Precipitation and the Portevin-Le Chatelier effect in Cu-5.5, 11.6 and 14.2 at.% Ga alloy

Scripta

METALLURGICA

Vol. 18, p p . I 0 4 1 - i 0 4 4 , Printed in the U.S.A.

1984

Pergamon Press Ltd. All rights reserved

P R E C I P I T A T I O N AND THE P O R T E V I N - L E C H A T E L I E R EFFEECT IN Cu-5.5, 11.6 AND 14.2 a t . % Ga A L L O Y

H. ERA, N. O H U R A ~, R. O N O D E R A and M. S H I M I Z U D e p a r t m e n t of Metallurgy, Faculty of Engineering, Kyushu University, Fukuoka, 812, Japan (Received (Revised

April 12, 1984) J u l y 23, 1 9 8 4 )

Introduction The P o r t e v i n - L e C h a t e l i e r effects (serrated flow) in s u b s t i t u t i o n a l f.c.c, alloys has been a t t r i b u t e d to the formation of solute a t m o s p h e r e s around d i s l o c a t i o n s (i). However, we pointed out in p r e v i o u s papers (2-4) that the cause might consist of the interaction b e t w e e n dislocations and clusters of solute atoms such as p r e c i p i t a t e s or short range o r d e r e d regions. A l t h o u g h the cluster theory could explain consistently the P-L effects in many binary systems, some questions which could not be i n t e r p r e t e d by the theory have remained. The most unfavourable e x p e r i m e n t a l result for the theory is the o b s e r v a t i o n of the P-L effect in alloys which were b e l i e v e d to be u n d e r s a t u r a t e d solid solution, such as Cu-Mn (2) and C u - G a (5). In the Cu-Ga system, which is studied in the present work, it has been r e p o r t e d by H u m e - R o t h e r y et al. (6) that the solid solubility of Ga in Cu is 18.6 atomic percent at 473 K. However, Stansbury and Brooks (7) o b s e r v e d a specific heat peak in the temperature range of 423-573 K in C u - 1 7 . 8 a t . % Ga alloy, and Sardar and Gupta (8) g u e s s e d the presence of short range order b e t w e e n 373 and 573 K in C u - 1 7 . 2 a t . % Ga alloy, b a s e d on electrical resistivity measurements. However, they also r e p o r t e d that an electron m i c r o g r a p h of the same alloy a n n e a l e d at 498 or 373 K for ten days r e v e a l e d n e i t h e r any unusual c o n t r a s t effect, like tweed contrast, nor any s u p e r l a t t i c e spots in the s e l e c t e d area d i f f r a c t i o n pattern. In this i n v e s t i g a t i o n the P-L effects in Cu-Ga alloys have been r e e x a m i n e d and evidence of cluster formation in the alloys with lower Ga content has been sought t r a n s m i s s i o n electron m i c r o s c o p y techniques.

further through

E x p e r i m e n t a l Procedure Three kind of m e l t e d and cast specimens with 20 or F u r n a c e - c o o l e d tempered at 673 K

alloys c o n t a i n i n g 5.5, 11.6 and 14.2at.% Ga (analysed composition) were vacuum using 99.9%Cu and 99.9%Ga. After b e i n g rolled and m a c h i n e d into plate-like mm gage length, 5 mm width and 1 mm thickness, these were quenched from 1073 K at a rate of about i00 K/h. For 11.6 and IA.2% Ga alloys, tensile specimens for 20 h were also prepared.

Tensile tests were carried out at temperature b e t w e e n 373 and 773 K. In a tensile test the specimens were pulled at strain rates 3 x i0 - and 5xlO -~ /s a l t e r n a t e l y in order to examine strain rate s e n s i t i v i t y of flow stress. But the strain rate s e n s i t i v i t y will not be discussed further in this paper. For electron m i c r o s c o p e o b s e r v a t i o n the Cu-ll.6 and 14.2at.% Ga thickness of 0.13 mm. The strips were a n n e a l e d for 1 h at 873 K water-quenched. The r e m a i n i n g strips were tempered at 573 or 673 followed by furnace cooling. For e l e c t r o - t h i n n i n g of the strips a and 1 part HNO was used. The e l e c t r o l y t e was kept below 273 K • 3 s p e c l m e n thinning.

alloys were c o l d - r o l l e d to a and then some of these were K for 2 h, 1 day and g days mixture of 2 parts methanol durin£ the ['[na] stages of

Results Fig.l shows the s t r e s s - s t r a i n curves for annealed and q u e n c h e d 5.5, 11.6 and 14.2% Ga alloys tested at 573 K. It should be noted that in the a n n e a l e d state the strength of the ]].6% alloy was higher than that of the i&.2% alloy whereas in the quenched state both alloys showed a

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nearly equal s-s curve. The quenched state revealed higher flow stress than the annealed state in all kinds of alloy; this was especially distinct in the 14.2% alloy. The difference in the flow stress of the annealed and quenched 14.2% alloys is due to both the differences of initial yield point and the work hardening rate. Results of the tensile tests in tempered specimens carried out at 473 K, which were conducted to comfirm that a slow cooling rate causes a lowering in the strength of the 14.2% alloy, are shown in Fig.2. In the 11.6% alloy the shapes of the s-s curves of a quenched and a 673 K tempered specimens were nearly same, though in the 14.2% alloy the flow stress of a 673 K tempered specimen was apparently lower than that of the quenched specimen. The appearance of the P-L effect in the annealed state became remarkable with an increase of the solute content of specimen, as shown in Fig.l, but in the quenched state it was extremely faint in the 5.5% alloy and very weak in the 11.6% alloy. The results mentioned above were nearly same in tests at 373 and 473 K. The comparison of s-s curves tested at 673 K led us to the same summary as above, except that the difference of strength between the annealed 11.6 and 14.2% alloy was less than that in the test at 573 K and the P-L effect in the quenched 5.5% alloy was more distinct than in the test at 573 K. At test temperature equal to 773 K the P-L effects did not occur in all the quenched specimens, and it was very faint in all the annealed specimens. The flow stress of each specimen tested "at 773 K was considerably lower than that at 673 K, which means that the deformation at 773 K depends upon recovery process. The electron micrographs of quenched samples revealed no image of precipitates, no short range order, and no superlattice spots in the diffraction pattern. However, the micrograph of the 14.2% alloy sample tempered at 673 K for 9 days revealed many rod-like images as shown in Fig.3. These images were observed clearly even in the 11.6% alloy tempered at 573 K for 2 h, but the reaction of precipitation in this sample is considered not to have reached equilibrium, for the size of the rod-like images grew after the sample was kept in the electron microscope for 6 min, as shown in Fig.4. The longitudinal direction of the rod-like image was determined to be parallel to <112> from Fig.3 and 4, in which the directions of the incident beam were parallel to <001> and , respectively. In the diffraction patterns of tempered samples, no extra spots were observed, but the matrix spots streaked as shown in Fig.5. Although the streak in the zone (Fig.5(a)) is very faint on the photograph, closer observation on the original film revealed that the streak was along directions and reached about half the distance to the adjacent spot. This observation suggests that the rod-like images observed in the bright field micrographs are due to coherent G-P zones and these have the form of long plates normal to directions and with thickness of 1-2 atoms. The above consideration could be supported by the form of streaks in the <001> and < i i ~ zones shown in Fig.5(b) and (c), respectively. The former revealed short streaks along directions and the wave vector in Fig.5(c) is not parallel to the [ii~ direction but slightly inclined toward the [ii0~ direction (perpendicular to the [i12] direction) so that in the left hand side of reciprocal lattice space three upward spikes, i.e., [ i i ~ , [ll~ and [iIi] spikes intersected the reflection sphere; on the other hand in the right hand side downward spikes, i.e., [i~I~, [ii~ and [ii~ spikes intersect the sphere. Discussion At first, it must be pointed out that the present paper is insufficient as a study relating to precipitation in the Cu-Ga system. Before the experiments it was assumed that even if a precipitation reaction occurred in these alloys, its rate was very slow or indistinct, because it has scarcely been reported in the literature. This is the reason that furnace cooling was preferred to quenching in the samples used for electron microscope observation. However, as it turned out from this work that the reaction rate was neither slow nor indistinct, the completion of a phase diagram for the precipitation of Ga in Cu could, we think, be expected. It is noteworthy that both the strength and work hardening rate of the quenched state were

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h i g h e r than t h o s e of the a n n e a l e d s t a t e Ln the a l l o y s w i t h h i g h Ga c o n t e n t ( e s p e c i a l l y in the 14.2'% a l l o y ) than t h o s e of the 679 K t e m p e r e d state. T h i s f a c t b e c o m e s m o r e i n t e r e s t i n g w h e n we consider that precipitates in the 673 K t e m p e r e d s p e c i m e n r e m a i n e d as s m a l l G-P zones, as s h o w n iri F L g . 3 . To e x p l a i n c o n s i s t e n t l y the a b o v e two e x p e r i m e n t a l r e s u l t s it s h o u l d be a s s u m e d that a G-! zon( in a C u - G a a l l o y can not be an e f f e c t i v e o b s t a c l e to d i s l o c a t i o n m o t i o n , or, al least, ur ]rlerease of s t r e n g t h due to the f o r m a t i o n of G-P z o n e s is less t h a n a d e c r e a s e o f st.pen~ v~Jte ( f" <'-i
I. 2. 3. 4. ~.;. 6. 7. 8.

A. R. R. E. V. M. E. M.

H. C o t t r e ] l , Phil. Mag. 44, 8 2 9 (1953]. Onod(>r'a, T. [ s h i b a s h i , M. K o ~ a a n d M. 8 h i m i z u , A e t a m e t a l l . 31, ()r]od<~na, iI. Era, T. I s h i b a s h i and M. S h i m i z u , A c t a m e t a l l . 31, ()rinu~ra, T. I s h i b a s h i , H. Era and M. S h i m i z u , A c t a m e t a l l . 32, W. [{uuc[( , <). V o h r i n g e r and E. M a c h e r a u c h , Z. M e t a l l k . 64, 296 Hails, r<, ( o n s t i t u t i o n of B i n a r y A l l o y s , p.582. M c G p a w Hill, N e w E. > ~ a n s b u r y and C. R. B r o o k s , A c t s m e t s ] l , l l , 1303 (1963). K. -<%amdar' and K. P. Gupta, S c r i p t s m e t a l [ . 14, B35 ~ 1930). i.

--

535 (1983). 1589 (1983). 8 1 7 (1984). (19"73). Y o r k (1958).

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11.6% Ga alloy

Ga

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x

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FIG.

|, II

3x /S

15

1

N o m i n a l s t r e s s - s t r a i n c u r v e s o f q u e n c h e d and annealed Cu-5.5, 11.6 and 14.2 at.% Ga a l l o y s t e s t e d at 573 K. S m a l l t r i a n g l e s in th< c u r v e s show the s t r a i n s at w h i c h the s t F a i n P a t e was c h a n g e d .

/ / t, II ,I

Test temperature

Strain

strain

FIG.

rate

4

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;

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N o m i n a l s t r e s s - s t r a i n c u r v e s of q u e n c h e d and 673 K tempered Cu-ll.6 and [4.2 a t . % Ga a l l o y s t e s t e d at 4 7 3 K. The m e a n i n g of" s m a l l t r i a n g l e s is s a m e as in Fig. i.

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Bright field micrograph of the rod-like G-P zones in Cu-14.2 at.% Ga alloy tempered at 673 K for 9 days, showing that the trace of the longitudinal direction of rod is parallel to <210>. The incident beam is parallel to a [001] direction.

FIG.

4

Bright field micrographs of a Cu-ll.6 at.% Ga alloy tempered at 573 K for 2 hours (a) and the same region as photo.(a) taken after keeping in microscope for 5 minutes (b). The incident beam is inclined to [110] from [110] directin by a few degrees.

FIG. 4(a)

FIG.

FIG. 4(b)

5(a)

FIG.

5(b)

FIG.

5

FIG.

Diffraction patterns of (a), <001> (b) and zone (c) in Cu-14.2 at.% Ga alloy tempered at 573 K for 9 days. Arrows in fig.(a) indicate long streaks.

5(e)