Dynamic softening of spinodal alloy Cu9Ni6Sn during hot deformation

Materials Science and Engineering, A 137 ( 1991 ) 185-188

185

Dynamic softening of spinodal alloy Cu9Ni6Sn during hot deformation A. Y. Abdellatief and P. Kratochvil Department of Metal Physics, Charles University, 12116, Prague 2, Ke Karlovu 5 (C~echoslovakia)

Abstract The response of polycrystalline Cu9Ni6Sn (i.e. 9 at.% Ni, 3.1 at.% Sn) to hot deformation was investigated in the temperature range 5 7 0 - 8 5 0 °C at initial strain rates g~ between 1.7 × 10 -2 and 3.3 × 10 -5 s t. The activation energy of deformation and the stress exponent were determined. Flow behaviour reveals softening by dynamic recovery in the temperature range 5 7 0 - 6 5 0 °C and by dynamic recrystallization between 700 and 850°C. The first range is characterized by higher values of the activation energy of the deformation process and of the stress exponent than in the second range. The fracture ductility was found to increase with increasing temperature of deformation. The effect of tin content on the dynamic softening in Cu9NixSn was studied in the alloys containing x = 0, 0.5, 1.2 and 3.1 at.% Sn.

1. Introduction Although the strengthening of the Cu9Ni6Sn alloy which is due to the so-called spinodal decomposition has been studied extensively (see for example refs. 1-6) very little is known about the dynamic softening mechanism during hot deformation [7]. It is the purpose of the present paper (i) to study the dynamic softening mechanisms in the Cu9Ni6Sn alloy in a wide range of deformation conditions, and (ii) to measure and understand the effect of the tin content x on the dynamic softening process of the Cu9NixSn alloys. The activation energy Q of deformation, the stress exponent n, the yield stress 00.2 and the fracture ductility em~~ were the quantities used to describe the investigated process. 2. Experimental results and discussion

2.1. Experimentalprocedure The alloys were prepared in the Institute of Metal Research, Panensk~ B~e~any. They are Cu9NixSn alloys with x = 0, 0.5, 1.2 and 3.1 at.%. As-cast alloys were cold rolled in several steps followed in each step by an annealing treatment. The last area reduction by cold rolling was 50% and the thickness of the strips was 0.35 +0.03 0921-5093/91/$3.50

mm. All specimens were sealed in silica tubes, solution treated at 800 °C for 30 min and water quenched. The specimens were deformed in tension in an argon atmosphere using Instron 1195 equipment at temperatures between 550 and 850°C and at initial strain rates go between 1.7 × 10 -2 and 3.3 × 10 -5 s-1. Specimens rapidly quenched after deformation were examined by optical microscopy and transmission electron microscopy. If it is assumed that the flow behaviour of the hot deformation obeys the relation [8]

Z= go exp(Q/RTd)=A sinh(aamax)"

(1)

the best fit of the experimental data ( To, amax and go) makes it possible to find the values of Q and n. Td is the temperature of deformation, and a and A are constants.

2.2. Dynamic softening of the CugNi6Sn Examples of the flow behaviour of the Cu9Ni6Sn alloy are shown in Figs. 1 and 2. It is obvious that the flow stress o increases with increasing go and decreasing Td. The plastic strain to the peak stress Umax decreases with increasing Td. It is also clear from the true stress-true strain o - e curves that when dynamic softening takes place by dynamic recrystallization (DRX) only, a single peak was observed. It is © Elsevier Sequoia/printed in The Netherlands

186

200

I

~

I

" --"""~o.

1.?x10_2S.1

I

1 0 -2

/ ~ ~ " ' , 6.7,~1o-'

t~

~o10~

100 ~ ~

~

.

.

"6.7~i70~10"3

Td-6500C

\ 6,7x10"5 0~

010

I 0.20

6

Fig. l. True stress-true strain curves of the Cu9Ni6Sn alloy at Ta = 650 °C for different strain rates go.

200

~

i

b 100c ~ ~ " ' - .

~0°c

~

I

I

600

'700

Td (°C)

800

Fig. 3. "Deformation map" for Cu9Ni6Sn alloy.

i #17"1°-3s -'

g

\650%

~ - - - ~ 0 0 ° C ~_

_a c --------_

~--

-"850%

O'

I

0.10

1

0.20

or~o C

I ] ,

Fig. 2. True stress-true strain curves of the Cu9Ni6Sn alloy measured with an initial strain rate ~0 = 1.7 x 10 -3 s 1 at different temperatures.

obvious that steady state flow stress is not attained during softening by DRX at TD < 700 °C. As the softening by dynamic recovery (DRV) can be distinguished from that by DRX with the help of the shape of the a - e curves, two regions divided by a boundary can be obtained (Fig. 3). In the temperature range 570-650°C DRV is dominant, while in the temperature range 700850°C DRX was observed. It is clear that the boundary indicating the onset of the DRX can be roughly established as Z c = 5 x 107 s- 1_ During DRX the particles of (CuxNil_x)3Sn phase having D03 structure are interacting with the grain boundaries (GBS) (Fig. 4(a)). The difference in the dislocation densities /9 on the two sides of the GB is obvious. Inside the gains the subgrain structure with stable dislocation configurations is typical (Fig. 4(b)). The dislocation substructure in the samples with the DRV-type a - e curve is formed by very dense unstable dislocation clusters (Fig. 5). The

Fig. 4. (a) T h e GB during the DRX. T h e deformation conditions were Td = 800 °C and e0 = 1.7 x 10 -3 s - 1. T h e strain is 18%. (b) Subgrain structure and stable dislocation configuration within the grain in the same sample as in (a). (Magnifications: (a) 9000 x ; ( b ) 2 0 0 0 0 x .)

bent dislocation segments are due to pinning oI dislocations on the internal stress fields caused by frozen-in small-scale spinodal-type inhomogeneities of the alloy components. A very important observation is the ditlerence between the samples deformed at Td<750°C

187 10 "1

.~oc, 16'.

+ Vo / / +XTo

///

/F

16 164

+7Zo

//V

./Iv

o

I

lO

100

6 (MPa)

(a) 101

Fig. 5. Dislocation configuration in the sample deformed by DRV. The deformation conditions were To = 600 °C and g.=l.7×10 3 s I The strain was 10%. (Magnification 12 000 × .) and those deformed at To > 750 °C. At high temperatures of deformation there exists a very small volume fraction of stable D03 particles, which dissolve at temperatures T > 740 °C, which is the solubility limit for tin in the quasi-binary C u 9 N i - S n diagram [9]. On the contrary, the density of these particles in samples deformed at lower temperatures is substantially higher (see Fig. 5). We suppose that the D03 particles have a retarding effect on the G B migration as well as on the redistribution of dislocations within individual grains. At Td > 700 °C second-phase particles dissolve substantially through the matrix and do not suppress the D R X and the dislocation redistribution in grains. T h e deformation process in the first temperature range 6 0 0 - 7 0 0 °C is characterized by higher values of both Q = 2.3 + 0.3 eV and n = 4.7 ± 0.3 compared with the second range ( 7 5 0 - 8 5 0 °C), for which Q = 1.9 ± 0.3 eV and n = 2.7 ± 0.1. T h e values were obtained by fitting to eqn. (1) (Fig. 6). This result indicates that hot deformation of this spinodal alloy can proceed more easily at temperatures TO> 750 °C at which the second-phase particles dissolve throughout the matrix. T h e effect of Td at different go on emax (plastic strain to fracture) is shown in Fig. 7. It is obvious that emax at all go used increases with temperature up to TO= 750 °C. We assume that the ductility is connected with the role of the second-phase particles in the GB sliding and crack nucleation.

2.2. The effect of the tin on dynamic softening T h e alloys with tin contents of 0, 0.5 and 1.2 at.% were deformed in tension. Examples of o - e

I

I +Vo

.~'ld' 1(

10'

I

I

10

lO0o~(MPa }

(b)

Fig. 6. The deformation data o.... To and g0 fitted with the help of eqn. (1): (a) temperatures of 600-700 °C; (b) temperatures of 750-850 °C.

0"30 / /A E to

/

0.20 / /

0.10

~t"

/ / ~

•

t

.. /

~/-/

//° / !/ / /

/ /~//

0

J

600

J 7OO

~

Td[°C ;

8OO

Fig. 7. The influence of T d o n tlae fracture ductility emax. All initial strain rates are included.

curves compared with that of the spinodal alloy are given in Fig. 8. For x = 0 and 0.5 at.% softening by D R V is most characteristic. Only at Td > 800 °C was softening by D R X observed for x = 0 at.%. On the contrary, softening by D R V is dominant for x = 1 . 2 at.% up to 700°C. At Td > 700 °C softening by D R X takes place. This difference in the softening mechanisms for different Cu9NixSn alloys is mainly due to the difference in the phase structure.

188

100

,

other copper-base solid solutions [10]. For x = 3.1 at.%, which is a two-phase alloy, particles dissolve around 740 °C and the value of 0o.2 is then comparable with that of other solid solutions. The values of both Q and n were calculated also for x = 0, 0.5 and 1.2 at.% by the same procedure used for the spinodal alloy and the values of Q and n are as follows.

4ol.7,10-3 ~-~ Td'750*C

-5 O_

x-3.1

50

x -12

I

01

I

0.10

(a) 20O

0.20

&

i

Lo" 12"10Z s 4 Td" 600°C

- ~ ~ 3 . 1 x -1.2

100

I

(b)

0

I

0-10

6

O20

Fig. 8. The comparison of the true stress-true strain curves with different contents x (at.%) of tin: (a) Td=750°C, ~, = 1.7 x 10 -3 s-1; (b) Td = 6 0 0 °C, g0 = 1.7 x 10 -4 s-1. 160

]

I

x (at.%)

Q (eV)

n

0 0.5 1.2

2.3 2.0 1.7

5.0 4.1 4.1

The results may be very important for providing a good choice of the proper technology of Cu9Ni6n alloys. The thermomechanical processing of such an alloy includes at present mainly the cold formation of a supersaturated solid solution obtained by the quenching of a solution-annealed product. By a proper use of our results, hot forming processing might also be used, enabling us to work with substantially smaller forces necessary for the process. The only condition is the omission of the region Td=350-500°C, for which the Cu9Ni6Sn is brittle according to the presence of coherent metastable particles (D022 structure, same stoichiometry). The latter makes the strength very high and/?max very low; see refs. 9 and 11.

I

6.,.1.7x10"3S -1

References

g

z 12o

80

600

700

800

T. I'c}

900

Fig. 9. The temperature dependence of the yield stress o0.2 for the Cu9Ni alloy with different tin contents (go = 1.7 x 10-3 s-I).

The different phase structure of the studied alloys is pronounced also in the different types of temperature dependences of the yield stress o0.2 (Fig. 9). The alloys with x = 0, 0.5 and 1.2 at.% are solid solutions for Td used in this experiment [9]. The dependence with a typical plateau is therefore observed at Td very near to those of

1 P. Kratochvil, J. Mencl, J. Pe~i6ka and S. N. Komnik, Acta Metall., 32(1984) 149. 2 S. Shektar, T. C. Lee and K. N. Subramanian, Phys. Status SolidiA, 91 (1985) 63. 3 M. Kato, S. Katsuta and A. Sato, Mater. Sci. Eng., 54 (1986) 95. 4 Z. K. Khamitov, J. Chumlyakov and A. D. Korotaev, Fiz. Met. Metalloved., 62 (1986) 362. 5 P. Kratochvil, M. Mandula, J. Mencl, J. Pe~i6ka and B. Smola, Phys. Status Solidi A, 104 (1987) 597. 6 K. Sato, S. Katsuta and M. Kato, Acta Metall., 36 (1988) 633. 7 P. Kratochvil, J. Mencl and A. Y. Abdellatief, in J. B. Bilde-Sorensen, N. Hansen, D. Juul Jensen, T. Leffers, H. Lilholt and O. B. Pedersen (eds.), Proc. lOth Riso Int. Symp. on Metallurgy and Material Science: Materials Architecture, National Laboratory, Roskilde, 1989, p.441. 8 E Garofalo, Fundamentals of Creep and Creep Rupture in Metals, New York, Macmillan, 1965. 9 J.T. Plewes, Metall. Trans. A, 6 (1975) 537. 10 L. Cizek, P. Kratochvil and B. Smola, J. Mater. Sci., 9 (1974) 1517. 11 M. Mandula, Diploma Thesis, Faculty of Mathematics and Physics, Charles University, Prague, 1987.