Al-Cu-Mg alloy.pdf

THE NUCLEATION AND GROWTH OF S’ PRECIPITATES IN AN ALUMINIUM-2.5 % COPPER-l.2 % MAGNESIUM ALLOY *

R. N. WILSONt and P. G. PARTRIDGE?

The nucleation and growth of S’ precipitates in an aluminium-2.5% copper-1.2% magnesium alloy aged at 190°C have been examined by transmission electron microscopy. The S’ precipitates have been shown to be heterogeneously nucleated at dislocations, and to grow as laths on {210} planes in a (001) direction. The shape of the laths is consistent with the degree of mismatch of the matrix and precipi- tate planes at the precipitate-matrix interfaces. Precipitate sheets are formed at dislocations by the growth of these laths on (210) planes with a common (001) growth direction; the growth planes conform to the strain energy requirement that the misfit vector of the precipitate should lie approximately parallel to the Burgers vector of the dislocation, and are those expected from consideration of the effect of the line tension of the dislocation. The possibility of refining the precipitate distribution is discussed, and an increase in the density of sites for heterogeneous nucleation (e.g. by prestrain prior to ageing) or a modification of the interfacial energy (e.g. by addition of silicon) are both shown to result in refinement of the S’ precipitate distribution.

GERMINATION ET CROISSANCE DE PRECIPITES S’ DANS UN ALLIAGE D’ALUMINIUM A 25% DE CUIVRE ET 1.2 % DE MAGNESIUM

La germination et la croissance des pr&cipit& S’ dans un alliage d’aluminium contenant 2,5 % Cu et 1,2 % Mg vieilli Q 190°C ont BtB Btudikes au moyen de lames minces en microscopic Blectronique. La germination des p&&pit& S’ se fait de fapon h&&rog&ne, aux dislocations, et ces prkipit& croissent en prenant la forme de lattes sur les plans {210} et suivant une direction (001). La forme de ces lattes est fonction du degr6 de d&orientation de la matrice et du plan du pr&ipitk It l’interface p&ipit&matrice. Des palquettes de pr6cipitk se forment aux dislocations par croissance des lattes suivant les plans {210} et avec une direction de croissance (001) commune. En accord avec les don&es d’bnergie de d&formation, les plans de croissance sont tels que le vecteur de d&orientation du pr&ipit& soit approximativement paralkle au vecteur de Burgers de la dislocation; dislocations permettaient de pr6voir ces faits.

des consid6rations sur les effects de tension de ligne des Les auteurs discutent ensuite de la possibilitk d’affiner la

distribution du pr&ipit& Un tel effet est produit soit par un accroissement de la densite des sites de germination h&&og&ne (par exemple par vieillissement aprk une pr&d&formetion) soit par une modifica- tion de 1’6nergie de l’interface (par exemple au moyen de l’addition de silicium).

KEIMBILDUNG UND WACHSTUM VON S’-AUSSCHIEDUNGEN IN EINER

ALUMINIUM-2.5 % KUPFER-1.2 % MAGNESIUM-LEGIERUNG

Mit Hilfe des Elektronenmikroskops wurden Keimbildung und Wachstum von S’-Ausscheidungen in einer Aluminium-2.5 o/o Kupfer-1.2 ‘A Magnesium-Legierung nach Alterung bei 190°C untersucht. Es wurde gezeigt, daO die S’-Ausscheidungen durch heterogene Keimbildung an Versetzungen entstehen. Sie wachsen in Form von Nadeln in der {210}-Ebene in (OOl)-Richtung. Die Gestalt dieser Nadeln ist konsistent mit dem Ma13 der Fehlorientierung von Matrix und Ausscheidungsebenen an den Zwischen- fliichen. Pliittchen von Ausscheidungen werden an Versetzungen durch Wachsen dieser Nadeln auf {210}-Ebenen mit einer gemeinsamen Wachstumsrichtung (001) gebildet. Die Wachstumsebenen stimmen iiberein mit der Bedingung fiir die Verzerrungsenergie, dalj der Fehlorientierungsvektor der Ausscheidung ungefkhr parallel zum Burgersvektor der Versetzung leigen sollte. Es sind genau die- jenigen, welche man auf Grund einer Betrachtung des Einflusses der Linienspannung auf die Versetzung erwertet. Es wird die MBglichkeit einer Verfeinerung der Ausscheidungsverteilung diskutiert. Es wird gezeigt, dal3 sowohl eine Zunahme der Dichte der Orte fiir heterogene Keimbildung (z.B. durch Vorver- formung vor der Alterung) als such eine Anderung der Zweischenfliichenenergie (z.B. durch Zusatz van Silizium) zu einer Verfeinerung der Verteilung der S’-Ausscheidungen fiihrt.

INTRODUCTION

The aluminium-copper-magnesium alloys containing

a copper : magnesium weight ratio of2.2 : 1 occur in the

pseudo-binary aluminium-S system. This S phase was

shown by Perlitz and Westgrencl) to have the compo-

sition Al,CuMg and to be face centred orthorhombic

(‘C’ face centered) with a = 4.00, b = 0.23 and c =

7.14 A. The supersaturated solid solutions of these

alloys precipitate on ageing in the following manner : Supersaturated solid solution -+ GPB zones -+ S’ -

S(A1,CuMg) and the S’ precipitates were shown by

Bagaryatskiit2) and by Silcockc3) to have an orientation

relationship to the aluminium matrix of:

strains in the matrix and the precipitates nucleate

heterogeneously upon dislocation loops and helices

formed during the quench.(4B5p7)

The purpose of the present work was to study by

transmission electron microscopy the morphology of

the precipitation of S’ in more detail and to discuss the

factors governing its growth and possible methods by

which its distribution might be improved.

MATERIALS AND EXPERIMENTAL METHODS

The aluminium alloy studied contained 2.52 wt. %

copper and 1.20 wt. o/o magnesium in the form of

0.005 in. strip. Specimens from the strip were solution

treated at 505°C for 1 hr in a salt bath and quenched

[1W, II PW_4,3 [0101, II [0211_4,, [001ls II ro121,, into water at 20°C. Some specimens were subsequently

The growth of this phase introduces small lattice aged at 190°C in a salt bath and air cooled. Foils for

study by transmission electron microscopy were pre-

pared from the as-solution treated specimens by elec- * Received March 13, 1965. t Chemistry, Physics and Metallurgy Department, Royal

tropolishing in a standard perchloric acid-glycerol-

Aircraft Establishment, Farnborough, Hants., England. alcohol solution at below 10°C. Foils of aged material

ACTA METALLURGICA, VOL. 13, DECEMBER 1965 1321

1322 ACTA METALLURGICA, VOL. 13, 1965

A slower quench or a short ageing treatment at eleva-

ted temperature resulted in the growth of these loops

as shown in Fig. 2. The formation of loops by climb

has been discussed by Embury and Nicholson@) who

showed that they lay on (110) planes with a/2(110)

type Burgers vector. The helical dislocations shown in

Fig. 1 also grew in diameter by a process of climb on

ageing at elevated temperature. On further ageing at

190°C small rod-shaped S’ precipitates were nucleated

at these loops and helices as shown in Fig. 3, and analy-

sis of their projected images with the electron diffrac-

tion patterns from the foils showed that they lay in the

plane of the loops and in a (100) direction (Fig. 3).

As ageing proceeded the rods grew in the (100) direc-

tion and widened to form laths. Trace analysis showed

that these laths lay on (210) planes having a common

(100) zone axis and so formed corrugated sheets of

FIG. 1. Al-2.5% Cu-1.20% Mg strip solution treated at 505°C and cold water quenched showing dislocation loops

and helices. x 50,000

were prepared by electropolishing in Lenoirs solution

(a chromic-phosphoric-sulphuric acid solution) at 70°C.

RESULTS

When the aluminium-2.5 % copper-l.2 % magne-

sium alloy was cold water quenched from the solution

treatment temperature many dislocation loops and

helices were produced in the matrix as shown in Fig. 1.

FIG. 3. Al-2.5% C~-1.207~ Mg strip S.T. and C.W.Q. and aged 0.1 days at 190°C. The growth of rods in the [OlO]

direction from dislocation loops. x 35,000

FIG. 2. Al-2.5% Cu-1.20% Mg strip aged 10 min at 70°C after solution treatment at 505°C and quenching into cold water showing the growth of large dislocation loops.

x 30,000

precipitate as shown in Figs. 4 and 5. The chevron

markings on the precipitate sheet, A, in Fig. 5,

clearly illustrate the corrugated nature of the sheet.

In Fig. 4, the laths can be seen to have grown inde-

pendently causing the growth front of the sheet to

become irregular. Where more complex dislocation

configurations were present, several (210) planes were

utilised for precipitate growth, all having a commom

(OOl),, growth direction (Figs. 6 and 7). Figure 6

shows a quench band of dislocation helices with S’

precipitates growing on them, some loops forming

completely closed rings of precipitate (A). Figure 7

shows precipitate growth at a dislocation which was

not lying in a single crystallographic plane and from

the different diffraction contrast at the individual

precipitate laths it can be seen that precipitates on

at least three (210) planes formed the complete pre-

cipitate sheet.

WILSON AND PARTRIDGE: S’ ‘RECIPITATES IN Al-Cu-Mg 1323

FIG. 4. Al-2.5% Cu-1.20% Mg strip S.T. and C.W.Q. aged 20 hr at 190°C showing the formation of composite precipitate sheets with irregular growth front. x 40,000

Under certain diffracting conditions producing low

precipitate contrast it was possible to reveal dislocation

contrast around the individual precipitate laths as

shown in Fig. 8 at A and B. These dislocations may

have formed as a result of the misfit strains between

the precipitates and the matrix (see Discussion).

DISCUSSION

The thin foil electron microscopy has shown that the

S’ precipitate laths grow on (210) planes. Depending

upon the nature of the heterogeneous nucleation site,

several (210) planes having a common (001) growth

direction may be used to produce a composite precipi-

tate sheet. From the stereographic projection (Fig. 9)

Fm. 5. Al-2.5% C&1.20% Mg strip S.T. and C.W.Q. aged 16 hr at 190°C showing chevron contrast on the composite

precipitate sheets A. x 60,000

FIG. 6. Same treatment as Fig. 5, S’ precipitate growth on quench bands to form closed precipitate loops (A).

x 40,000

it can be seen that for a given [OOl] growth direction

four possible (210) pl anes may contribute to the for-

mation of the composite sheet, viz. @lo), (i20), (120)

and (210). By selecting two of these planes a com-

posite sheet may form on {OlO}, (110) or (130) planes

as shown in Table 1.

Factors affecting the growth of S’ laths

The nucleation of precipitates in a metastable mat-

rix is dependent upon three energy changes, the volume

free energy change, the strain energy change produced

in the lattice by the precipitate and the increase in

FIG. 7. Formation of a sheet of S’ precipitates at a dis- location. Contrast effects show that precipitation has oc- curred on at least three {ZlO} planes. Al-2.5% C~-1.20~/~ Mg strip S.T. and C.W.Q. and aged 16hr at 190°C.

x 75,000

1324 ACTA METALLURGICA, VOL. 13, 1965

FIG. 8. Al-2.52% Cu-1.20% Mg strip, ST. and C.W.Q. aged 20 hr at 190°C showing dislocation loops A, B associ-

ated with individual S’ precipitate laths. x 30,000

iol

FIG. 9. Stereographic projectionshowingthepossible (210) precipitate planes containing the [OOl] growth direction. Also shown are the possible composite {OlO}, {l 10) and

{130} precipitate sheets they may form.

TABLE 1

Apparent composite {210} planes sheet plane selected

Angle between {210} planes

selected

(010) (i20) (120) 0 8’ (210) (210) l& 52’

(110) (210) (120) 36” 52’ (720) (2iO) 143” 8’

(130) (210) (720) 90”

interfacial energy due to the formation of the new pre-

cipitate-matrix interface i.e. the free energy change

becomes,

AF = AFvo~ume + AFsutiace + AFstrain A precipitate will grow therefore if the release of

volume free energy is sufficient to provide the neces-

sary surface and strain energy, provided the nucleus is

above critical size. When dislocations are present in

the matrix their elastic strain accommodates the strain

energy of the precipitate and the resistance to growth

of the precipitate is lowered.

The elastic strains caused by the growing precipitate

result from the misfit of the precipitate and matrix

lattices. The crystallographic data for the S’ phase was

first obtained using X-ray diffraction techniquest3) and

their values agree well with those obtained by

Weatherly and Nicholson(7) from transmission electron

diffraction patterns. From a comparison of these data

with the spacing of the corresponding aluminium planes

a misfit value 6 for the S’ phase in the aluminium

matrix may be calculated.

The values using the X-ray data are shown relative

to the aluminium in Table 2. The apparent similarity

between the lattice misfit in the a- and c-axes of the

precipitate suggests similar compressive stresses in the

matrix in these directions, and although the misfit

values obtained in this way give an indication of the

strain energy resulting from the formation of the pre-

cipitate they do not clearly indicate the relative

interfacial energies of the precipitate faces which con-

trol the precipitate shape.@) This is best studied by

comparing the number of aluminium lattice planes

TABLE 2

Interplanar spacing (A)

R’ S’

Electron

Misfit y0 relative to Al

8= ndal - ds’

ndnl x 100

Al X-ray datats’

- 4.06 2::: = 0.906

a = 4.00 b = 9.23

doi, = 0.906 c = 7.14

* +ve indicates compressive stress in aluminium matrix --ve indicates tensile stress in aluminium matrix

diffraction”’

a = 4.04 b = 9.25 c = 7.18

(n = 1) + 1.23* (n = 10) - 1.88 (n = 8) + 1.49

WILSON AND PARTRIDGE: S’

TABLE 3

Periodicity p = No. of Direction in S’ Al planes: 1 plane of S’

a 0.99

b 10.2

c 7.9

corresponding to one lattice spacing of the S’ precipi-

tate, i.e. the matching periodicity p.

The periodicity values along the a-, b- and c-axes of

the S’ precipitate are shown in Table 3. While coher-

ency is possible along the a-axis, the periodicity of

matching p. along the b- and c-axes is large and no easy

matching is possible : therefore the low misfit values 6

along the c-axis as shown in Table 2 is misleading.

The interfacial energy would be expected to increase

as the periodicity increases above unity and these

periodicities therefore permit the relative interfacial

energies of the faces of the S’ lath to be compared. The

S’ precipitate is drawn schematically in Fig. 10 and, as

would be expected from the above relationship, the S’

phase grows preferentially along the a-direction, the

direction with a periodicity of matching nearest unity.

The interfacial energy of plane B is proportional to the

periodicity along the a- and c-axes and since these are

the directions of best matching this plane will have the

minimum interfacial energy and therefore will be the

predominant plane of the precipitate. Similarly the A

plane will have the highest interfacial energy due to

- - --- 8’ PRECIPITATE PLANES

ALUMINIUM MATRIX PLANES

PRECIPITATES IN Al-Cu-Mg 1325

poor matching in both the b- and c-directions and wil

be relatively small in area. The observed morphology

is therefore consistent with the relative interfacial

energies of its faces deduced from a consideration of

the periodicities of matching of precipitate and matrix

planes at the precipitate faces.

The formation of corrugated precipitate sheets

Where a dislocation loop lies in a (110) plane with

a/2 [l lo] Burgers vector, a composite sheet of S’ pre-

cipitates may grow by the formation of laths on only

two (210) planes as is illustrated in Fig. 11. From the

stereographic projection (Fig. 9) it can be seen that

either of two pairs of (210) planes would satisfy this

condition, i.e. (210) and (120) or (210) and (i20). Which

of these two pairs is used is determined by two factors,

(a) the misfit vector of the precipitate and its relation-

ship to the Burgers vector of the dislocation and (b) the

line tension of the dislocation.

The misfit vector of a precipitate is the direction of

maximum misfit around the precipitate. Its import-

ance has been discussed by Kelly and Nichloson@) and

it has recently been calculated for the S’ precipitate in

an Al-Z.7 ‘A Cu-1.35 % Mg alloy by Weatherly and

Nicholson.(7) The misfits 6 of the S’ precipitate in the

principal directions have been calculated from the X-

ray results of perlitz and Westgrent3) and are shown in

Table 2. It can be seen that the misfit vector occurs in

the b-direction of the S’ precipitate producing a tensile

strain upon the aluminium matrix.

I IO.2 Al:

FIG. 10. Diagram showing the periodicity of aluminium and S’ planes at the surfaces of an S’ precipitate lath.

1326 ACTA NETALLURGICA, VOL.

GROWTH

13, 1965

PRECIPITATE PLANES

COMPOSITE PLANE

FIG. 11. The formation of a composite precipitate sheet on a (110) plane by the growth of precipitate lath on (120) and (210) planes in the [OOl] direction.

Since maximum relief of the misfit strains occurs

when the Burgers vector of the dislocation and the

misfit vector are parallel, the precipitate planes most

nearly satisfying this condition will be employed. Thus

the (210) and (120) planes would be used and not (2iO)

and (i20) planes. This condition is illustrated in Fig. 12

(a) ; in all the platelets or laths, the c-axis are lying in

the (210) and (120) planes and the b-axes are normal

to the platelet and inclined lS&0 from the [llO] direc-

tion of the Burgers vector of the dislocation. Com-

pressive stresses must exist over the platelet surface

with tensile stresses along the c-axes. The compressive

stresses could be reduced by vacancy loops lying paral-

lel to the precipitate with a dislocation line along the

precipitate and either alternately above and below the

precipitate sheet or all on one surface as shown in

Fig. 12(a).

An alternative precipitate morphology is that shown

in Fig. 12(b) in which alternate plates have the c- and

b-axes lying in the (210) and (120) planes: the tensile

FIG. 12. Possible models for the relief of coherency strains at the precipitate-matrix interface by formation of vacancv tvne dislocation loous (a) where b-axis of both {ZlO} lithi k-e approximatelynokal to the sheet and (b)

where the Z+ and c-axes alternate along the sheet.

and compressive components in these planes would

then tend to relieve each other over the sheet, and

alternate laths would exert tensile and compressive

stresses normal to the sheet surface. The latter stresses

could be relieved by vacancy loops above and below

those laths exerting compressive stresses on the matrix

as shown in Fig. 12(b).

Although the second precipitate arrangement can-

not be ruled out, the only evidence so far obtained

favours the arrangement shown in Fig. 12(a), and this

configuration conforms to the general rule that the

misfit vector prefers to lie parallel to the Burgers vec-

tor of the dislocation.

The line tension of the dislocation would also be

expected to influence the selection of (210) planes since

the nucleation and growth of precipitates on (210)

planes at a dislocation loop must deviate the loop from

its (1 lo} plane and increase its length. The dislocation

would prefer to be of minimum length due to its line

tension, and so the precipitate orientation lying close

to (110) would be chosen, i.e. again planes (210) and

(120) would be chosen for a dislocation loop on a (110)

plane and of a/2 (110) Burgers vector.

Rejinement of the S’ precipitate distribution

From the above discussion it would appear that the

S’ precipitate distribution might be refined in several

ways. Since the formation of the S’ precipitate is by

heterogeneous nucleation at dislocations, increasing

the dislocation density by prestrain prior to ageing

would be one method of refinement and since it would

increase the number of nucleating sites for a given

solute content, the final precipitates would be smaller.

This has indeed been observed and has been discussed

in detail elsewhere.c5) Silicon (0.25 %) may play a sim-

ilar role in the alloy. Figure 13(a) and (b) shows the

the effect of the addition of 0.25 % silicon upon the size

of S’ precipitates in an Al-2.5% Cu-1.2 % Mg alloy

aged 16 hr at 190°C ; the S’ precipitates are smaller

and more finely distributed in the silicon bearing alloy.

WILSON AND PARTRIDGE: S’

(b) FIG. 13. Effect of 0.25% Si on the size of S’ precipitates in the Al-2.5% C~-1.20~/~ Mgalloy S.T.,C.W.Q. andaged 16 hr at 19O’C (a) silicon free (b) 0.25% silicon. x 17,000

Measurement(l”) of the length of the S’ lath has shown

that silicon increases the incubation period for nucle-

ation and lowers the rate of growth of the precipitates.

Increasing the periodicity parameter, p, would affect

the rate of growth of the S’ and for a given change in S’

parameters this would be most marked in the a-direc-

tion. Segregation of silicon to the precipitatematrix

PRECIPITATES IN Al-Cu-Mg 1327

interface would also affect the interfacial energy and

affect the growth rateof the precipitate. Areductionof

the strain energy would make heterogeneous precipita-

tion at dislocations less favourable. However, the lattice

strain associated with the S’ precipitate is already low

and moreover it changes from a compressive to a ten-

sile strain with direction in the lath. Thus by reducing

the misfit in one direction one may of necessity in-

crease the misfit in another.

The two most profitable lines of approach therefore

might be to increase the number of nucleation sites

and to increase the periodicity, p, at the precipitate

matrix interface normal to the a-direction, thus reduc-

ing the rate of growth of the precipitates in this direc-

tion. CONCLUSIONS

1. The S’ precipitates grow as laths on (210) planes.

2. Composite sheets of S’ precipitates can be formed

on dislocations by laths lying on (210) planes having

common (100) growth direction.

3. The shape of the laths is consistent with the rela-

tive interfacial energies of the precipitate faces deduced

from the periodicity of matching of planes at the pre-

cipitate-matrix interfaces.

4. The precipitates grow at dislocations along direc-

tions and on planes consistent with the requirement

that their misfit vector must lie as close as possible to

the Burgers vector of the dislocation.

5. The line tension of the dislocation may also play a

part in determining which (210) planes are precipitate

growth planes.

6. Refinement of the precipitate distribution may be

obtained by increasing the number of heterogeneous

nucleating sites by prestrain, or by modification of the

interfacial energy by addition elements.

1.

2.

3. 4. 5.

6.

7.

a.

9. 10.

REFERENCES

H. PERLITZ and A. WESTGREN, Arkiv. Kemi. Min. Geol. B16 13 (1943). Yu. A. BAGARYATSKII. Zhm. Tekhn. Eiziki Fulmer Res. Inst. Trans. Nos. 11, 12, 55, 66. 1948. J. M. SILCOCK. J. Inst. Metals 89. 203 (19611. D. VAUGHAN, private oommu&&ion. ’ ’ R. N. WILSON and P. J. E. FORSYTH, J. Inst. Metals (In Press). J. D. EMBURY and R. B. NICHOLSON, Fifth Int. Congress fo? Electron Microscopy, Philadelphia. Academic Piess, New York p. Jl. 1962. G. C. WEATHERLY and R. B. NICHOLSON. The Structure and Properties of Al-&-Mg Alloys. Prog. Rep. on M.O.A. No. PD/29/025 1963. A. KELLY and R. B. NICHOLSON, Progr. Mat. Sci. 10, 151 (1963). D. W. PASHLEY. Phil. Mug. 5, 173 (1956). R. N. WILSON, D. M. MOORE and P. J. E. FORSYTH, to be published.