elevated temperature deformation behaviour of alumina-dispersed p/m copper

Materials Chemistry and Physics, 36 (1994) 311-376 371

Elevated temperature deformation behaviour of alumina-dispersed P/M copper

Govind, R. Balasubramaniam and G.S. Upadhyaya ~epa~ent of ~e~aiLu~.~ffl Engineeting Indian instate of Technology, Kanpur 208016 (India)

(Received May 15, 1993; accepted July 12, 1993)

Abstract

Alumina-dispersed copper was tested under compression in the temperature range 25-800 “C, at a constant strain rate of 10m3 s-‘. Tests were also performed at a constant temperature of 600 “C using different strain rates (lo-’ and 10m4 s-l). At all temperatures and strain rates, the stress-strain curves showed flow softening and the microstructures revealed the existence of shear bands. At temperatures of 500 “C and above, dynamic recrystallization was observed at shear bands. The formation of shear bands in the alloy is discussed based on the instability of texture during subsequent deformation.

Introduction

The strength of dispersion-hardened alloys at high temperatures, far in excess of that found in conventional alloys, renders them suitable for high-temperature ap- plications. The advantage of dispersion strengthening is that, in addition to increasing in strength, the matrix retains most of its other properties.

Dispersion-strengthened copper possesses good mechanical properties. Its creep resistance is better than other copper alloys. The dispersoids provide better mechanical properties, both at room and elevated temperatures, without exhibiting significant loss in electrical conductivi~.

During deformation of metals and alloys at elevated temperatures, two processes, namely deformation and recrystallization, operate simultaneously. The occurrence of recrystallization modifies the appearance of the flow curves produced at a constant strain rate [l]. At high strain rates, in the hot working range, the flow stress increases to a maximum, and thereupon, as a result of dynamic recrystallization, it drops to a value intermediate between the yield stress and peak stress. At low strain rates, in the hot working range, dynamic rec~stall~ation causes ovulations in the flow stress.

The presence of second-phase particles tends to stabilize the fine substructure formed as a result of deformation [2] and prevents grain boundary migration [3]. Such inhibition or retardation of dynamic recrystallization has been observed in austenite containing NbC particles and in Udimet 700 with y’ precipitates [4]. The present study is an attempt to understand the

role of finely distributed alumina particles in powder- metallurgically produced copper during high temperature deformation.

Experimental

The material for investigation (alumina-dispersed copper) was supplied, in the form of hot extruded rods of 12.5 mm diameter, by SCM Glidden Metals, USA. The Al,O, content in the two grades of ODS (oxide- dispersion-strengthened) copper, namely Glidcop Al- 20 and Al-60, were 0.9 and 2.7 vol.%, respectively. The size of alumina particles in the alloys, as reported by Stephens et al. [5], was in the range of 100 to 1000 A. These were prepared by internal oxidation through the P/M route as described by Nadkarni et al. [6]. A commercially pure copper rod was also investigated for comparative purposes.

The compression test specimens were made from the as-received (hot extruded) rods. The diameter of the specimens was 5.3 mm and the L/D ratio was kept at 1.5. The specimens were machined such that their ~ompre~ion axis was parallel to the longitudinal axis of the rod.

In the first set of experiments, compression tests were performed on the as-received ODS copper and commercially pure copper at different temperatures (25 to 800 “C), keeping strain rate constant at 10e3 s-l. In the next set of experiments, the test temperature was maintained constant at 600 “C and the tests were

0254-0584/94/$07.00 0 1994 Elsevier Sequoia. All rights reserved

372

a

k 300 .

,” 2 zi Qt 2 200 t

100

Oi I I f t I I I 1 _J 0 0.1 0.2 0.3 0.4

True Strain

Fig. 1. Compressive flow cmves for as-received Al-20 ahoy at different temperatures.

t

01 , 8 8 8 1 I I

0 0.10 0.20 0.30 0 True Strain

Fig. 2. Compressive flow curves for as-received Al-60 alloy at different temperatures.

performed using strain rates of 10e2 and low4 s-l. The tested specimens were sectioned longitudinally and were observed under an optical microscope after etching. One specimen from the as-received Al-60 rod was cut along the transverse direction and tested in compression

Fig. 3. Optical microstructures of as-received Al-20 alloy after compression testing at 600 “C corresponding to different strain rates: 10-a s-’ [(a), (c) and (d)] and 10m4 s-’ (b). (a) and (b) exhibit shear bands, (c) is an enlarged view of the shear band region and (d) is an enlarged view of the region away from the shear band.

at 600 “C at a strain rate of foe3 s-‘. Another set of as-received Al-60 alloy was annealed at 900 “C for 24 h and was tested in compression at 600 “C and a constant strain rate of lop3 s-‘. Microscopic examination of these specimens was also carried out to see whether such annealing aligns the axes of the matrix grains.

Results

Figures 1 and 2 show the true stress-strain curves for the as-received Al-20 and Al-60 ODS copper tested in compression in the temperature range 25-800 “C at a constant strain rate of lop3 s-l. The trend in the flow curves is similar at all temperatures, and shows a peak followed by a drop. Microstructural examination of tested specimens at all temperatures revealed the existence of shear bands. Typical examples of shear bands observed in the as-received Al-20 alloy after compression testing at 600 “C and strain rates of lo-* and lop4 s-l are given in Fig. 3(a) and 3(b) respectively. The shear band region in Fig. 3(a) is shown in enlarged view in Fig. 3(c), revealing the occurrence of dynamic recrystallization. Dynamic recrystallization was observed in both ODS coppers tested in compression at temperatures of 500 “C or above. Moreover, such a feature was observed only in the region of shear bands. Figure 3(d) shows the microstructure of a region away from the shear band. The presence of elongated grains in the microstructure reveals that recrystallization did not occur in this region. The microstructure of the specimens, interrupted near the peak stress, revealed shear bands with no sign of re~stallization. The deformation behaviour of both ODS alloys was similar. The flow stress level, however, was higher for Al-60 than for Al- 20.

The annealed (900 “C, 24 h) specimen of the Al-60 alloy when tested in compression at 600 “C using a strain rate of low3 s-’ showed a similar flow behaviour. In addition it exhibited shear bands giving evidence of dynamic recrystallization. A shear band is visible in Fig. 4(a). The enlarged view of shear band region, shown in Fig. 4(b), reveals the occurrence of dynamic re~~stallization in the shear bands. Regions removed from the shear bands did not show dynamic recrystallization (Fig. 4(c)). This indicates that the annealing did not align the grain axes.

Figure 5 shows the flow curves obtained in compression for the ODS alloys of both grades, i.e. Al-20 and Al-60, at 600 “C and strain rates of lo-* and 10F4 S -‘. The curves reveal that ODS alloys are sensitive to strain rate, such that the higher the strain rate, the higher the flow stress level. At any strain rate, the drop in stress is sharper in Al-20 than in Al-60. A

373

Fig. 4. Optical microstructures of annealed (950 “C for 24 h) Al-60 alloy after compression testing at 600 “C and strain rate lo-’ s-‘. (a) and (b) exhibit shear band regions, while (c) shows a region away from the shear bands.

close look at the Aow curves, including those corresponding to the 10m3 s-’ strain rate, shows that the behaviour for the Al-60 alloy follows a pattern, while this is not the case for Al-20. The anomaly in the latter case is, however, not very clear.

Figure 6 shows the variation in yield stress (0.2% offset) with temperature for commercially pure copper and alumina-dispersed copper. Pure copper shows a drastic drop in yield stress at temperatures of 500 to 800 “C, whereas alumina-dispersed copper maintains higher strength. The occurrence of dynamic recrystallization (Fig. 7(a)) and grain growth (Fig. 7(b) and 7(c)) was found to coincide with the drastic drop in yield stress of pure copper (Fig. 6). At 200 “C no grain growth (Fig. 7(d)) was noted.

374

2508

200

g

z" I-60, strain rate

?! 150 z

0, 2 I-

501 I I I I 0 0.05 0.10 0.15 0.20 0.25

True Strain

Fig. 5. Compressive flow curves for as-received ODS copper at different strain rates and at a temperature of 600 “C.

500, , Strain rate 1.0 ~16~ 5-l

I

Temperature, 'C

Fig. 6. Variation of yield stress of as-received copper and its ODS alloy as a function of temperature.

Discussion

Elevated temperature flow behaviour It is apparent from the results (Figs. 1 and 2) that

the flow curves in compression show work-softening at all the test temperatures. The softening takes place below 5% strain. The onset of softening is observed at lower strains with increasing temperature. The specimens in which straining was stopped near peak stress did not show any sign of recrystallization, and their microstructures revealed the presence of shear bands.

Fig. 7. Optical microstructures for pure copper after compression testing at different temperatures: (a) 500 “C, (b) 600 “C, (c) 800 “C and (d) 200 “C.

This implies that the flow softening in the present case is not a result of dynamic recrystallization, but rather is due to localized deformation in the alloy system.

315

There are many factors reported in the literature [7-121 which promote localized deformation in a material. These factors are: (1) non-operation of strengthening mechanisms during hot working; (2) existence of a temperature gradient within the material; (3) microstructural inhomogeneity; (4) presence of texture.

Each of them is discussed below in relation to the ODS copper alloys investigated here.

There are reports [7, 81 of localized deformation due to non-operation of the strengthening mechanisms during hot working. During hot working, the diffusion rates in alloys are generally observed to be several orders of magnitude higher than those simply maintained at high temperature. This may lead to the coarsening of fine dispersoids present in the matrix, which finally results in flow softening. This can be ruled out in the present case, since the A&O, dispersoids are very stable and the chances of their coarsening are negligibly small. Moreover, the shear bands were observed even at lower temperatures (25 “C and 200 “C), at which coarsening of Al,O, dispersoids is absent.

Semiatin et al. [ 131 reported that a small temperature gradient (10 “C) during testing can lead to localized flow in titanium aluminide, steel and Sic. However, the temperature gradient effect can also be ruled out in the present case, as the accuracy of temperature control (& 1 “C) during the test was sufficient to avoid any localized deformation. Furthermore, the alloy (Cu-A&O,) possesses fairly good thermal conductivity.

Microstructural inhomogeneity can also trigger localized flow. However, such inhomogeneities were not observed in the present study.

Another types of flow softening is that due to texture, which is operative at both cold and hot working temperatures. Aernoudt et ~2. [14] found that powder- metallurgically processed copper alloys containing 0.5 vol.% of fine Al,O, particles retained a strong room temperature deformation-induced texture, which they found was responsible for the mechanical anisotropy of the alloys. There are reports [15] of the presence of strong wire texture [{loo} (li2)] in the hot-rolled consolidated Cu-ThO, system. It seems that the ODS copper investigated in this study also possesses texture which is unstable under subsequent compression load- ing, thus leading to shear band formation. This was further confirmed by a compression test conducted on a specimen cut from the transverse direction, which did not show any shear band formation. Earlier experiments with longitudinal specimens, in which the compression axis was parallel to the extrusion direction, involved a change in the strain path as the elongated grains were being compressed. The recent experiment

on the transverse specimen, in which the compression axis was perpendicular to the extrusion direction, did not involve any change in the strain path as the elongated grains were in a stress state similar to that present during the earlier hot extrusion. Gill and coworkers [ll] found that when the textured alloy is subjected to strain path changes, it can promote the formation of shear bands by decreasing the mean Taylor factor (&) along that path. In the present case, when the textured alloy was subjected to a change in strain path, shear bands were formed because the texture could not respond instantaneously to the change in strain path.

The occurrence of dynamic recrystallization in the shear bands can be explained on the basis of conclusions drawn by Korbel et al. [16] in a study on plastic flow behaviour of copper single crystals. In an extensive study using transmission electron microscopy, they found that strain localization in essence provides the miso- rientation needed for the nucleation of recrystallized grains. This appears to be also valid in the present alloy systems.

Effect of strain rate An increase in the strain rate increases the flow

stress level during compression, which confirms that the flow stress in ODS copper is strain-rate sensitive. The relation between the strain rate (i) and the intrinsic material parameter (dislocation density and its average velocity) is given by the following relation [17]:

i=bpu (1)

where p is the density of mobile dislocations and V is the average dislocation velocity. The dislocation density increases with strain, and its average velocity (V) is a strong function of stress, as given by:

d = ( T/TfJrn ’ (2) where TV is the resolved shear stress corresponding to unit velocity, T is the resolved shear stress corresponding to velocity V, and m’ is an exponent describing the stress dependence of dislocation velocity. For metals and alloys with a low initial mobile dislocation density or with strong pinned dislocations, as is the situation in the present case, the only way that bpz? can achieve the imposed strain rate is for V to be high. This can only occur at high stress.

Effect of dispersoids An increase in the alumina content in the ODS alloy

leads to an increase in peak strain and thus helps to delay localized deformation. The sharp drop in yield stress with increasing temperature in the case of pure copper (Fig. 6) is the result of dynamic recrystallization

376

and grain growth. In alumina-dispersed copper, dispersoids stabilize the substructure formed as a result of deformation, thus inhibiting dynamic recrystallization and maintaining high strengths at all temperatures.

1) The flow softening observed in the present study can be attributed to textural instability during subsequent changes in strain path.

2) Dispersoids delay flow localization in alumina- dispersed P/M copper.

3) Alumina-dispersed copper maintains high yield stresses at high temperatures, whereas the yield stress of commercially pure copper decreases drastically with increasing temperature.

References

H.J. McQueen and J.J. Jonas, in R.J. Arsenault (ed.), 7’rea~ire on Materials Science and Technology, Vol. 6, Academic Press, New York, 1975, p. 393. F.J. Humphreys and J.W. Martin, Phil. Mug., 17 (1968) 365. I. Baker and J.W. Martin, in N. Hansen, A.R. Jones and T. Leffers (eds.), Proc. 1st Rird Int. Symp. Metallurgy and Materials Science, Risd National Laboratory, Roskilde, 1980, p. 27. J.M. Oblak and W.A. Owczarski, MetaN. Trans., 3 (1972) 617.

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J.J. Stephens, J.A. Romero and C.R. Hills, in H.J. Cialoni, M.E. Blum, G.W.E. Johnson and G.F. Vandervoort (eds.), Microstructural Science, Vol. 16, International Metallographic Society, Columbus, OH and ASM International, Metals Park, OH, 1988, p. 245. A.V. Nadkami, E. Klar and W.M. Shafer, Met. Eng. Q., August (1976) 10. J.J. Jonas and M.J. Luton,Advances in Deformation Processing, Plenum Press, New York, 1974, p. 215. I. Weiss and J.J. Jonas, Metall. Trans., 1lA (1980) 215. T. Chandra, I. Weiss and J.J. Jonas, Met. Sci. J., 16 (1982) 97. T. Chandra, I. Weiss and J.J. Jonas, Can. Metall. Q., 20 (1981) 421. J. Gill Sevillano, P.V. Houtte and E. Aernoudt, fiogr. Mater. Sci., 25 (1981) 69. J.J. Jonas and S.L. Semiatin, Formability and workabili@ of metals: Plastic instability and flow localization, ASM, Metals Park, OH, 1984, p. 43. S.L. Semiatin, S.I. Oh and J.J. Jonas, Metall. Trans., 23A (1992) 963. E. Aernoudt, R. Garvin-Salazar, P.V. Houtte, M. Follon and P. Neutjens, in P. Ramakrishnan (ed.), PowderMetallurgy and Related High Temperature Materials, Trans Tech Publications, Aerdennannsdorf, 1987, p. 373. E.R. Kimmel, L.G. Peterson and R.A. Queeney, Thermo- mechanical Processing of Oxide Dkpersion-strengthened Coppec The International Copper Research Association, New York, January, 1971. A. Korbel, W. Bochniak, L. Blaz and J.D. Embury, Met. Sci., I8 (1984) 216. G.E. Dieter, Mechanical Metallurgy McGraw-Hill, London, 1988, p. 200.

elevated temperature deformation behaviour of alumina-dispersed p/m copper

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