +a)slip on the 1/3(11L3){10I1} system in ad-dition
to the classical (a)-type basal slipby testing Zr-2.5 Nb in pure
shear usinl.compressed helical springs or twisted tubesat 573 K.
They also showed a weak depen-dence of creep rate on dislocation
density.Irradiation creep in Zircaloy-2 is aniso-tropic and
dependent on texture as shownby Harbottle (1978).
For practical purposes and for the lim-ited range of operating
parameters, the ir-radiation creep strain rate may be de-scribed
accurately with a simple equationof the form:
e=A . ( t)m. (Tn. e-Q/(RT) (7-2)
where the effect of flux, stress and temper-ature can be
separated. The usual values ofthose exponents are m =0.6 to 1, n
close to1. The fact that n is close to 1 means thatlarge creep
strains can be sustained with-out failure. The activation energy,
Q, islow, in the range of 5 to 15 kJ . mol-I, i.e.,0.05 to 0.16 eV'
at- 1 (Franklin 1982).
The task of finding a mechanistic modelthat explains irradiation
creep in Zr alloysis a daunting one, in view of the complexit.of
the alloy its inherent single-crystal andtexture anisotropy and the
effect of theirradiation field. For example, ince void
40001,.,,.,"uuu
//
/ t=2x1014 n.cm b.s-/v
/ v ~
/ ~ t_4X1013 n·cm-
// .. --- . -- _-- t=O-----If ..'.'.. 'o
o
0.2
7.3.1.5 Irradiation Creep
Irradiation creep refers to the slow de-formation under an
external stress, experi-enced by a material under irradiation.
Inanisotropic materials, such as Zr, creep hasalways to be
separated from growth in ex-perimental situations. A convenient way
todo this is to assume that creep and growthare linearly additive
and define the irradia-tion creep strain as the additional
strainthat results when the deformation processtakes place under an
external stress. Thecreep strain is then the total strain minusthe
growth strain.
As shown in Fig. 7-22, the creep defor-mation rate of Zr alloys
is increased underneutron flux. Due to the objective of usingZr
alloys in reactor, this phenomenon hasbeen subject to a large
amount of experi-mental work, reviewed by Fidleris (1988),
0.6
2000 3000
time (hours)
Figure 7-22. Diametral creep of RX Zircaloy underinternal
pressure (330 °e, CTH = 150 MPa). The irradia-tion affects both the
primary and the secondary creep,and the creep rate is increased in
proportion to thedose rate.
*-; 0.4c.~
U5
-
welling is practically nonexistent in Zr al-loys (Sec. 7.3.1.2),
the standard rate theorymodel applicable to cubic metals
couplingirradiation creep and void swelling (Brails-ford and
Bullough, 1972), cannot be usedfor Zr alloys. In addition, many of
the pa-rameters needed for mechanistic models,such as defect-defect
interactions, are notknown for Zr, further complicating the
en-deavor.
Matthews and Finnis (1988) recently re-viewed the literature on
mechanisms of ir-radiation creep. It is thought that defor-mation
during irradiation creep occurs bya combination of dislocation
climb andglide, the climb being controlled by thestress-modified
absorption of point defectsat dislocations. According to the
so-calledSIPA (stress induced preferential absorp-tion) mechanism
(Bullough and Willis,1975), dislocations that have their
Burgersvectors parallel to the applied uniaxialstress,
preferentially annihilate interstitialsthan vacancies, leading to
dimensionalchanges due to the dislocation climb itselfand to the
subsequent dislocation glide(Woo, 1979). Matthews and Finnis
(1988),analyzing the possible origins for SIPAconclude that the
elastodiffusion modelproposed by Woo (1984) is the
strongestcandidate, since it is a first-order effect,involving the
migration anisotropy of in-terstitials in an applied stress
field.
The climb-assisted glide (I-creep) andSIPA mechanisms have
recently been em-ployed by Woo et al. (1990) to derive ana-lytic
expressions of their contribution tothe single crystal stress
compliance tensorand used a self-consi tent model that treateach
grain as an inclusion embedded in ananisotropic medium, for
deriving expres-sions that relate the polycrystalline
creepcompliance tensor with that of a singlecrystal. Using a self
consistent model, anda numerical technique, and using the ana-
7.3 In-Reactor Behavior 33
lytic expressions derived by Woo et al.(1990), Christodoulou et
al.(1992) derivedthe relative contributions of basal, pyrami-dal
and prismatic slip systems to the totalstrain in Zircaloy-2 and
Zr-2.50/0 Nb. Itwas concluded that
-
34 7 Zirconium Alloys in Nuclear Applications
o
1018 1019 1020 1021 1022
(a) Integrated neutron flux (E>1 MeV) (n cm-2)
7.3.1.7 Charged-Particle Irradiation
For the sake of completeness, the use ofcharged-particles
(electrons and ions) forexperimental and theoretical studies of
ir-
radiation effects in Zr alloys is discussed inthis section. The
motivation for studyingirradiation effects with charged-particles
isseveral fold: firstly, because their displace-ment rates are
orders of magnitude higherthan under neutron irradiation,
equivalentdoses in dpa are achieved correspondinglyfaster;
secondly, the effect of experimentalvariables such as temperature,
particletype and dose rate can be studied with rel-ative ease, and,
finally, the irradiated sam-ples are usually not radioactive,
makingthem much easier to handle.
In general, neither the results obtainedin such experiments have
a one-to-one cor-respondence with the results from
neutronirradiation nor should such close corre-spondence be
expected in principle, as theirradiations are quite different.
Electron ir-radiation differs from neutron irradiationboth in the
absence of collision cascades,and in that the electron beam is
focused,causing a much higher electron flux anddisplacement rate.
For ion irradiation, theprincipal difference is that the rate of
ionenergy loss per unit target thickness ismuch higher than that
for neutron irradia-tion, which causes only a relatively thinlayer
close to the surface to be irradiated,albeit at a much higher dose
rate. Sinceirradiation effects are dependent on thebalance between
irradiation damage andthermal annealing, any shift in the
dis-placement rate can affect the observed ef-fects. Also, bulk
effects of neutron irradia-tion, such as irradiation-induced creep
andgrowth, and irradiation hardening are dif-ficult to study with
near-surface irradia-tion. If, however, the characteristics ofeach
type of irradiation are taken into ac-count, then charged-particle
irradiationcan provide valuable information on the(mechanisms of
irradiation damage and mi-crostructural evolution, which can be
ap-plied to neutron irradiation.
Annealed
UTS
YS
Cold worked
UTS ::_:----------------------:::::::::::
YS
~30
c0
~ 200)c0a;~ 10-enc~
This increase in yield strength is associ-ated with a reduction
in ductility, affectingboth the uniform elongation (through
areduction of the strain hardening expo-nent), and the total
elongation, decreasingfrom about 200/0 to 2-40/0 (Morize,
1984).This effect is clearly illustrated in Fig. 7-23from Price and
Richinson (1978).
-
Most of the effects reviewed in Sec. 7.3.1have also been studied
with charged-parti-cle irradiation, often with different
results.Extensive studies of microchemical evolu-tion in Zircaloy-2
and Zircaloy-4 underproton irradiation have been performed byKai et
ai. (1990, 1992). It was found thatproton irradiation may induce
variationsin the chemical composition of intermetal-lic
precipitates and increase the solute con-tent in the matrix. These
changes appear toincrease the nodular corrosion resistanceof the
alloy. Other investigations of theeffect of irradiation on uniform
oxidationin Zircaloy-2 and 4 have been conductedby Pecheur et aI.
(1992) and are reported inSec. 7.3.2.
Irradiation-induced segregation and pre-cipitation has been
observed in at least twodifferent circumstances. Cann et aI.
(1992)have observed that 3.6 MeV proton irradi-ation of annealed
Zr-2.5 Nb at 720 K-770 K to 0.94 dpa causes a fine dispersionof
Nb-rich platelike precipitates to appearwithin the r:x grains that
is absent uponequivalent heat treatment of a controlsample, in
agreement with the neutron ir-radiation results of Coleman et ai.
(1981),reported in Sec. 7.3.1.3. Segregation of tinand
precipitation as ~-Sn at the surface ofthin Zircaloy-2 foils has
also been reportedafter 5.5 MeV proton irradiation to 1 dpaat 350 K
(Motta et aI., 1992) a result thathas not been observed after
neutron irradi-ation of calandria tubes at equivalent
tem-peratures.
Amorphization of intermetallic precipi-tates in Zircaloy-2 and 4
under chargedparticle irradiation has been extensivelystudied, both
experimentally (Motta et aI.,1989; Lefebvre and Lemaignan, 1989),
and
eoretically (Motta and Olander, 1990).The critical temperature
for amorphiza-tion of Zr(Cr,Fe)2 precipitates under elec-tron
irradiation was found to be 300 K
7.3 In-Reactor Behavior 35
lower than under neutron irradiation. ForAr ion irradiation,
although the criticaltemperature was similar to that under neu-tron
irradiation, the transformation mor-phology was quite different,
amorphiza-tion no longer starting at the precipitatematrix
interface.
As mentioned above, bulk effects suchas irradiation creep and
growth and hard-ening, are more difficult to simulate withcharged
particle irradiation, but some at-tempts have been made, especially
inthe area of irradiation creep and growth.Parsons and Hoelke
(1989) developed a"cantilever" beam method, which relatesthe
irradiation-induced creep and growthof a cantilevered Zr beam to
its deflectionwhen exposed on one side to 100 keV Neion
irradiation. The results of the experi-ment differed from those of
neutron irradi-ation, but no detailed modeling effort hasbeen
undertaken as yet to rationalize thedifferences.
When both bulk and near-surface irradi-ation can be understood
in terms of theatomic level mechanisms, then neutron ir-radiation
can be simulated with chargedparticle irradiation. For example,
forma-tion of voids under electron irradiationwas observed in a Zr
sample that had beenpreviously charged with He (Faulkner andWoo,
1980) and dislocation loop forma-tion can be studied with electron
irradia-tion (Woo et aI., 1992). This last study inci-dentally,
supported the idea that micro-structural evolution in Zr alloys is
affectedby the large diffusion anisotropy, as men-tioned in Sec.
7.2.3: irradiation-inducedvoids changed their shape under
electronirradiation, preferentially shrinking in the
-
36 7 Zirconium Alloys in Nuclear Applications
7.3.2 Corrosion Behavior
7.3.2.1 General Corrosion Behavior
Zr alloys are highly resistant to corro-sion in common media and
are used forthat reason in the chemical industry (Tri-cot, 1989).
Those alloys are however notimmune to oxidation and in the
hightemperature water environment found ina power reactor (280 -
340°C at 10-15MPa), corrosion and hydriding controlsthe design life
of fuel rods and other com-ponents. Several international
meetingshave been organized to discuss this impor-tant design
parameter, the latest ones hav-ing been reviewed by Franklin and
Lang(1991).
In the early stages, a thin compact blackoxide film develops
that is protective andinhibits further oxidation. This dense
layerof zirconia is rich in the tetragonal al-lotropic form, a
phase normally stable athigh pressure and temperature. As
de-scribed in Fig. 7-24, the growth of the ox-ide layer thickness
d, follows a power law,usually described by a quadratic
relation-ship
d ex:. Jl (7-3)
The activation energy of 130 kJ . mol- i
(1.35 eV . at-i) for corrosion in the denseoxide regime (Billot
et al., 1989) is equiva-lent to activation of the diffusion of
oxy-gen in zirconia (Smith, 1969). Oxygen isconsidered to diffuse
from the free surfaceof the oxide as 0 2 -, by a vacancy mecha-nism
through the zirconia layer, and toreact with the zirconium at the
matrix-ox-ide interface, but recent studies support thefact that
those ions diffuse mostly throughgrain boundaries (Godlewski et
a1., 1991).
For very thin oxide layers, the zirconiagrows in epitaxy on the
metallic zirconi~m(Ploc, 1983). The Pilling- Bedworth ratIO,or
ratio of the oxide volume to the parentmetal volume, is equal to
1.56 for zirconia.Thus as the oxide grows, the stress buildu.due to
the volume expansion associatedwith oxidation induces a
preferential oxideorientation that reduces the compressivestresses
in the plane of the surface (Davidet al., 1971). This gives rise to
various fi-brous textures. Those compressive stressesare one of the
factors that explain the sta-bilization of the tetragonal form of
zirco-nia in this layer. A chemical effect of thealloying elements
could also be invoked toexplain that stabilization.
Figure 7-24. Uniform corrosionbehavior of Zirca10y-4 in water
at633 K. The oxide thickness fol-lows a power law UP to a
transi-tion (1.5 to 2 mm thickness) andthen remains linear. The
metallur-gical state of the material affectsalso the corrosion rate
(Claytonand Fisher, 1985).
500400200 300Time (days)
100
(Summary of 1600 data points)
200 ....-----.--.....---r-----r-,---r-""I""-01 ---r----,o Alpha
annealed (RXA)180 0 stress relieved annealed (SRA) SR
~ 160 least square fitted line'"0 2or § data spread around
averagea, 140E-120c'g,100
-§, 80'Qj
~ :~!~~Oo
-
As the oxidation proceeds, the compres-sive stresses in the
oxide layer cannot becounterbalanced by the tensile stresses inthe
metallic substrate and plastic yield inthe metal limits the
compression in the ox-ide. The tetragonal phase becomes un-stable
and the oxide transforms to a mon-oclinic form (Godlewski et aI.,
1991). Thismartensitic-type transformation is associ-ated with the
development of a very fineinterconnected porosity that allows the
ox-idizing water to access closer to the corro-sion interface (Cox,
1969). The size ofthose pores has been measured to be verysmall:
Cox (1968) has used a modified mer-cury porosimeter to determine
pore sizessmaller than 2 nm. His work has been con-irmed by
nitrogen absorption kinetics to
pore sizes smaller than 1 nm, correspond-ing to a volume
fraction of the order of1 0/0 (Ramasubramanian, 1991). Once
thistransition has occurred, only a portion ofthe oxide layer
remains protective. Thecorrosion kinetics are therefore
controlledby diffusion of oxygen only through thedense protective
oxide layer next to themetal substrate. Since the thickness of
thislayer remains constant, in the range of1 Jlm (Garzarolli et aI.
1991) the corro-sion rate is constant after this transition(Fig.
7-24).
In this dense oxide layer the structure ofthe zirconia, which
controls the posttransi-tion corrosion kinetics, is complex and
stillunder discussion. Starting from the metal-oxide interface, a
very thin layer of amor-phous oxide has been reported under
par-ticular conditions, about 10 nm in thick-ness (Warr et aI.,
1991). The existence ofepitaxy shows that this layer, if
present,should not be continuous. It is followed by
zone of very small zirconia crystallites,10-20 nm, that become
larger in diameterand columnar in shape further into the ox-ide
layer (Bradley and Perkins, 1989).
7.3 In-Reactor Behavior 37
For thick oxide layers, in excess of50 Jlm, the oxide may spall,
leaving zirco-nia particles free to flow in the coolingwater, and
giving rise to a much thinneroxide film. On the one hand, this
processcould be beneficial since it reduces themetal-oxide
interface temperature, whichis the parameter controlling oxidation
ki-netics. However, design considerations onremaining cladding
thickness after oxida-tion does not allow for massive spalling
incommercial reactors. Other undesirableconsequences of spalling
are the buildup ofactivity in the coolant and safety
consider-ations, since those particles can interferewith the
functioning of valves.
In BWRs, as shown in Fig. 7-25, nodu-lar corrosion is the
limiting design consid-eration. This behavior, specific to the
boil-ing water reactor, can be reproduced bytesting alloys in steam
at 500°C (Schemel,1987). Several mechanisms have been pro-posed for
the nucleation of those nodules,leading to various possible sites
for nodulenucleation: metallic matrix grain bound-aries, local
rupture of the continuous denseoxide at an early stage of growth,
localvariations in composition and precipitatedensities or
crystallographic orientation ofclusters of grains (Charquet et aI.,
1989 b;Ramasubramanian, 1989). As the corro-sion progresses the
nodules grow in sizeand thickness and their number densityincreases
leading to a complete coverageof the metal. The ~ quench structure
de-scribed in Sec. 7.2.2.2 significantly im-proves the resistance
of Zircaloy-4 tonodular corrosion, however as higher bur-nups are
reached uniform corrosion couldthen become a problem.
7.3.2.2 Oxidation of the Precipitates
Since most of the metallic alloying ele-ments pre ent in
Zircaloys are added for
-
38 7 Zirconium Alloys in Nuclear Applications
(a)
(b)
(c)
Figure 7-25. (a) Typical aspect of nodular corrosionof
Zircaloy-4 obtained in a 500°C steam test. (b)Higher magnification
of a cracked nodule in Zirca-loy-4, obtained during a 10.3 MPa
steam test at500°C for 25 h, showing the extensive oxide
crackingassociated with the process. In addition to
circum-ferential cracks in the flanks of the nodules, a
verticalcrack can also be seen running between the nodules[Courtesy
of NFIR (Nuclear Fuel Industry ResearchGroup)]. (c) Greater detail
of the nodule, showing thatthe upper portion of the nodule consists
of a series ofparallel, cracked, oxide layers (Courtesy of
NFIR).
improving corrosion resistance, the mech-anism of their
interaction with the oxida-tion front is of great importance. Due
tothe fine structure of the precipitates nospecific observation of
this process wasavailable until recently, when a few studieshave
been performed using advancedSTEM. It was shown that the
precipitatesare incorporated in metallic form into theoxide layer,
and oxidize only afterwards,deeper into the oxide layer. In
particular,iron has been shown to remain unoxidizedin the dense
oxide layer (Garzarolli et al.,1991; Pecheur et al., 1992). The
precipi-tates slowly release part of their iron in theoxidized
matrix so oxide chemistrychanges as the oxidation proceeds. It
isfound that the oxidation of iron coincidewith the oxide
transition from tetragonalto monoclinic.
Also, some of the crystalline intermetal-lic precipitates are
found to be amorphousafter incorporation into the oxide
layer(Pecheur et al., 1992). This transformationis not caused by
irradiation since those arealso observed in nonirradiated
oxidizedmaterial. The crystal-to-amorphous trans-formation in the
oxide layer could becaused by changes in chemistry, notablyhydrogen
intake (the nearest neighbor dis-tance in the amorphous phase is
biggerthan in the corresponding precipitatesamorphized by
irradiation in the Zr ma-trix).
The effect of the release of iron in theoxide is still under
consideration. A corre-lation has been found between the oxida-tion
of the precipitates and the tetragonal-to-monoclinic transformation
of the zirco-nia, but without clear knowledge of itsorigin
(Godlewski et al., 1991). A chemicaleffect on the stabilization of
the tetragonphase of the oxide may be present.
An additional point is the role played bythe precipitates as
possible short circuits
-
7.3 In-Reactor Behavior 39
*
0.8
*
*
0.1
o
o
,I
*,*1** II
// 500°C/16h
/
~/ *o it ..., ......g> ----'
o *0
out-at-pile
3
2 2oL..
10
co'inoL..o 0.02u
~ 30
-
40 7 Zirconium Alloys in Nuclear Applications
gen to the cooling water. Then a generalreduction of stable
radiolytic species oc-curs and a reduction in corrosion rate ofthin
oxide ensues. However this effect isnot as efficient in boiling
conditions com-pared to pressurized ones, due to the segre-gation
of Hz in the steam phase (Ishigureet aI., 1987).
Radiolytic enhancement of corrosioncan also occur in the case of
PWR's, inpost transition thick oxide films, or duringcorrosion
testing of coupon specimens in areactor corrosion loop. There,
although nosteam is expected to be present, the waterchemistry in
the pores of the thick oxideshould be much like that present in
BWRs,i.e., a two-phase regime (Johnson, 1989).
Some particular cases of localized corro-sion enhancement have
recently beenlinked to an increase in metastable oxidiz-ing species
due to ~- radiolysis. The irradi-ation enhancement of corrosion is
well ob-served for thick oxide films, where a three-to four-fold
enhancement is commonlyseen (Marlowe et aI., 1985).
The reported occurrences are character-ized by the presence of
dissimilar materialsin the vicinity of the corroding material.As
recently reviewed, enhanced corrosionhas been observed in front of
stainlesssteel, copper cruds, platinum inserts and inthe case of
gadolinia bearing poison rods(Lemaignan, 1992). In all the
reportedcases, strong ~ emitters are present and thelocal energy
deposition rates of those par-ticles within the coolant are much
largerthan the bulk radiolysis contribution dueto the neutrons and
gammas. Thus, inreactor, the local intensity of the radiolysishas
to be considered, as it will change thelocal chemistry for alloy
corrosion.
In addition to the radiolysis describedabove, corresponding to
chemical evolu-tion by reaction in the bulk, specific
con-siderations should be given to the fact that
the pores have a large surface to volumeratio, leading to
additional reactions at thesurface of the pores, instead of a
recombi-nation between them as in the bulk. Forinstance, the HzOz
molecules can be ad-sorbed on the surface of the pores leadingto
the reaction:
HzOz ~ HOads . +HOads .
instead of recombining in two steps with2 H as the standard back
reaction to waterafter radiolysis. This leads to a higher oxy-gen
potential at the surface of the ZrOz ,giving a higher corrosion
rate.
7.3.2.4 Hydrogen Pickup
During oxidation of Zr alloy compo-nents in reactors or in
autoclaves, the re-duction of water by the Zr alloy followsthe
general reaction scheme:
2HzO+Zr ~ ZrOz +4H'
As described above, the reaction pro-ceeds in several steps and
the oxidationprogresses at the metal-oxide interface bydiffusion of
Oz - ions through the oxide.The reduction of the water molecules
atthe coolant-oxide interface releases hy-drogen as radicals H+.
They are chemi-cally adsorbed at the tips of the oxidepores and
their evolution controls thebehavior of the hydrogen. Most of
themrecombine, creating hydrogen moleculesthat escape through the
pore and dissolveinto the coolant. A limited amount caningress in
the oxide and diffuse through tothe metallic matrix and then react
with Zrfor the formation of hydrides, when theterminal solid
solubility is exceeded. Cor-rosion experiments performed using
tritium-doped water have shown that the Hdissolved in the coolant
is not trapped bythe matrix, but that radicals obtained dur-
-
ing the reduction of water are necessary forhydrogen pick-up
(Cox and Roy, 1965).
The diffusion coefficient of H in Zr0 2has been measured with
difficulty becauseof its very low value and because of
thecontribution of grain boundary diffusionand surface diffusion to
diffusion in thepores of the oxide. Recent measurementsby Khatamian
and Manchester (1989)have confirmed earlier results in the rangeof
10- 17 to 10- 15 m 2 . S-1 at 400°C, de-pending on the condition of
measurementsand chemical composition of the alloy onwhich the oxide
is grown. With such lowvalues, the penetration of H in the oxide
islimited, and the zirconia layer is indeedorotective with respect
to H ingress.
Care should be taken not to introducecatalyzers of the hydrogen
molecule disso-ciation reaction that could enhance Hpick-up (or
hydrogen uptake) - i.e., thefraction of the hydrogen produced by
re-duction of water that is trapped into the Zralloy.