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Heterogeneous nucleation of γ′ in Al-Ag and Al-Ag (Cd or Cu) alloys
HETEROGENEOUS
NUCLEATION
OF y’ IN Al-Ag
AND
Al-A9
(Cd
OR Cu) ALLOYS*
and G. S. ANSELL?
D. E. PASSOJAt$
Transmission electron microscopy studies of heterogeneous nucleation of the m&a&able y’ phase were made in binary and ternary alloys of Al-6 at. % Ag, Al-4.8 at. % Q-0.15 at. % Co! and Al-4.8 at. ‘A Ag0.2 at. % Cu. Heterogeneous nucleation of y’ was found to be associated with four types of dislocations regardless of the quenching treatment or alloy additions which were used. Nucleation was found to occur on: (a) jogged dislocations; (b) helical dislocations; (c) dislocation loops; (d) heterogeneousclimb sources. Only one nucleation mechanism was found to be unique to the ternary alloys : that of nucleation on heterogeneous climb sources. NUCLEATION
HETEROGENE DE y’ DANS LES ALLIAGES Al-A&Cd OR Cu)
Al-Ag
ET
Les &udes au microscope Blectroniquede la nucleation h&&rog&nede la phase m&stable y’ ont BtB effect&es pour les alliages binaires et ternaires de Al-5 % at. Ag, Al-4,8% at. Ag-0,16 % at. Cd et Al4,s % at. Ag-0,2 x at. Cu. Les auteurs ont trouvb que lanuclbation h&Brog&e de y’ est associbe B quatre types de dislocations, quels que soient le traitement de trempe ou les compositions des alliages. La nucleation se produit sur: (a) les dislocations portant des crans; (b) les dislocations en hblices; (c) les Pour l’alliage ternaire les auteurs boucles de dislocations; (d) 1es sources de montbe h&rog&nes. ont trouvb seulement un m&anisme de nucleation unique : celui de la nucleation sur les souroesde montbe h&t&og8nes. HETEROGENE
y’-KEIMBILDUNG
IN Al-Ag-
und Al-Ag-(Cd
ODER
Cu)-LEGIERUNGEN
Elektronenmikroskopische Untersuohungen der heterogenen Keimbildung der metastabilen y/-Phase wurden an den bin&en und tern&en Legierungen Al-5 At. % Ag, Al-+3 At. % Ag-0,15 At. ok Cd und Al-4,8 At. ‘A Ag-0,2 At. ‘A Cu durchgefiihrt. Die heterogene y’-Keimbildung htigt unabhiingig von der Abschreckbehandlung oder von den Zulegierungen mit vier Versetzungstypen zusammen. Keimbildung fand statt an: (a) Versetzungen mit Versetzungsspriingen; (b) spiralf&migen Versetzungen; (c) Versetzungsringen; (d) heterogenen Kletterquellen. Nur ein Keimbildungsmechanismus wurde gefunden, der all&n in der tern&en Legierung auftritt : Keimbildung an heterogenen Kletterquellen.
INTRODUCTION
Decomposition of metastable aluminum-rich aluminum-silver alloys occurs by the formation of both spherical silver-rich Guinier-Preston zones and the hexagonal close-packed intermediate y’ phase.(l) Precipitation of G. P. zones in a 5 at.% alloy, for example, occurs immediately upon quenching the alloy from above the coherent solvus temperature(2s3) (Fig. 1) and also at higher temperatures which are below the coherent solvus. Aging the alloy after quenching results in the heterogeneous nucleation of y’ platelets on dislocations contained in the metallic crystal. Growth of y’ proceeds by dissolution of the G. P. zones in the immediate vicinity of the growing y’ platelets.(4) The AI-Ag system is a particularly interesting one for the study of the heterogeneous nucleation of precipitates for several reasons: (a) there is a small difference in size between the aluminum and silver atoms in the binary aluminumrich allays,(5) resulting in a minimal strain-nergy interaction between silver atoms, dislocations and vacancies; * Received June 7, 1968; revised April 5, 1971. t Materials Division, Rensselaer Polytechnic Institute, Troy, New York. ow at: Chase Brass & Copper Co., 11000 Cedar Ave., ClJvzand, Ohio 44106. ACTA METALLURGICA,
VOL.
19, NOVEMBER
1971
(b) there is an extremely small vacancy-solute binding energy’s) which is probably electronic in origin;“) (c) the crystallographic nature of the y’ nucleation (f.c.c. ---f h.c.p.) can be easily described. Thus the aluminum-rich aluminum-silver alloys are particularly useful for the study of heterogeneous nucleation of precipitates since the factors governing nucleation appear to be limited. Several investigators (*-lo) have studied, by means of transmission electron microscopy, precipitation of the y’ phase in the aluminum-rich aluminum-silver system. It has been found that platelets of the y’ nucleate heterogeneously on dislocations of all types, are semi-coherent, and have a (OOOl),, (1(ill),, [llZO],, 11[liO], orientation relationship. Several y’ nucleation mechanisms have been observed by the various investigators. Nicholson and Nutting’ll) found that in water quenched samples, y’ nucleates at helical dislocations and at isolated stacking faults associated with Frank sessile dislocation loops formed during quenching from the solution treatment temperature. Hren and Thomas(12) studied the precipitation of y’ in thin foils quenched and aged in situ in the electron microscope. These investigators found that y’ nucleates at dislocations of all types; however, helices and Frank sessile dislocation loops
1253
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700 660 600
Atomic
percent
FIG. 1. The aluminum-silver
Silver
phase diagram.
were most effective as nucleation sites for the y’ in freshly quenched specimens. Nemoto and Koda(13) utilized specimens that had been first deformed in tension, rapidly heated to 5OO”C, and held at that temperature for three minutes. By using cooling rates of 40 and lO”C/min, nucleation of y’ in the as-quenched foils was obtained. It was found that dislocations in subboundary networks formed jogs and the nucleation of y’ occurred only on those dislocation segments lying on (111) planes. It is clear that heterogeneous nucleation of y’ has been observed to occur by several mechanisms in the Al-AS system. However, the details of some of the mechanisms and how these mechanisms depend upon quench rate and alloy content are not well established. The purpose of this investigation was to study, by means of transmission electron microscopy, the heterogeneous nucleation and growth of the y’ precipitate in aluminum-rich Al-Ag alloys. It was desirable to place emphasis upon understanding how the spectrum of nucleation processes depended upon To acvacancy supersaturation and migration. complish this goal, the relative amounts of vacancy supersaturation and migration were controlled through ; (a) variation in quench rates in binary Al-Ag alloys; and (b) additions of third elements to the binary alloy system. EXPERIMENTAL
DETAILS
Alloy preparation
Alloys of Al-5 at. % Ag were prepared from 99.999 % Al and 99.99 % Ag. The ingots were prepared by air
VOL.
19,
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melting and vacuum induction melting, and chill casting the melts into heavy aluminum molds. The cast ingots were then homogenized at 500% for 2 weeks in air. After the homogenizing treatment, the ingots were water-quenched and cold-rolled into 0.006 in. thick foils. No substantial differences in aging response could be detected between the two melting techniques used. Alloys of A14.8 at.% Ag-0.2 at.% Cu and A14.8 at. % Ag-0.15 at. % Cd were prepared from 99.999 y0 Al, 99.99 % Cd, 99.999 % Cu and 99.999 % Ag. Cylindrical ingots were prepared by air melting and then chill casting the melts into heavy aluminum molds. The ingots were homogenized at 550°C for 2 weeks in air. After the homogenizing heat treatment, the ingots were water-quenched and then cold-rolled into 0.006 in. thick foils.
Solution annealing heat treatment and quenching
A portion of the Al-5 at.% Ag foils, wrapped in groups of twenty in aluminum foil to minimize oxidation, were solution annealed in air at 550 & 5°C for 24 hr, and then quenched into ice water at 3°C. The remaining samples, wrapped in groups of ten in aluminum foil, were solution annealed in air at 550 -& 5°C for 24 hr and then air-cooled. The cooling rate for the water-quenched samples was approximately 1500”C/min as compared to approximately 40”C/min for the air-quenched samples. The structures of the air-cooled specimens were studied as-cooled, and were given no subsequent heat treatment. The structures of the water-quenched samples were studied both in the as-quenched condition and after various aging heat treatments.
Aging
All the water-quenched samples were aged in either air or silicone oil at 160 f 5°C for various time intervals up to 1000 hr. The specimens aged in air were wrapped in lots of ten in aluminum foil to minimize oxidation. No difference in the nucleation and growth response of the y’ precipitate was discernible between the two types of aging treatments.
TABLE 1. Alloys and heat treat,ments for samples listed in the text Air quenched at 550°C Sample Sample Sample Sample
No. No. No. No.
1 2 3 4
Yes NO No No
H,O quenched at 550°C v ;:* Yes Yes
Aging time at 160°C (hr) 20: 48 48
At. y0 solute Ag
i.8 5 4.8
Cd
i.15 0 0
CU
:: 0 0.2
PASSOJA
Specimen
AND
ANSELL:
NUCLEATION
OF
y’
IN
Al-Ag
ALLOYS
1255
prepamtion
Foils suitable for transmission electron microscopy were prepared from the aa-quenched and aged samples by first electro-thinning in a modified Lenoir solution at 70-80°C followed by dipping the foils into a hot (7O’C) dilute solution of chromic acid, phosphoric acid and Ha0 in order to remove surface films. A Hitachi HU 125 electron microscope with a 130” tilting stage and ele~troma~etie deflection coils for dark field microscopy was used to observe the specimens.
In addition to those previously reported, several other nucleation mechanisms for the precipitation of y’ were observed in the Al-5 at. % Ag and Al4.8 at. % Ag, 0.2 at.% Cu or 0.15 at.% Cd alloys. Of these, only one type of y’ nucleation was unique in that it was observed to occur in the ternary alloys and not in the binary alloys. The other nucleation mechanisms were found to be present in all the specimens. However, the relative frequencies with which each nucleation mechanism was observed was found to be dependent upon quench rate and alloy content Since most of the nucleation mechanisms were common to all of the specimens examined, the observations of the heterogeneous nucleation of y’ are presented by utilizing the best examples of nucleation mechanisms observed from the various alloys and the corresponding heat treatment. A Thomson tetrahedron is superimposed on some of the mi~rog~phs to facilitate visualization of the geometric relationships. (a) Nucleation on jogged dislocations--stair rod formation. Figures 2 and 3 are examples of y’ nucleation frequently observed in air-quenched samples of Al-5 at. % Ag (sample No. I). Observation of these
PIG. 2. Al-5 et. % Ag air-quenched from 550°C. Climbed extended jogs containing stair rods at A-H. 8
FIG. 3. Al-5 at. % Ag air-quenched from 550°C. Climber extended superjogs. Obtuse and acute at A-2 and A-l, respectively.
st~ctures in the water-quenched and aged alloys was relatively infrequent. This precipitate morphology is similar to that observed by Hren and Thomas(14) in slowly quenched foils studied in &u in the electron microscope. Figure 2 is a ribbon-like structure formed from stacking faults on two sets of (111) planes. It is believed that this structure resulted from the absorption of silver during the queneh, on a jogged edge dislocation, thus causing dissociation of the dislocation. Since the jogged segments are extended, this structure must have stair rod dislocations at the intersections between the jogged sections. The reactions between the Shockley partials which give rise to the stair rods are either #ll] or
+- $[lE]
$[112] -I- ,[llE]
+ +[I101
(I)
-+ i[llO]
(2)
Reaction (1) has the greatest elastic energy decrease if it proceeds to the right, hence it is inferred that an a/G[iiO] stair rod forms at the junctions of the acute angles of intersections A-F. Figure 3 (at A) shows a section of a stacking fault which has climbed out of the plane of stacking faults B and C. The A-l segment has formed acute jogs at the junction of the faults between A and C and is constricted. The other segment A-2 has formed a jog with an obtuse angle and is extended. This dissoeiation reaction could have occurred by means of vacancy condensation and climb of an extended edge dislocation. Barnes(15) has suggested that an extended edge dislocation can climb by the addition of vacancies to the bounding partial dislocations. The dislocations climb by vacancy condensation, and the fault sequence is restored by a shift of a/12[112]; thus the stacking sequence is maintained on alternate (111) planes. It is not likely that this type of climb process
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can occur extensively in regions where the G. P. zone density is high, however, since the fault would be pinned by the zones in the immediate vicinity of the fault. One must conclude therefore that dislocation jogging occurs more readily at high temperatures, above the coherent solvus temperature, giving rise to a dislocation containing large super jogs. Subsequent precipitation on the dislocation causes a decrease of the stacking fault energy and extension of the partial dislocations. In Fig. 3 the fact that the obtuse bend A-2 is extended whereas the acute bend A-l is not completely extended, is related to the length of the jog as well as the type of stair rods formed at each junction, If a stair rod of Burgers’ vector $lOOl] forms at an obtuse jog, and a +[IlO] stair rod forms at an acute jog, then the energies of the stair rods are in the ratio driving of ~~~~c~~~~~cc~ = 2I1. The therm~~amic force for the f.c.c. -+ h.c.p. transformation is apparently not great enough to overcome the line tension forces of the Shockley partials and hence a constricted junction formed at A-l. Figure 4 (sample No. 2) shows a dark field micrograph of this type y’ nucleation imaged from the (Iii) plane. The displacement fringes from the precipitate are clearly visible on three sets of (111) planes. The fringe between A-B is “invisible” as a result of having overlapping stacking faults, giving rise to @ * .i? = nb (where n = 0, I, 2 . . .).(lQ Between B and C two displacement fringes are visible in the (111) reflection inasmuch as the total displacement from the stacking faults gives Q - R # nb, the b is different for the two faults shown at B-C. These observations suggest that the number of different orientations of y’ which nucleate from a single edge dislocation containing superjogs is de-
FIQ. 4. Aluminum--4.8 at. %Ag-0.16 at.%Cd H,Oquenched from 550°C and aged at 16O’C for 200 hr. Dark field photomicrograph showing displacement fringes and G.P. zones, 9 = @Ii).
VOL.
19, 1971
FIQ. 5. Two
types of y’ nucleation(stair).
pendent upon the nature of the stair rods which form at the junctions of the faults. The stability of stair rod dislocations which form at the junctions of extended dislocations lying on different slip planes is dependent upon the angle between the Burgers’ vector and the line of the initial unjogged dislocation. If the stair rods are unstable relative to a constricted node, only one orientation of y’ will form from one dislocation line (Fig. 5). The observation of jogged-extended dislocations as heterogeneous nucleation sites for y’ in these alloys suggests that the jog-type of nucleation as observed by Nemoto and Koda(l7) is operational in all of these alloys. The number of jogged dislocations observed in water-quenched specimens, were far fewer than those observed in the air-quenched specimens, however. This indicates that jogging occurs more readily during slow cooling from the solution treatment temperature probably due to the more extensive vacancy migration and annihilation occurring above the coherent solvus tern~rature. In rapidly quenched specimens there is extensive dislocation loop formation and helical dislocation formation; i.e. structures resulting from high vacancy supersaturation. The extent of the stair nucleation process is, therefore, strongly dependent upon the quench rate as well as the jog formation processes which occurs at high temperatures above the coherent solvus temperature. (b) ~~c~u~~on on helicaE d~8lo~~~n~. Nucleation of y’ on helical dislocations was frequently observed in the water-quenched and aged binary and ternary alloys. heterogeneous nucleation of y’ on helical dislocations has been observed previously in an investigation by Nicholson and Nutting;(l*) however, some of the details of the decomposiition of the helix and nucleation of y’ on the helix were found to be somewhat different than those reported previously. The decomposition of a helical dislocation proceeds first, by the migration of silver t,o the dislocation. It is di~euIt to ascertain precisely when y’ nucleates on
PASSOJA
XUCLEATION
ANSELL:
AND
OF
y’
IN
,41-Ag
1257
ALLOYS
FIG. 6. Aluminum-5 at. % Ag H,O-quenched from 550°C and aged at 160°C for 48 hr. the helix since the solute
migration
and nucleation
processes appears as one continuous well defined
steps.
Figures
show the decomposition sive segments
6 and 7 (sample No. 3)
of the helix when the succes-
of the dislocation
Decomposition
of
solute build-up
in the vicinity
segments;
the
process without
helix
are not co-planar.
occurs
gradually
of curved
this is usually indicated
by
dislocation
by strong absorp-
tion contrast in the bright field image, resulting from a localized increase in silver concentration At longer aging times the absorption stronger
close to the dislocation
assumes
polygonal
shape
(Fig.
shape which the y’ ultimately thin
hexagonal
shaped
line can be seen surrounding
contrast becomes line, and the line
7).
The
polygonal
assumes is that of a
platelet
bounded by (110) directions
(e.g. Fig. 7).
whose
edges
(the faint hexagonal out-
B in Fig. 7).
When the segments of the helix decompose platelets,
FIG. 550°C
decomposition
are
occurs
by
the
into y’
continuous
7. Aluminum-5 at. % Ag. H,O-quenched and aged at 160°C for 48 hr. Polygonal segments (A and B).
from shaped
FIG. 8. Aluminum-4.8 at. % Ag-0.2 at. % Cu. H,Oquenched and aged for 48 hr at 160°C. Loop decomposition on helix.
localized build-up of solute within the concave portions of dislocation dislocation
Often the helical
segments on the helix. line
decomposes
into
dislocation
which are in the same plane (Fig. 8, sampleNo. when co-planar composition,
dislocation
loops 4); and
loops form from the de-
solute build-up
occurs gradually
the loop, resulting in loop expansion
within
and loop colaes-
cence (Fig. 9, sample No. 4). These types of decomposition be independent frequently
content
for they
could
be observed in both the binary and ternary
water-quenched
FIG. 9.
of impurity
processes appeared to
and aged samples.
Aluminum-4.8 at. %Ag-0.2 at. %Cu. H,Oquenched and aged for 48 hr at 160°C. Loop coalescence on helix.
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FIG. 12. The dislocation node analysis of photomicrograph lo at D. H,OFlo. 10. Aluminum-4.8 at. %Ag-0.15 at. %Cd. quenched from 550°C~aged 200 hr at lBO”C, helix decomposition, 9 = (200)
Figures 10 and 11 are bright fieId and dark fieId photographs (sample No. 2) of y’ which had nucleated on a helix. The nature of the ~splacement vector ofthe precipitate can be obtained from these two pictures by methods outlined by Hashimoto et aZ.(19) and Amelinckx.(20) The phase factor a’ w&s found to be 2~13 sinoe the first displacement fringe of the precipitate at the top of the foil was bright (the top and bottom of the foil are marked T and B, respectively, Fig. 11). Utilizing the method outlined by Amelinckx,(21) it was found that the fault was extrinsie with a shear vector of *(iii). It can be seen that in the region G of Fig. 10 no fringes are observed; furthermore, the dislocation which separates the fringed section from the unfringed section is not visibie, By utilization of the phase factor, cc, dislocation node equations and contrast criteria, it was possible to infer what dislocat,ions are present at the boundary of the precipitate.
Figure 12 shows the dislocation node analysis which was performed on region D of Fig. 10. The nature of the dislocations present at the precipitate boundaries are important with regard to both nucleation and growth mechanisms of the precipitate. I?rom the analysis shown in Fig. 12 it was found that the dislocation reaction at D could be described by the reaction:
giolj 3 gill] + -@I~
(3)
which gives a Shockley partial plus a Frank partial; the stacking fault is extrinsic (Fig. 13) and is similar to a reaction observed by Silcock(2s) in stainless steels. The observation of precipitates with dispia~ement vectors which resemble shear vectors equivalent to that of the extrinsic stacking fault is important since
\
b=-&[iOi]
-------7 trrtt
{ /’
Fro. Il. Dark field pho~micro~&ph of Pig. 10, r* = %/s/3. The top and bottom of the foil are marked 2” and B’, respectively.
Fro. 13. Decomposition of a hefical segment into tEn extrinsic frank partial.
PASSOJA
AND ANSELL:
NUCLEATION
the nucleation mechanism of 7’ depends on the par-
ticular type of stacking fault which is formed. Nucleation of y’ on an intrinsic fault has been suggested(z2) to occur by means of a Suzuki mechanism i.e. the local excess of silver in the vicinity of a dislocation line locally lowers the stacking fault energy, resulting in the extension of the dislocation; the stacking fault then forms the heterogeneous nucleus for nucleation of y’. On the basis of the observations presented in this section an alternate mechanism can be proposed for the nucleation of y’ on helices. During quenching, vacancies are eliminated in the vicinity of the dislocation line due to formation of the helix. Due to the interaction between the vacancies and the silver atoms, the silver concentration increases in the vicinity of the dislocation as a result of the loss of vacancies to form the helix. By increasing the local silver concentration the stacking fault energy around the helix is lowered resulting in the splitting of the dislocation into partials. The dislocation cannot split entirely along its length, however, since it no longer lies in a single plane. Solute thus is attracted to the various extended segments via a Suzuki mechanism. An extrinsic fault may appear if there is no other mechanism of generating a new layer of precipitate. At the aging temperature the composition of the alloys are such that the free energy vs. composition curve has a negative curvature,(za) and thus uphill diffusion can occur. The local extended extrinsic faults can act as continuous sources of vacancies to assist uphill diffusion by operating as climb sources. Thus the precipitate can act as a Bardeen-Heering climb source resulting in the production of a silver-rich extra half plane of atoms bounded by a Frank partials whose shear vector is equivalent to the extrinsic displacement vector of the y’ precipitate. An extrinsic Frank dislocation can grow on (111) planes by dissociation of the helix and expansion by an outward climb of the segments of the Frank partial; the extrinsic Frank partial can then grow by vacancy emission (Fig. 13). This growth mechanism will operate until a new intrinsic fault forms at the precipitate matrix interface by shearing. New precipitate layers can thus be built up by an intrinsic --f extrinsic growth mechanism. Figure 10 at E shows a Frank partial which overlaps segment C on a parallel (111) plane. The phase factor for both segments A and C is 2rr/3, resulting in a phase factor in the overlapping region at E equal to -2rr/3. This results in a reversal of the fringe contrast observed at E. The dislocation bounding the fault A-E has sections which lie parallel to the (110) directions in the (111) plane. The contrast in the over-
OF
y’
IN
Al-Ag
1259
ALLOYS
lapped section at A shows no contrast, since there are three overlapping faults in this region, giving rise to a phase factor equal to zero. It may be concluded, therefore, that multiple extrinsic faults bounded by Frank partials are present on parallel (111) planes. Extrinsic precipitate displacement vectors were not present in every nucleation event which was observed. It is likely that both extrinsic and intrinsic faults are present depending upon the particular conditions which obtain at the time the observation is made. (c) y’ nucleation on dislocation loops. In the waterquenched binary and ternary alloys, nucleation of y’ on prismatic dislocation loops and Frank sessile dislocation loops was by far the most common y’ nucleation mode. The asquenched dislocation loop density of the alloys was in the order of Al-Ag4d > Al-Ag-Cu > Al-AS, whereas the as-quenched loop size was in the reverse order. These trends are probably due to two influences; the ease of dislocation loop nucleation and the vacancy supersaturation possibly as it is related to the vacancy-solute binding energy. If the addition of an impurity atom increases the heterogeneous nucleation of dislocation 10ops(~4) and if, in addition, the solute atoms (Cu and Cd) have a high vacancy-solute binding energy then the dislocation loop density would be high but the loop size would be small. Recent work in ternary Al-Ag-Zn alloys(25) indicates that the dislocation loop density can be strongly influenced by the nature of ternary G. P. zones. Both the heterogeneous nucleation of dislocation loops as well as the vacancy migration in the water-quenched alloys used in this study could be a result of the presence of numbers of ternary G. P. zones having anisotropic strain fields and trapping vacancies.(26) Heterogeneous nucleation of the dislocation loops could also be facilitated by the assistance of the strain field surrounding the G. P, zones. Hence the observed trends of the as-quenched dislocation loop size and density can be explained on a qualitative basis by consideration of the impurities’ independent influences on nucleation and growth of the dislocation loops. (d) Nucleation on heterogeneous climb sources. There was one mode of y’ nucleation which was unique to the ternary Al-AgXu and Al-Ag+d alloys. Figure 14 shows y’ precipitation on a heterogeneous climb source found in the Al-Ag-Cd- (sample No. 2) specimens; Fig. 15 shows a similar precipitate morphology found in Al-AgXu (sample No. 1). In both cases the precipitation associated
had occurred on dislocations with dark impurity
particles
15(b)]. This type of y’ nucleation in the binary Al-Ag alloys.
which were [Figs.
14(a),
was never observed
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In both the binary and ternary alloys heterogeneous nucleation of y f was found to occur on: (a) (b) (c) (d)
FIG. 14. Aluminum-4.8 at. %Ag-OS15 at. %Cd. quenched from 550°C-aged at 160°C for Xuoleation of heterogeneous climb source.
HaO36 hr.
Fra. 16. Atuminum-4.8 at. %Ag-0.2 at. %Cu. quenched from 550”C-aged at 160°C for Nuoleabion on heterogeneous climb source.
H2048 hr.
climb sources have been previously observed in A1-Mgtz7) alloys; decoration of heterogeneous climb sources by precipitates has also been reported for the Al-% system.(2a) In Al-AS alloys, it appears that y’ nucleation will occur on heterogeneous climb sources only if the insoluble impurity particles are present to initiate the heterogeneous climb source. ~e~~~~~~~e~~~
CONCLUSIONS
Decomposition of metastable binary AI-AS and ternary Al-Ag-Cd or Cu alloys occurs by the formation of G. P. zones and nucleation of y’ precipitate phase.
jogged dislocations; helical dislocations; dislocation loops; heterogeneous climb sources.
Only one nucleation mechanism was found to be unique to t.he ternary alloys-nucleation on dislocation climb sources associated with impurity particles. The frequency with which a nucleation event was observed on these dislocations depended upon the relative amounts of dislocation superjog formation, vacancy supersaturation and migration, and was influenced by quench rate and alloy impurity additions. The nucleation controlled by a stair rod process was found to operate by the initial formation of superjogs on edge dislocations by vacancy absorption above the coherent solvus temperature. Nucleation of y’ in one or two orientations on the jogged dislocations depended upon the relative stability with respect to constricted nodes of the extended-jogged segments of the dislocations. A discrete nucleation step for y’ was never observed for nucleation on helices; nucleation of y’ was found to occur by the continuous local solute build-up in the vicinity of the helical dislocations. This resulted in the development of segments of dislocation lines having a polygonal shape. Some of the displacement vectors of these polygonal dislocation segments were found to have extrinsic displacement vectors of &[i%iJ. A dislocation pole mechanism which can act as a continuous source of vacancies and thus generate new layers of y’ on the faces of the y’ platelet, is characterized by the reactions
gior]-+ griTi + gi2il and could control the nucleation process on helices. This disfocation dissociation (a Bardeen-Heering climb source) can act to generate vacancies end new layers of the precipitate by means of climb and an intrinsic --+ extrinsic -+ intrinsic stacking fault growth mechanism. ACKNOWLEDGMENTS
The work reported in this paper was supported by both the National Aeronautics and Space Administration and the office of Naval Research. The studies were conducted in the NASA Interdiscip~a~ Materials Research Center at, Rensselaer Polytechnic Institute, Troy, New York.
PASSOJA
AND
ANSELL:
NUCLEATION
REFERENCES 1. C. S. BARRETT, Structure of Metals, p. 541. McGraw-Hill (1952). 2. M. HILLERT, B. L. AVERBACH and M. COHEN, Acta Met. 4, 31 (1956). 3. R. BAIJR and V. GEROLD, Acta Met. 10, 637 (1962). 4. R. B. NICHOLSON and J. NUTTINQ, Acta Met. 9,332 (1961). of Lattice Spacings and 5. A. PEARSON, A Handbook Structurea of Metals and A&ye. Pergamon Press (1958). 6. D. R. BEAMAN, R. W. BALLUFFI and R. 0. SIMMONS, Phys. Rev. 184, 532A (1964). 7. A. BLADIN and J. L. DEPLANTE, MetaZZic SoZid SoZutiolas, edited by J. FREIDEL and A. GUINIER. Benjamin (1963). R. B. NICHOLSON and J. NUTTINO, Acta Met. 0,332 (1961). :: J. A. H&EN and G. TRODIAS, 2lran.s Am. Inat. Min. Engv8 237, 308 (1963). 10. M. NEMOTO and S. KODA, J. Inst. Metals 98,164 (1964-66). 11. R. B. NICHOLSON and J. NUTTINQ, Acta Met. Q, 332 (1961). 12. J. A. HREN and G. THOMAS, Trans. Am. Inst. Min. Engrs
227, 308 (1963). M. NEMOTO and S. KODA, J. Inst. Metal8 93, 164. :43: J. A. HREN and G. THOMAS, Trans. Am. Inst. Min. Engrs 227, 308 (1963). 15. R. S. BARNES, Acta Met. 2, 380 (1954). 16. HIRSH et al., Electron Microscopy of Thin Crystals, p. 339. Butterworths
(1965).
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ALLOYS
1261
(1964-66). 17. M. NE~OTO and S. KODA, J. Inst. Metal8 98,164 18. R. B. NICHOLSON and J. NUTTING, Acta Met. 9,332 (1961). 19. H. HASHI~UOTO,A. HOWIE and M. J. WEHLAN, Proc. R. Sot. A269, 80 (1962). 20. S. AMELINCKX, The Direct Observation of DieZocationa, Solid State Phyaica Se&e, edited by F. SEITZ and D. TURNBULL, Vol. 6. Academic Press (1964). 21. S. AMELINCKX, The Direct Observation of Dislocations, Solid State Physics Selies, edited by E. SEITZ and D. TURNBULL, p. 459. Academic Press (1964). 22. R. NICHOLSON and J. NUTTINO, Acta Met. 9, 332 (1961). 23. M. HILLERT, B. L. AVERBACK and M. COHEN, Acta Met. 4, 31 (1956). 24. T. L. DAVIS and J. P. HIRTH, J. appl. Phys. 37, 2112 (1966). 25. D. E. PASSOJA and R. B. NICHOLSON, University of Manchester Report. 26. D. E. PASSOJA, S. POPOVIC and P. BARRAND, to be published. 27. R. B. NICHOLSON, Metallic Solid Solutiorw, edited by J. FRIEDEL and A. GUINIER, PXLIV-10. Benjamin (1963). 28. M. 0. SPEIDEL, Boeing Scientific Research Laboratories, Solid State Physics Laboratory, Review July-December (1966).
29. J. SILCOCK and W. (1964).
J. TUNSTALL, Phil.
Mag.
10,
361
NUCLEATION
OF y’ IN Al-Ag
AND
Al-A9
(Cd
OR Cu) ALLOYS*
and G. S. ANSELL?
D. E. PASSOJAt$
Transmission electron microscopy studies of heterogeneous nucleation of the m&a&able y’ phase were made in binary and ternary alloys of Al-6 at. % Ag, Al-4.8 at. % Q-0.15 at. % Co! and Al-4.8 at. ‘A Ag0.2 at. % Cu. Heterogeneous nucleation of y’ was found to be associated with four types of dislocations regardless of the quenching treatment or alloy additions which were used. Nucleation was found to occur on: (a) jogged dislocations; (b) helical dislocations; (c) dislocation loops; (d) heterogeneousclimb sources. Only one nucleation mechanism was found to be unique to the ternary alloys : that of nucleation on heterogeneous climb sources. NUCLEATION
HETEROGENE DE y’ DANS LES ALLIAGES Al-A&Cd OR Cu)
Al-Ag
ET
Les &udes au microscope Blectroniquede la nucleation h&&rog&nede la phase m&stable y’ ont BtB effect&es pour les alliages binaires et ternaires de Al-5 % at. Ag, Al-4,8% at. Ag-0,16 % at. Cd et Al4,s % at. Ag-0,2 x at. Cu. Les auteurs ont trouvb que lanuclbation h&Brog&e de y’ est associbe B quatre types de dislocations, quels que soient le traitement de trempe ou les compositions des alliages. La nucleation se produit sur: (a) les dislocations portant des crans; (b) les dislocations en hblices; (c) les Pour l’alliage ternaire les auteurs boucles de dislocations; (d) 1es sources de montbe h&rog&nes. ont trouvb seulement un m&anisme de nucleation unique : celui de la nucleation sur les souroesde montbe h&t&og8nes. HETEROGENE
y’-KEIMBILDUNG
IN Al-Ag-
und Al-Ag-(Cd
ODER
Cu)-LEGIERUNGEN
Elektronenmikroskopische Untersuohungen der heterogenen Keimbildung der metastabilen y/-Phase wurden an den bin&en und tern&en Legierungen Al-5 At. % Ag, Al-+3 At. % Ag-0,15 At. ok Cd und Al-4,8 At. ‘A Ag-0,2 At. ‘A Cu durchgefiihrt. Die heterogene y’-Keimbildung htigt unabhiingig von der Abschreckbehandlung oder von den Zulegierungen mit vier Versetzungstypen zusammen. Keimbildung fand statt an: (a) Versetzungen mit Versetzungsspriingen; (b) spiralf&migen Versetzungen; (c) Versetzungsringen; (d) heterogenen Kletterquellen. Nur ein Keimbildungsmechanismus wurde gefunden, der all&n in der tern&en Legierung auftritt : Keimbildung an heterogenen Kletterquellen.
INTRODUCTION
Decomposition of metastable aluminum-rich aluminum-silver alloys occurs by the formation of both spherical silver-rich Guinier-Preston zones and the hexagonal close-packed intermediate y’ phase.(l) Precipitation of G. P. zones in a 5 at.% alloy, for example, occurs immediately upon quenching the alloy from above the coherent solvus temperature(2s3) (Fig. 1) and also at higher temperatures which are below the coherent solvus. Aging the alloy after quenching results in the heterogeneous nucleation of y’ platelets on dislocations contained in the metallic crystal. Growth of y’ proceeds by dissolution of the G. P. zones in the immediate vicinity of the growing y’ platelets.(4) The AI-Ag system is a particularly interesting one for the study of the heterogeneous nucleation of precipitates for several reasons: (a) there is a small difference in size between the aluminum and silver atoms in the binary aluminumrich allays,(5) resulting in a minimal strain-nergy interaction between silver atoms, dislocations and vacancies; * Received June 7, 1968; revised April 5, 1971. t Materials Division, Rensselaer Polytechnic Institute, Troy, New York. ow at: Chase Brass & Copper Co., 11000 Cedar Ave., ClJvzand, Ohio 44106. ACTA METALLURGICA,
VOL.
19, NOVEMBER
1971
(b) there is an extremely small vacancy-solute binding energy’s) which is probably electronic in origin;“) (c) the crystallographic nature of the y’ nucleation (f.c.c. ---f h.c.p.) can be easily described. Thus the aluminum-rich aluminum-silver alloys are particularly useful for the study of heterogeneous nucleation of precipitates since the factors governing nucleation appear to be limited. Several investigators (*-lo) have studied, by means of transmission electron microscopy, precipitation of the y’ phase in the aluminum-rich aluminum-silver system. It has been found that platelets of the y’ nucleate heterogeneously on dislocations of all types, are semi-coherent, and have a (OOOl),, (1(ill),, [llZO],, 11[liO], orientation relationship. Several y’ nucleation mechanisms have been observed by the various investigators. Nicholson and Nutting’ll) found that in water quenched samples, y’ nucleates at helical dislocations and at isolated stacking faults associated with Frank sessile dislocation loops formed during quenching from the solution treatment temperature. Hren and Thomas(12) studied the precipitation of y’ in thin foils quenched and aged in situ in the electron microscope. These investigators found that y’ nucleates at dislocations of all types; however, helices and Frank sessile dislocation loops
1253
ACTA
1254
METALLURGICA,
700 660 600
Atomic
percent
FIG. 1. The aluminum-silver
Silver
phase diagram.
were most effective as nucleation sites for the y’ in freshly quenched specimens. Nemoto and Koda(13) utilized specimens that had been first deformed in tension, rapidly heated to 5OO”C, and held at that temperature for three minutes. By using cooling rates of 40 and lO”C/min, nucleation of y’ in the as-quenched foils was obtained. It was found that dislocations in subboundary networks formed jogs and the nucleation of y’ occurred only on those dislocation segments lying on (111) planes. It is clear that heterogeneous nucleation of y’ has been observed to occur by several mechanisms in the Al-AS system. However, the details of some of the mechanisms and how these mechanisms depend upon quench rate and alloy content are not well established. The purpose of this investigation was to study, by means of transmission electron microscopy, the heterogeneous nucleation and growth of the y’ precipitate in aluminum-rich Al-Ag alloys. It was desirable to place emphasis upon understanding how the spectrum of nucleation processes depended upon To acvacancy supersaturation and migration. complish this goal, the relative amounts of vacancy supersaturation and migration were controlled through ; (a) variation in quench rates in binary Al-Ag alloys; and (b) additions of third elements to the binary alloy system. EXPERIMENTAL
DETAILS
Alloy preparation
Alloys of Al-5 at. % Ag were prepared from 99.999 % Al and 99.99 % Ag. The ingots were prepared by air
VOL.
19,
1971
melting and vacuum induction melting, and chill casting the melts into heavy aluminum molds. The cast ingots were then homogenized at 500% for 2 weeks in air. After the homogenizing treatment, the ingots were water-quenched and cold-rolled into 0.006 in. thick foils. No substantial differences in aging response could be detected between the two melting techniques used. Alloys of A14.8 at.% Ag-0.2 at.% Cu and A14.8 at. % Ag-0.15 at. % Cd were prepared from 99.999 y0 Al, 99.99 % Cd, 99.999 % Cu and 99.999 % Ag. Cylindrical ingots were prepared by air melting and then chill casting the melts into heavy aluminum molds. The ingots were homogenized at 550°C for 2 weeks in air. After the homogenizing heat treatment, the ingots were water-quenched and then cold-rolled into 0.006 in. thick foils.
Solution annealing heat treatment and quenching
A portion of the Al-5 at.% Ag foils, wrapped in groups of twenty in aluminum foil to minimize oxidation, were solution annealed in air at 550 & 5°C for 24 hr, and then quenched into ice water at 3°C. The remaining samples, wrapped in groups of ten in aluminum foil, were solution annealed in air at 550 -& 5°C for 24 hr and then air-cooled. The cooling rate for the water-quenched samples was approximately 1500”C/min as compared to approximately 40”C/min for the air-quenched samples. The structures of the air-cooled specimens were studied as-cooled, and were given no subsequent heat treatment. The structures of the water-quenched samples were studied both in the as-quenched condition and after various aging heat treatments.
Aging
All the water-quenched samples were aged in either air or silicone oil at 160 f 5°C for various time intervals up to 1000 hr. The specimens aged in air were wrapped in lots of ten in aluminum foil to minimize oxidation. No difference in the nucleation and growth response of the y’ precipitate was discernible between the two types of aging treatments.
TABLE 1. Alloys and heat treat,ments for samples listed in the text Air quenched at 550°C Sample Sample Sample Sample
No. No. No. No.
1 2 3 4
Yes NO No No
H,O quenched at 550°C v ;:* Yes Yes
Aging time at 160°C (hr) 20: 48 48
At. y0 solute Ag
i.8 5 4.8
Cd
i.15 0 0
CU
:: 0 0.2
PASSOJA
Specimen
AND
ANSELL:
NUCLEATION
OF
y’
IN
Al-Ag
ALLOYS
1255
prepamtion
Foils suitable for transmission electron microscopy were prepared from the aa-quenched and aged samples by first electro-thinning in a modified Lenoir solution at 70-80°C followed by dipping the foils into a hot (7O’C) dilute solution of chromic acid, phosphoric acid and Ha0 in order to remove surface films. A Hitachi HU 125 electron microscope with a 130” tilting stage and ele~troma~etie deflection coils for dark field microscopy was used to observe the specimens.
In addition to those previously reported, several other nucleation mechanisms for the precipitation of y’ were observed in the Al-5 at. % Ag and Al4.8 at. % Ag, 0.2 at.% Cu or 0.15 at.% Cd alloys. Of these, only one type of y’ nucleation was unique in that it was observed to occur in the ternary alloys and not in the binary alloys. The other nucleation mechanisms were found to be present in all the specimens. However, the relative frequencies with which each nucleation mechanism was observed was found to be dependent upon quench rate and alloy content Since most of the nucleation mechanisms were common to all of the specimens examined, the observations of the heterogeneous nucleation of y’ are presented by utilizing the best examples of nucleation mechanisms observed from the various alloys and the corresponding heat treatment. A Thomson tetrahedron is superimposed on some of the mi~rog~phs to facilitate visualization of the geometric relationships. (a) Nucleation on jogged dislocations--stair rod formation. Figures 2 and 3 are examples of y’ nucleation frequently observed in air-quenched samples of Al-5 at. % Ag (sample No. I). Observation of these
PIG. 2. Al-5 et. % Ag air-quenched from 550°C. Climbed extended jogs containing stair rods at A-H. 8
FIG. 3. Al-5 at. % Ag air-quenched from 550°C. Climber extended superjogs. Obtuse and acute at A-2 and A-l, respectively.
st~ctures in the water-quenched and aged alloys was relatively infrequent. This precipitate morphology is similar to that observed by Hren and Thomas(14) in slowly quenched foils studied in &u in the electron microscope. Figure 2 is a ribbon-like structure formed from stacking faults on two sets of (111) planes. It is believed that this structure resulted from the absorption of silver during the queneh, on a jogged edge dislocation, thus causing dissociation of the dislocation. Since the jogged segments are extended, this structure must have stair rod dislocations at the intersections between the jogged sections. The reactions between the Shockley partials which give rise to the stair rods are either #ll] or
+- $[lE]
$[112] -I- ,[llE]
+ +[I101
(I)
-+ i[llO]
(2)
Reaction (1) has the greatest elastic energy decrease if it proceeds to the right, hence it is inferred that an a/G[iiO] stair rod forms at the junctions of the acute angles of intersections A-F. Figure 3 (at A) shows a section of a stacking fault which has climbed out of the plane of stacking faults B and C. The A-l segment has formed acute jogs at the junction of the faults between A and C and is constricted. The other segment A-2 has formed a jog with an obtuse angle and is extended. This dissoeiation reaction could have occurred by means of vacancy condensation and climb of an extended edge dislocation. Barnes(15) has suggested that an extended edge dislocation can climb by the addition of vacancies to the bounding partial dislocations. The dislocations climb by vacancy condensation, and the fault sequence is restored by a shift of a/12[112]; thus the stacking sequence is maintained on alternate (111) planes. It is not likely that this type of climb process
1256
ACTA
METALLURGICA,
can occur extensively in regions where the G. P. zone density is high, however, since the fault would be pinned by the zones in the immediate vicinity of the fault. One must conclude therefore that dislocation jogging occurs more readily at high temperatures, above the coherent solvus temperature, giving rise to a dislocation containing large super jogs. Subsequent precipitation on the dislocation causes a decrease of the stacking fault energy and extension of the partial dislocations. In Fig. 3 the fact that the obtuse bend A-2 is extended whereas the acute bend A-l is not completely extended, is related to the length of the jog as well as the type of stair rods formed at each junction, If a stair rod of Burgers’ vector $lOOl] forms at an obtuse jog, and a +[IlO] stair rod forms at an acute jog, then the energies of the stair rods are in the ratio driving of ~~~~c~~~~~cc~ = 2I1. The therm~~amic force for the f.c.c. -+ h.c.p. transformation is apparently not great enough to overcome the line tension forces of the Shockley partials and hence a constricted junction formed at A-l. Figure 4 (sample No. 2) shows a dark field micrograph of this type y’ nucleation imaged from the (Iii) plane. The displacement fringes from the precipitate are clearly visible on three sets of (111) planes. The fringe between A-B is “invisible” as a result of having overlapping stacking faults, giving rise to @ * .i? = nb (where n = 0, I, 2 . . .).(lQ Between B and C two displacement fringes are visible in the (111) reflection inasmuch as the total displacement from the stacking faults gives Q - R # nb, the b is different for the two faults shown at B-C. These observations suggest that the number of different orientations of y’ which nucleate from a single edge dislocation containing superjogs is de-
FIQ. 4. Aluminum--4.8 at. %Ag-0.16 at.%Cd H,Oquenched from 550°C and aged at 16O’C for 200 hr. Dark field photomicrograph showing displacement fringes and G.P. zones, 9 = @Ii).
VOL.
19, 1971
FIQ. 5. Two
types of y’ nucleation(stair).
pendent upon the nature of the stair rods which form at the junctions of the faults. The stability of stair rod dislocations which form at the junctions of extended dislocations lying on different slip planes is dependent upon the angle between the Burgers’ vector and the line of the initial unjogged dislocation. If the stair rods are unstable relative to a constricted node, only one orientation of y’ will form from one dislocation line (Fig. 5). The observation of jogged-extended dislocations as heterogeneous nucleation sites for y’ in these alloys suggests that the jog-type of nucleation as observed by Nemoto and Koda(l7) is operational in all of these alloys. The number of jogged dislocations observed in water-quenched specimens, were far fewer than those observed in the air-quenched specimens, however. This indicates that jogging occurs more readily during slow cooling from the solution treatment temperature probably due to the more extensive vacancy migration and annihilation occurring above the coherent solvus tern~rature. In rapidly quenched specimens there is extensive dislocation loop formation and helical dislocation formation; i.e. structures resulting from high vacancy supersaturation. The extent of the stair nucleation process is, therefore, strongly dependent upon the quench rate as well as the jog formation processes which occurs at high temperatures above the coherent solvus temperature. (b) ~~c~u~~on on helicaE d~8lo~~~n~. Nucleation of y’ on helical dislocations was frequently observed in the water-quenched and aged binary and ternary alloys. heterogeneous nucleation of y’ on helical dislocations has been observed previously in an investigation by Nicholson and Nutting;(l*) however, some of the details of the decomposiition of the helix and nucleation of y’ on the helix were found to be somewhat different than those reported previously. The decomposition of a helical dislocation proceeds first, by the migration of silver t,o the dislocation. It is di~euIt to ascertain precisely when y’ nucleates on
PASSOJA
XUCLEATION
ANSELL:
AND
OF
y’
IN
,41-Ag
1257
ALLOYS
FIG. 6. Aluminum-5 at. % Ag H,O-quenched from 550°C and aged at 160°C for 48 hr. the helix since the solute
migration
and nucleation
processes appears as one continuous well defined
steps.
Figures
show the decomposition sive segments
6 and 7 (sample No. 3)
of the helix when the succes-
of the dislocation
Decomposition
of
solute build-up
in the vicinity
segments;
the
process without
helix
are not co-planar.
occurs
gradually
of curved
this is usually indicated
by
dislocation
by strong absorp-
tion contrast in the bright field image, resulting from a localized increase in silver concentration At longer aging times the absorption stronger
close to the dislocation
assumes
polygonal
shape
(Fig.
shape which the y’ ultimately thin
hexagonal
shaped
line can be seen surrounding
contrast becomes line, and the line
7).
The
polygonal
assumes is that of a
platelet
bounded by (110) directions
(e.g. Fig. 7).
whose
edges
(the faint hexagonal out-
B in Fig. 7).
When the segments of the helix decompose platelets,
FIG. 550°C
decomposition
are
occurs
by
the
into y’
continuous
7. Aluminum-5 at. % Ag. H,O-quenched and aged at 160°C for 48 hr. Polygonal segments (A and B).
from shaped
FIG. 8. Aluminum-4.8 at. % Ag-0.2 at. % Cu. H,Oquenched and aged for 48 hr at 160°C. Loop decomposition on helix.
localized build-up of solute within the concave portions of dislocation dislocation
Often the helical
segments on the helix. line
decomposes
into
dislocation
which are in the same plane (Fig. 8, sampleNo. when co-planar composition,
dislocation
loops 4); and
loops form from the de-
solute build-up
occurs gradually
the loop, resulting in loop expansion
within
and loop colaes-
cence (Fig. 9, sample No. 4). These types of decomposition be independent frequently
content
for they
could
be observed in both the binary and ternary
water-quenched
FIG. 9.
of impurity
processes appeared to
and aged samples.
Aluminum-4.8 at. %Ag-0.2 at. %Cu. H,Oquenched and aged for 48 hr at 160°C. Loop coalescence on helix.
125x
ACTA
METALLURGICA,
VOL.
19,
1971
FIG. 12. The dislocation node analysis of photomicrograph lo at D. H,OFlo. 10. Aluminum-4.8 at. %Ag-0.15 at. %Cd. quenched from 550°C~aged 200 hr at lBO”C, helix decomposition, 9 = (200)
Figures 10 and 11 are bright fieId and dark fieId photographs (sample No. 2) of y’ which had nucleated on a helix. The nature of the ~splacement vector ofthe precipitate can be obtained from these two pictures by methods outlined by Hashimoto et aZ.(19) and Amelinckx.(20) The phase factor a’ w&s found to be 2~13 sinoe the first displacement fringe of the precipitate at the top of the foil was bright (the top and bottom of the foil are marked T and B, respectively, Fig. 11). Utilizing the method outlined by Amelinckx,(21) it was found that the fault was extrinsie with a shear vector of *(iii). It can be seen that in the region G of Fig. 10 no fringes are observed; furthermore, the dislocation which separates the fringed section from the unfringed section is not visibie, By utilization of the phase factor, cc, dislocation node equations and contrast criteria, it was possible to infer what dislocat,ions are present at the boundary of the precipitate.
Figure 12 shows the dislocation node analysis which was performed on region D of Fig. 10. The nature of the dislocations present at the precipitate boundaries are important with regard to both nucleation and growth mechanisms of the precipitate. I?rom the analysis shown in Fig. 12 it was found that the dislocation reaction at D could be described by the reaction:
giolj 3 gill] + -@I~
(3)
which gives a Shockley partial plus a Frank partial; the stacking fault is extrinsic (Fig. 13) and is similar to a reaction observed by Silcock(2s) in stainless steels. The observation of precipitates with dispia~ement vectors which resemble shear vectors equivalent to that of the extrinsic stacking fault is important since
\
b=-&[iOi]
-------7 trrtt
{ /’
Fro. Il. Dark field pho~micro~&ph of Pig. 10, r* = %/s/3. The top and bottom of the foil are marked 2” and B’, respectively.
Fro. 13. Decomposition of a hefical segment into tEn extrinsic frank partial.
PASSOJA
AND ANSELL:
NUCLEATION
the nucleation mechanism of 7’ depends on the par-
ticular type of stacking fault which is formed. Nucleation of y’ on an intrinsic fault has been suggested(z2) to occur by means of a Suzuki mechanism i.e. the local excess of silver in the vicinity of a dislocation line locally lowers the stacking fault energy, resulting in the extension of the dislocation; the stacking fault then forms the heterogeneous nucleus for nucleation of y’. On the basis of the observations presented in this section an alternate mechanism can be proposed for the nucleation of y’ on helices. During quenching, vacancies are eliminated in the vicinity of the dislocation line due to formation of the helix. Due to the interaction between the vacancies and the silver atoms, the silver concentration increases in the vicinity of the dislocation as a result of the loss of vacancies to form the helix. By increasing the local silver concentration the stacking fault energy around the helix is lowered resulting in the splitting of the dislocation into partials. The dislocation cannot split entirely along its length, however, since it no longer lies in a single plane. Solute thus is attracted to the various extended segments via a Suzuki mechanism. An extrinsic fault may appear if there is no other mechanism of generating a new layer of precipitate. At the aging temperature the composition of the alloys are such that the free energy vs. composition curve has a negative curvature,(za) and thus uphill diffusion can occur. The local extended extrinsic faults can act as continuous sources of vacancies to assist uphill diffusion by operating as climb sources. Thus the precipitate can act as a Bardeen-Heering climb source resulting in the production of a silver-rich extra half plane of atoms bounded by a Frank partials whose shear vector is equivalent to the extrinsic displacement vector of the y’ precipitate. An extrinsic Frank dislocation can grow on (111) planes by dissociation of the helix and expansion by an outward climb of the segments of the Frank partial; the extrinsic Frank partial can then grow by vacancy emission (Fig. 13). This growth mechanism will operate until a new intrinsic fault forms at the precipitate matrix interface by shearing. New precipitate layers can thus be built up by an intrinsic --f extrinsic growth mechanism. Figure 10 at E shows a Frank partial which overlaps segment C on a parallel (111) plane. The phase factor for both segments A and C is 2rr/3, resulting in a phase factor in the overlapping region at E equal to -2rr/3. This results in a reversal of the fringe contrast observed at E. The dislocation bounding the fault A-E has sections which lie parallel to the (110) directions in the (111) plane. The contrast in the over-
OF
y’
IN
Al-Ag
1259
ALLOYS
lapped section at A shows no contrast, since there are three overlapping faults in this region, giving rise to a phase factor equal to zero. It may be concluded, therefore, that multiple extrinsic faults bounded by Frank partials are present on parallel (111) planes. Extrinsic precipitate displacement vectors were not present in every nucleation event which was observed. It is likely that both extrinsic and intrinsic faults are present depending upon the particular conditions which obtain at the time the observation is made. (c) y’ nucleation on dislocation loops. In the waterquenched binary and ternary alloys, nucleation of y’ on prismatic dislocation loops and Frank sessile dislocation loops was by far the most common y’ nucleation mode. The asquenched dislocation loop density of the alloys was in the order of Al-Ag4d > Al-Ag-Cu > Al-AS, whereas the as-quenched loop size was in the reverse order. These trends are probably due to two influences; the ease of dislocation loop nucleation and the vacancy supersaturation possibly as it is related to the vacancy-solute binding energy. If the addition of an impurity atom increases the heterogeneous nucleation of dislocation 10ops(~4) and if, in addition, the solute atoms (Cu and Cd) have a high vacancy-solute binding energy then the dislocation loop density would be high but the loop size would be small. Recent work in ternary Al-Ag-Zn alloys(25) indicates that the dislocation loop density can be strongly influenced by the nature of ternary G. P. zones. Both the heterogeneous nucleation of dislocation loops as well as the vacancy migration in the water-quenched alloys used in this study could be a result of the presence of numbers of ternary G. P. zones having anisotropic strain fields and trapping vacancies.(26) Heterogeneous nucleation of the dislocation loops could also be facilitated by the assistance of the strain field surrounding the G. P, zones. Hence the observed trends of the as-quenched dislocation loop size and density can be explained on a qualitative basis by consideration of the impurities’ independent influences on nucleation and growth of the dislocation loops. (d) Nucleation on heterogeneous climb sources. There was one mode of y’ nucleation which was unique to the ternary Al-AgXu and Al-Ag+d alloys. Figure 14 shows y’ precipitation on a heterogeneous climb source found in the Al-Ag-Cd- (sample No. 2) specimens; Fig. 15 shows a similar precipitate morphology found in Al-AgXu (sample No. 1). In both cases the precipitation associated
had occurred on dislocations with dark impurity
particles
15(b)]. This type of y’ nucleation in the binary Al-Ag alloys.
which were [Figs.
14(a),
was never observed
ACTA
1260
METALLURGICA,
VOL.
19,
1971
In both the binary and ternary alloys heterogeneous nucleation of y f was found to occur on: (a) (b) (c) (d)
FIG. 14. Aluminum-4.8 at. %Ag-OS15 at. %Cd. quenched from 550°C-aged at 160°C for Xuoleation of heterogeneous climb source.
HaO36 hr.
Fra. 16. Atuminum-4.8 at. %Ag-0.2 at. %Cu. quenched from 550”C-aged at 160°C for Nuoleabion on heterogeneous climb source.
H2048 hr.
climb sources have been previously observed in A1-Mgtz7) alloys; decoration of heterogeneous climb sources by precipitates has also been reported for the Al-% system.(2a) In Al-AS alloys, it appears that y’ nucleation will occur on heterogeneous climb sources only if the insoluble impurity particles are present to initiate the heterogeneous climb source. ~e~~~~~~~e~~~
CONCLUSIONS
Decomposition of metastable binary AI-AS and ternary Al-Ag-Cd or Cu alloys occurs by the formation of G. P. zones and nucleation of y’ precipitate phase.
jogged dislocations; helical dislocations; dislocation loops; heterogeneous climb sources.
Only one nucleation mechanism was found to be unique to t.he ternary alloys-nucleation on dislocation climb sources associated with impurity particles. The frequency with which a nucleation event was observed on these dislocations depended upon the relative amounts of dislocation superjog formation, vacancy supersaturation and migration, and was influenced by quench rate and alloy impurity additions. The nucleation controlled by a stair rod process was found to operate by the initial formation of superjogs on edge dislocations by vacancy absorption above the coherent solvus temperature. Nucleation of y’ in one or two orientations on the jogged dislocations depended upon the relative stability with respect to constricted nodes of the extended-jogged segments of the dislocations. A discrete nucleation step for y’ was never observed for nucleation on helices; nucleation of y’ was found to occur by the continuous local solute build-up in the vicinity of the helical dislocations. This resulted in the development of segments of dislocation lines having a polygonal shape. Some of the displacement vectors of these polygonal dislocation segments were found to have extrinsic displacement vectors of &[i%iJ. A dislocation pole mechanism which can act as a continuous source of vacancies and thus generate new layers of y’ on the faces of the y’ platelet, is characterized by the reactions
gior]-+ griTi + gi2il and could control the nucleation process on helices. This disfocation dissociation (a Bardeen-Heering climb source) can act to generate vacancies end new layers of the precipitate by means of climb and an intrinsic --+ extrinsic -+ intrinsic stacking fault growth mechanism. ACKNOWLEDGMENTS
The work reported in this paper was supported by both the National Aeronautics and Space Administration and the office of Naval Research. The studies were conducted in the NASA Interdiscip~a~ Materials Research Center at, Rensselaer Polytechnic Institute, Troy, New York.
PASSOJA
AND
ANSELL:
NUCLEATION
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