Metastable phase formation and decomposition in a rapidly solidified aluminium-platinum alloy

Materials Science and Engineering, 38 (1979) 7 - 17 © Elsevier Sequoia S.A., Lausanne -- Printed in the Netherlands

7

Metastable Phase Formation and Decomposition in a Rapidly Solidified Aluminium-Platinum Alloy K. CHATTOPADHYAY and P. RAMACHANDRARAO Department o f Metallurgical Engineering, Banaras Hindu University, Varanasi - 221005 (India) (Received July 19, 1978)

SUMMARY

a metastable phase on annealing the quenched foils. We have reinvestigated the effect of rapid solidification on A1-Pt alloys as part of a larger programme on aluminium-based alloys and the present paper deals with the results obtained on an AI-2 at.% Pt alloy.

The metastable phase formation and decomposition behaviour of a rapidly solidified A12 at.% Pt alloy were studied by electron microscopy and X-ray diffraction techniques. Two metastable phases were detected. One of them has a cubic structure with a = 5.67 £ and can be considered as a metastable extension of the e q u i l i b r i u m A12 Pt phase. The other phase, of composition A16Pt, has an orthorhombic structure with lattice parameters a = 15.762 A, b = 12.103 £ and c = 8.318 A and is isostructural with the equilibrium phase Gas Pt. In the as-splat foils, the A12 Pt phase forms at the interdendritic regions whereas A16 Pt forms at the predendritically solidified regions. After prolonged aging the A16 Pt phase grows with the dissolution of the Al 2 Pt phase. The mechanisms of formation of the A1e Pt phase during artificial aging and natural aging are different. While normal nucleation processes control the decomposition during artificial aging, sympathetic nucleation of A16 Pt is the dominant mode of decomposition in the naturally aged foils and leads to a star-like morphology. The reasons for the formation of the above metastable phases are also discussed.

Small (about 150 rag) samples of the AI-Pt alloy prepared by induction melting high purity {five 9's) constituent elements in an inert atmosphere were quenched from the liquid state using the " g u n " technique [4]. The resultant foils were thin enough in many areas to be examined directly in a Philips EM 300 electron microscope. X-ray examination was conducted using both D e b y e - S c h e r r e r and Guinier cameras and Cu K~ radiation. Annealing experiments were carried out both in the h o t stage of the microscope and in evacuated and sealed quartz capsules for X-ray examination purposes. Some samples were aged at room temperature for as long as two years and they were subjected to X-ray and electron microscope examination from time to time.

1. INTRODUCTION

3. RESULTS

Under equilibrium conditions platinum has negligible solid solubility in aluminium and several intermediate phases are formed on alloying the two elements [ 1 ]. However, rapid solidification of aluminium-transition metal alloys is known to result in an increase in the solid solubility of the transition metal in aluminium [ 2 ] . In an earlier investigation, Tonejc et al. [3] studied A1-Pt alloys containing up to 3 at.% platinum and reported the extension of solid solubility as well as the formation of

3.1. As-quenched foils X-ray diffraction patterns from the asquenched foils showed the presence of only a few faint extra lines in addition to those from the aluminium solidsolution. The solid solution lines were broad and not amenable to accurate lattice parameter determination. Significant information regarding the microstructure and constitution of the as-quenched foils could thus be obtained only with the aid of transmission electron microscopy.

2. EXPERIMENTAL

All the areas which were highly transparent to the electron beam exhibited a typically dendritic microstructure with considerable interdendritic segregation (Fig. 1). Very often the secondary dendrite arms were found to be closely parallel to each other with very little development of ternary arms {Fig. 2). In rela-

Fig. 3. Microstructure of a thick area where the cooling rate is comparatively low.

Fig. 1. Microstructure of as-splat foils.

Fig. 4. Predendritic region with a star-shaped precipitate in the predendritic disc.

Fig. 2. Parallel dendrites g!ving rise to a modulated appearance. tively thicker sections the dendrite morphology was coarser and the existence of a precipitate phase in the interdendritic region was noticeable (Fig. 3). The rate of cooling estimated from the spacing of the secondary dendrite arms varied from about 10 l° ks -1 to about l 0 s ks -1 depending upon the thickness of the foil. In several other thin areas we could observe predendritic solidification of the type reported earlier [5]. Invariably the central portion of the predendritic structure was characterized by the existence of a star-like precipitate (Fig. 4).

Electron diffraction patterns obtained from regions characterized by parallel secondary dendrite arms with interdendritic segregation consisted of a spot pattern due to the aluminium solid solution and a superimposed ring pattern {Fig. 5(a)) which could n o t be indexed in terms of stable intermediate phase A14 Pt which should form under equilibrium conditions. Dark field observation (Fig. 5(b)) showed t h a t the phase responsible for the rings is in the interdendritic regions. The interplanar spacings corresponding to the ring pattern are shown in Table 1. These could be indexed on the basis of a cubic phase with a lattice spacing of 5.67 + 0.03 A. It is significant that the X-ray diffraction patterns do not show any reflections corresponding to this phase.

was observed that aging the foils for about 15 days led to the formation of star-shaped precipitates similar to those observed in the predendritic regions. These precipitates were found to grow at the expense of the interdendritic phase discussed earlier (Fig. 6). Prolonged aging resulted only in an increase in the density of the star-shaped precipitates.

(a)

Fig. 6. G r o w t h o f A16Pt stars w i t h t h e s i m u l t a n e o u s d i s s o l u t i o n of t h e i n t e r d e n d r i t i c A12Pt phase. N o t e t h e d e n u d e d zone. (b) Fig. 5. (a) D i f f r a c t i o n p a t t e r n f r o m t h e area s h o w n in Fig. 1. ( b ) D a r k field m i c r o g r a p h o b t a i n e d b y placing t h e a p e r t u r e o n t h e i n t e n s e ring. TABLE 1 O b s e r v e d a n d c a l c u l a t e d i n t e r p l a n a r spacing o f t h e Al2Pt phase obtained from electron diffraction Ring

dob s

dca 1

no.

(A)

(A)

1 2 3 4 5

3.25 1.99 1.70 1.41 1.23

6

1.01

3.27 2.00 1.71 1.42 1.27 1.16 1.00

hkl

Intensity

111 220 311 400 420 422 440

s, b vs vw, b m w, v b w -m

a = 5.67 ± 0.03 A. vs, very s t r o n g ; s, s t r o n g ; m , m e d i u m ; w, w e a k ; b, b r o a d ; vb, very b r o a d .

3.2. R o o m temperature aged foils R o o m temperature aging of as-quenched foils was carried out for up to two years. It

X-ray diffraction patterns from samples aged for up to six months showed a large number of reflections which could n o t be explained on the basis of equilibrium phases at this composition. The patterns obtained could be indexed in terms of an orthorhombic unit cell w i t h a = 15.762 £ , b = 12.103 £ a n d c = 8.318 A. The results are shown in Table 2 together with the metastable phase lines reported by Tonejc et al. [3]. These investigators assigned a tetragonal unit cell with a = 13.58 £ and c = 16.64 £ for the metastable phase. Except for the first two reflections obtained by Tonejc et al., the agreement between the calculated interplanar spacings on the basis of our cell and those observed by us and by Tonejc et al. is excellent. It should be noted t h a t the first two reflections reported by Tonejc et al. were at extremely small angles and have n o t been detected by us. At these small angles errors in the determination of the interplanar spacings can be very large. The fact t h a t our patterns did n o t show these lines could also mean t h a t their intensity is very low. Thus the lack of agreement between calculated and observed

10 TABLE 2

Tonejc

Present investigation dc81

et al. dobs

dobs (A)

13.522 10.500 8.27 6.841 6.074 5.533 5.259 4.907

4.816

w -m

4.160

4.177

vw

3.935

3.961

w

(A)

2.024 1.979

vsa w

1.964

1.966

vw

1.939

1.939

vw

1.92~ 1.909

sa w

1.885 1.789 -

w vw -

1.756

w

-

The interplanar spacings of the Al6Pt phase observed by Tonejc et al. [3 ] and in the present investigation together with indexing on the basis of an orthorhombic cell having a = 15.762 A, b = 12.103 A and c = 8.318 k

Iobs

-

hkl

(A)

-

-

-

3.819 3.768 3.642 3.510 3.433 3.397 3.274

3.835 3.778 3.652 3.511 3.437 3.276

3.232 3.209 3.145

3.209 3.138

vw sa

3.040

3.036

w

2.965 -

m

2.901 2.855 2.775 2.720 2.683 2.621

-

2.452

2.452

vw, b vw, b

2.262

2.409 2.336 2.289 -

2.147

2.175 2.149 2.082

w vw vw, b

2.048

2.050

w

-

-

-

vw m m vw m w -m

v s

a

vw

12.103 9.600 8.318 6.855 6.052 5.650 5.254 4.894 4.819 4.800 4.160 4.170 3.933 3.941 3.967 3.813 3.747 3.630 3.519 3.428 3.416 3.261 3.297 Al2Pt (?) 3.200 3.143 3.149 3.026 3.050 2.970 2.948 2.870 2.860 2.773 2.718 2.675 2.622 2.627 2.452 2.447 2.403 2.337 2.285 2.261 2.268 2.174 2.139 2.079 2.082 2.049 2.052

010 110 001 011 020 120 3O0 021 310 220 0 0 2 ,2 2 1 311 012 40O 320 112 410 031 212 022 411 302 231 330 222 312 O4O 510 140 501 322 402 OO3 232 241 340 600 303 042 313 242 033 133 403 701 711 0O4 5O3 014 513

1.791 1.779

-

2.017 1.978 1.980 1.967 1.966 1.945 1.943 1.933 1.932 1.920 1.908 1.909 1.883 1.790 1.777 1.780 1.778 1.755

060 243 702 024 730 161,810 352 304 542 550 224 314 360,613 642 253 650 732 461

The intensities are taken from the present investigation. vs, very strong; s, strong; m, medium; w, weak; vw, very weak ; b, broad. aLines superimposed on aluminium or silicon lines. Silicon was used as a standard. interplanar seriously.

spacings need

not

be taken

3.3. Elevated temperature aging o f quenched foils Hot stage electron metallography of the foils r e v e a l e d t h a t the i n t e r d e n d r i t i c segregat i o n o b s e r v e d in t h e a s - q u e n c h e d f o i l s w a s stable for a period of 10 min up to a temperat u r e o f a b o u t 2 5 0 °C ( Fi g . 7(a)). A f u r t h e r i n c r e a s e in t e m p e r a t u r e t o 3 0 0 °C n u c l e a t e d a phase at the grain b o u n d a r i e s (Fig. 7(b)). M a i n t a i n i n g t h e t e m p e r a t u r e a t 3 7 5 °C w a s s u f f i c i e n t t o r e p l a c e t h e i n t e r d e n d r i t i c segretion by a uniform distribution of fine precipitates (Fig. 7(c)). L o n g e r t i m e s at l o w e r t e m p e r a t u r e s a l s o r e s u l t e d in t h e f o r m a t i o n o f f i n e p r e c i p i t a t e s b o t h in t h e g r a i n a n d a t g r a i n b o u n d a r i e s . E l e c t r o n d i f f r a c t i o n p a t t e r n s obt a i n e d f r o m t h e fine p r e c i p i t a t e s s h o w e d rings c o r r e s p o n d i n g t o t h e c u b i c p h a s e n o t i c e d in the a s - q u e n c h e d foils. A g i n g a t t e m p e r a t u r e s c l o s e t o 4 0 0 °C o r for prolonged times at lower temperatures r e s u l t e d in a b i m o d a l d i s t r i b u t i o n o f p r e c i p i t a t e s . T h e p r e c i p i t a t e s d i f f e r e d b o t h in size a n d m o r p h o l o g y ( Fi g . 8). O n e o f t h e p r e c i p i tates often assumed a rectangular cross section and formed to about 1000 - 2000 A within the grain and up to 4 0 0 0 A at the grain boundaries. E l e c t r o n d i f f r a c t i o n patterns obt a i n e d f r o m t h e s e p r e c i p i t a t e s c o u l d b e in-

11

(a) Fig. 8. Bimodal distribution of coarse and fine precipitates obtained by aging above 400 °C.

sections could be interpreted in terms of the cubic phase discussed earlier (Fig. 10).

4. DISCUSSION

(b)

(c) Fig. 7. The same area as Fig. 2 heated for 10 rain at (a) 250 °C, (b) 300 °C and (c) 375 °C.

dexed in terms of the orthorhombic phase which was noticed in foils aged for long times at room temperature (Fig. 9). Single-crystal electron diffraction patterns from the smaller precipitates with square or rectangular cross

The earlier results obtained by Tonejc et al. [3] using the X-ray diffraction technique only suggest that alloys containing up to 2 at.% platinum should consist of a homogeneous solid solution. However, our electron metaUographic results showed that in no case was a homogeneous solid solution obtained. Platinum appeared to segregate rapidly to the interdendritic areas of the solidifying film. Alternatively, the microstructure of the transparent areas could have resulted from a spinodai decomposition of a single phase solid solution. Several attempts were made to obtain diffraction evidence for a possible spinodal decomposition. Even at the highest camera constants or at the greatest enlargement of the diffraction spots no trace of any satellite reflections near the matrix spots could be detected. We are therefore constrained to conclude that the microstructure of the quenched foils represents interdendritic segregation only. According to the model due to Sare and Wood [6] highly parallel dendritic arms could result from a characteristic heat transfer in the plane of the foil. These investigators proposed that the thinnest areas need n o t necessarily represent the points of contact with the substrate. They argued that areas which are intermediate to points of contact of melt with substrate (lift-

12

(a)

(b)

i 2°°AI

4"04•

"IAla

3,~o0

-

0 1 4 42,4 • • •

4,0Q + •

[3

iliA[

711

013

,]T •

•

•

o7°

ar

•

•

702

~

O

--

• 70-2

0 (,33)AI

- o OI3

_e 71~

[]

•

--Ill

}OOAI

ZON E ZONE

•

r2 ~--i 7] Ale• p t

n [o~ 0 A= • Et3,']A[6P t

(e)

(d)

Fig. 9. (a, b) Two prominent zones of the AI6Pt phase. (c, d) Schematic illustrations of patterns (a) and (b) respectively.

off areas) cool by transfer of heat to the point of contact. Under these conditions heat transfer is mainly in the plane o f the foil and could encourage the formation of columnar grain structure in lift-off areas. In agreement with their proposals we obtained columnar grains in the smaller lift-off areas. The dendrite arms were parallel to the columnar grains. In contrast to the lift-off areas, the predendritic areas (Fig. 4) represent areas which have cooled by direct contact with the substrate. As pointed o u t by Biloni and Chalmers [5] and Aptekar and Kamanetskaya [ 7 ] , the central part of the predendritic region (termed the disc) undergoes a diffusionless transformation. The boundary of the disc represents points at which heat of recalescence could n o t be conducted away sufficiently fast. At these points diffusion becomes favourable and typical dendritic morphology arises. The disc itself represents an area of supersaturated solid

solution with a concentration equivalent to the nominal composition of the alloy. Our observation that star-shaped precipitates are present in these disc regions supports the hypothesis that the disc was originally supersaturated and hence acted as a preferential nucleation site for metastable phase precipitates. The mechanism of nucleation of these precipitates is discussed later. The present investigation clearly shows that two metastable precipitates can be obtained in these dilute alloys. The structure and equilibrium isomorphs of these phases will n o w be discussed in detail.

4.1. The cubic phase A cubic phase with a lattice constant of 5.67 + 0.03 A forms in the interdendritic areas of the quenched foils in the form of an extremely fine precipitate. In the as-quenched condition both its size and quantity are such

13

as to preclude the possibility of securing any X-ray evidence for its existence. Hence it is n o t surprising that Tonejc et al. were unable to detect its presence. Even in the present study it could only be established on the basis of electron diffraction evidence. It is interesting to note that a metastable cubic phase with a comparable lattice parameter was obtained b y Jacobs et al. [8] in rapidly solidified A1-Fe alloys. In these alloys the phase was also indexed on the basis of electron diffraction patterns. Owing to the absence of (200) and (222) reflections in the patterns, Jacobs et al. concluded that the phase has a diamond cubic structure. In our patterns also the (200) and {222) reflections could n o t be detected. However, the assumption of a diamond cubic structure leads to an anomalously large volume per atom in the structure. The resultant atomic volume of 22.79 h 3 is far greater than the atomic volu m e of pure aluminium (16.6 A s) and pure platinum (15.10 As). A similar anomaly exists in the interpretation of Jacobs e t al. concerning the metastable A1-Fe phase [ 9 ] . Alternatively, the observed lattice parameter of 5.67 + 0.03 A can be considered to be close to the lattice parameter of 5.927 A of the A12Pt equilibrium phase. A12Pt has a CaF2 t y p e cubic structure. The atomic volume calculated on the basis of the CaF2 structure is 15.19 k 3 which is far closer to the atomic volumes of the constituent elements than that obtained on the basis of a diamond cubic structure. On the basis of the CaF2 structure the (200) and (222) reflections are permitted. However, structure factor calculations based on the assumption that the atomic positions in the metastable phase are akin to those in the equilibrium A12Pt phase show that the (200) and {222) reflections are inherently weak (Table 3). Furthermore, examination of TABLE 3 Calculated structure factor for the first f o u r reflections of the cubic A12Pt phase on the basis of electron diffraction hkl

IFI 2

111 200 220 222

1425 145 1765 126

the structure factor equations shows that any deficiency of platinum in the A12Pt structure affects reflections of the type h + k + l = 2n where n is odd to a greater extent than other reflections for which all the indices are even or h, k and I are all odd. On the basis of these observations it may be reasonable to conclude that the cubic phase is A12Pt. As has already been mentioned, aging for short times at elevated temperatures results in a more uniform distribution of the A12Pt phase in the foils. The precipitates appear to grow at the expense of the interdendritic segregation. At this stage single-crystal electron diffraction evidence could be obtained. Figure 10 shows two such zones and their schematic illustrations. The orientation of the A12Pt phase with respect to the aluminium matrix could be determined from patterns such as that shown in Fig. 11, the details of which are given elsewhere [8, 9 ] . The relation can be expressed as (Ill)precipitate

]l ( 0 0 1 ) m a t r i x

[ - 1 1 0 ] precipitate I] [ 0 1 0 ] m a t r i x

The orientation relation is again similar to that obtained by Jacobs et al. for the metastable A12Fe phase. If the existence of a phase of composition Al2Pt is accepted, its occurrence in the interdendritic spaces can be understood. Apparently the concentration of platinum in the interdendritic space is high enough to nucleate A12Pt. However, such a phase should be unstable at the nominal composition of the alloy. It should therefore disappear and give rise to the equilibrium A14Pt phase. 4.2. T h e o r t h o r h o m b i c phase

As the A12Pt phase becomes unstable on annealing at temperatures close to 400 °C, a second metastable phase with an orthorhombic structure appears. We have been able to assign this unit cell to the phase on the basis of X-ray diffraction evidence. It is interesting to note that the lattice parameters of this phase are very similar to those of the phase Ga6Pt [10]. Both gallium and aluminium belong to the same group in the periodic table and the A1-Pt and G a - P t equilibrium diagrams exhibit many similarities which have recently been discussed in detail [ 1 1 ] . For example, in both systems a C a F 2 phase occurs at 33.33 at.% platinum, Ni2Als type phases appear at 40 at.% platinum

14

(a)

(b)

222 Ill •

•

aTo +

Ts~T ~,~.o °

•

~§, •

•

•

o_

402 3hi 220 131

--

242 0

•

• e l IT

4- 0 0 0 • oTo

ZONE

ZONE

• A t 2 P t 1"112]

• At2Pt [123]

(c)

(d)

Fig. 10. (a, b) Two prominent zones of AI2Pt. (c, d) Schematic illustrations of patterns (a) and (b) respectively.

Fig. 11. Selected area diffraction pattern used for obtaining the orientation relation.

and FeSi type phases appear at 50 at.% platinum. In the aluminium-rich side, the phase GasPt has previously been regarded as analogous to A14Pt. In view of this close similarity between the two systems and the close agreement between the lattice parameters of the metastable phase

under discussion and GasPt, it is reasonable to assign a formula A16Pt to the metastable phase. In order to obtain a reasonable atomic volume (16.18 £3) fourteen formula units have to be included in each unit cell of the phase. The calculated atomic volume of 16.18 £3 is less than the value obtained on the basis of Zen's law for stoichiometric AlsPt. Edshammer [12] has considered the atomic volume of aluminium-rich phases in the AI-Pt system and pointed out a tendency for negative deviation from Zen's law. The atomic volume obtained by us agrees well with the scheme suggested by Edshammar. Single-crystal electron diffraction patterns obtained during the later stages of annealing at elevated temperatures or during long term aging at room temperature clearly demonstrate the correctness of the proposed structure. Figure 9 shows the two most prominent zones of AlePt and the schematic illustrations of the indices assigned to the diffraction spots. A tentative orientation relation can be obtained

15 from Fig. 9(b) which can be expressed as follows: [011]A1 ~ II [141]Al6Pt (200)A 1 ~

II

(014)A16Pt

AS has already been pointed out, A16Pt nucleates preferentially (even after a short period at r o o m temperature) in the central regions of the predendritic areas of the quenched foil. Invariably the morphology of AlePt precipitates is star-like in foils aged at room temperature while the precipitates assume a massive rectangular section when formed by aging at elevated temperatures. The star-shaped AlePt particles are usually surrounded by a denuded zone and then an Al2Pt network. Thus it appears that the growth of A16Pt particles is associated with the dissolution of the A12 Pt precipitates. In order to understand the star-like morphology careful electron microscopy was carried o u t to identify the early stages of precipitate growth. Figure 12 shows a micrograph obtained from a sample aged for three days at room temperature. A row of platelets can be seen in strain contrast which confirms that originally only single platelets nucleated. Further nucleation occurs on top of the platelets which are already nucleated resulting in a projected star-like shape. Mehl and Marzke [13] and Massalski [14] have observed a similar precipitation m o d e and morphology of precipitates in copper-based alloys. Aaronson

--

_~

.~-

|._

|~

' O.2p

I

Fig. 12. Foils aged at room temperature for three days (A16Pt plates in strain contrast). and Wells [15] have described the origin of such precipitates in terms of "sympathetic nucleation". The AIePt phase appears to be stable over a large temperature range. According to Tonejc et al. the metastable phase detected by them (which we consider to be AlsPt) decomposes only at temperatures close to the melting point. Our results confirm this observation.

4.2. Mechanism o f metastable phase formation Under equilibrium conditions the hypereutectic alloy investigated should contain primary A14Pt and the a aluminium solid solution, The results obtained in the present investigation are summarized in Fig. 13. It is im-

DfF U ~ 4 L E SS 5,:JLIOIF~..AT~ON

•

l

MEI~ AJ - ~PT

ALLOYL

E-C,~NS

J

1

•

__

R.T AC.~I~G ICTiC (At 4"AI4PT )i

/~TFK::IAI. AC~IING

T~TUI~

PLATEI.ETS

I

I AND (~SSOLU- I

Fig. 13. Scheme of the phase transition in the A1-2 at.% Pt alloy due to rapid solidification and subsequent aging.

16

portant to note that the formation of the equilibrium A14Pt is totally suppressed by quenching from the liquid state. Even artificial aging of the foil does not result in the formation of A14Pt except at temperatures close to the melting point. The as-quenched foils consist of "equilibrium"A12 Pt and metastable phase A16Pt. Significantly, the formation of the other equilibrium intermediate phase AI3Pt is also suppressed on both quenching and aging. Under equilibrium conditions all three equilibrium phases A14Pt, A13Pt and A12Pt form peritectically. The first two phases are line compounds while A12Pt is stable over a small composition range. Another point of difference arises with reference to the composition of the phases vis d vis the composition of the liquid with which they are in equilibrium. While the stoichiometric composition of A14Pt and Al3Pt are far removed from the composition of the liquids with which they are in equilibrium at the peritectic temperatures, the composition of A12Pt is close to that of the corresponding liquid. Earlier investigators [2] have shown that quenching from the liquid state can easily suppress the formation of intermediate phases resulting from peritectic reactions under equilibrium conditions. Suppression appears to become easier when the composition of the peritectically forming phase and that of the liquid in equilibrium with it are very different. Such a hypothesis would explain the suppression of the nucleation of A14Pt and A13Pt in preference to A12Pt. It is therefore tentatively suggested that a metastable equilibrium exists between aluminium and A12Pt. Possibly, a eutectic results between aluminium solid solution and A12Pt at some undercooled temperature with a concentration greater than that of the alloy studied. As a consequence, aluminium nucleates as a primary phase while A12Pt forms in the interdendritic phase by a eutectic reaction. While the present investigation shows the existence of a very finely dispersed A12Pt phase in the interdendritic spaces of aluminium, the evidence is as yet insufficient to prove t h e a b o v e hypothesis conclusively. 5. CONCLUSIONS T h e present study has established several interesting aspects of the formation and sta-

bility of metastable phases in a rapidly solidified AI-2 at.% Pt alloy. (1) Extension of solid solubility is limited to highly localized regions in the foil. Such regions are at the centre of the predendritic regions. Elsewhere, diffusion of platinum is rapid and platinum is segregated at the interdendritic regions. (2) The high concentration of platinum in the interdendritic areas is capable of nucleating the "equilibrium" Al2Pt phase. This phase is, however, metastable at the composition of the alloy investigated and redissolves to form another metastable phase Al6Pt. (3) The metastable phase A16Pt is isostrucrural with Ga6Pt and has an orthorhombic unit cellwith the parameters a = 15.762 A, b = 12.103 A and c = 8.318 A. This phase appears to be highly stable and can be retained up to temperatures close to the melting point of the alloy where it transforms to the equilibrium A14Pt phase. (4) During natural aging, sympathetic nucleation of the AlePt phase results in a characteristic star-like morphology. ACKNOWLEDGMENTS T h e authors are grateful to Dr. S. Lele for many stimulating discussions throughout this investigation and for reading the manuscript. They also thank Professor S. Ranganathan for his help in preparing the alloy and Professor H. I. Aaronson for helpful comments.

REFERENCES 1 M. Hansen and K. Anderko, Constitution of Binary Alloys, McGraw-Hill, New York, 1958. 2 T.R. Anantharaman, P. Ramachandrarao, C. Suryanarayana, S. Lele and K. Chattopadhyay, Trans. Indian Inst. Met., 30 (1977) 423. 3 A. M. Tonejc, A. Tonejc and A. Bonefacic, J. Mater. Sci., 9 (1974) 523. 4 P. Duwez and R. H. Willens, Trans. Metall. Soc. AIME, 227 (1963) 362. 5 H. Biloni and B. Chalmers, Trans. Metall. Soc. AIME, 233 (1965) 373. 6 J. V. Wood and I. R. Sare, Metall. Trans., 6A (1975) 2153. 7 I. L. Aptekar and D. S. Kamenetskaya, Phys. Met. Metallogr. (USSR), 14 (1962) 33. 8 M. H. Jacobs, A. G. Doggett and M. J. Stowell, J. Mater. Sci., 9 (1974) 1631. 9 K. Chattopadhyay, S. Lele and P. Ramachandrarao, J. Mater. Sci., 13 (1978) 2730.

17 10 S. Bhan and K. Schubert, Z. Metallkd., 51 (1960) 327. 11 P. Guex and P. Feschotte, J. Less-Common Met., 46 {1976) 101. 12 L. E. Edshammar, Acta Chem. Scand., 20 (1966) 2683.

13 R. F. Mehl and O. T. Marzke, Trans. Am. Inst. Min. Metall. Pet. Eng., 93 (1931) 123. 14 T. B. Massalski, Acta Metall., 6 (1958) 243. 15 H. I. Aaronson and C. Wells, Trans. Metall. Soc. AIME, 206 (1956) 1216.