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M in-100 superalloy
Journal of Mechanical Working Technology, 3 (1980) 331--339 © Elsevier Scientific Publishing Company, Amsterdam -- Printed in The Netherlands
331
T H E E F F E C T O F T E M P E R A T U R E AND S T R A I N R A T E ON T H E S U P E R P L A S T I C B E H A V I O U R O F P/M IN-100 S U P E R A L L O Y
A.A. AFONJA* Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, Mass. (U.S.A.) (Received February 14, 1979; accepted in revised form June 15, 1979)
Industrial S u m m a r y IN-100 is a very important superalloy, used in such industrial high-temperature, highstrength, applications as gas turbine blades, nozzles and wheels. Such parts are normally produced from cast alloy and complex heat-treatment is required to obtain the desirable fine-grained structure. An alloy of superior microstructure can be produced by powder metallurgy techniques. This greatly simplifies subsequent processing and enhances the hot ductility of the alloy within a certain range of temperature and strain rate. This paper reports the results of an investigation designed to establish the optimum range of conditions within which a low-carbon P/M IN-100 superalloy is superplastic. Enhanced hot ductility of the alloy is of considerable practical importance since it extends the range of possible manufacturing techniques to such precision processes as pressing, extrusion and forging.
Introduction IN-100 is a nickel-base p r e c i p i t a t i o n - h a r d e n a b l e superalloy d e v e l o p e d basically as a cast alloy. T h e high h o t strength and specific strength m a k e it particularly suitable f o r such h i g h - t e m p e r a t u r e applications as gas t u r b i n e blades, nozzles and wheels. H o w e v e r , the c o m p l e x i t y o f the alloy c o m p o s i t i o n is such t h a t severe segregation occurs during casting, resulting in p o o r h o t ductility. Castings o f the alloy are usually coarse-grained and n u m e r o u s heat t r e a t m e n t and working cycles are required to h o m o g e n i s e and refine the microstructure. However, a u n i f o r m , ultra-fine-grained alloy with i m p r o v e d h o t d u c t i l i t y can be p r o d u c e d b y p o w d e r m e t a l l u r g y techniques. T h e P/M alloy is also super-ductile if w o r k e d u n d e r certain c o n d i t i o n s o f t e m p e r a t u r e and strain rate. This has immense practical implications since it widens the scope o f possible m a n u f a c t u r i n g processes t o such precision processes as forging and extrusion. Smith et al. [1, 2] studied the h o t d e f o r m a b i l i t y o f P/M IN-100 m a d e b y e x t r u d i n g pre-alloyed argon gas a t o m i z e d p o w d e r at 1177°C. T h e alloy was superplastic b e t w e e n 927°C and 1093°C and strain rates o f 10 -5 and 10 -3 s -~. Present address: Metallurgical and Materials Engineering, University of Ife, Nigeria.
332
Strains in excess of 500% were achieved. Athey and Moore [3] reported superplastic strains of over 1300% and described a method of production of full-scale turbine parts. Moscowitz et al. [4] studied the microstructure of IN-100 alloys produced by hot extrusion of pre~alloyed powders produced by three techniques: inert-gas atomization, vacuum atomization, and the rotating electrode process. The product had ultra-fine structure of gamma-prime particles of the order of 0.2 tam uniformly dispersed in a gamma matrix of grain size 8 tam. On the basis of the room- and elevated-temperature mechanical properties, the alloy made from powder produced by the rotating electrode process showed the greatest propensity for superplastic deformation in the range of temperature and strain rate, 982°C--1149°C and 10-3--10 s- ~, respectively. There was also a substantial growth in the average grain size of the gamma-prime particles. Similar results were obtained by Kobayashi [5] in his study of the post-superplastic deformation microstructure. In addition to gamma-prime grain coarsening, Kobayashi observed large gamma-prime free areas in the gamma matrix, mainly at grain boundaries normal to the axis of the applied stress. These morphological changes were attributed to stressinduced diffusional creep. The investigations carried out so far on the superplastic behaviour of P/M IN-100 have been on materials of varying chemical and microstructural properties under varying experimental conditions. Also, different methods have been used in determining the strain-rate sensitivity index. There is no c o m m o n basis therefore for comparing and correlating the results of these investigations. The object of this study was to determine the limits of temperature and strain rate within which a P/M IN-100 superalloy may be expected to deform superplastically. Experimental procedure Material Pre-alloyed IN-100 powder was prepared by the N M Powder Rotating Electrode Process [6] and extruded in an evacuated stainless steel container at 1150°C and at an extrusion ratio of 30:1. The chemical composition of the alloy is presented in Table 1. The as-extruded alloy had a fine-grained, equiaxed, uniform structure of extremely fine gamma-prime particles dispersed in a g a m m a matrix (Fig. I). The grain size of the g a m m a matrix was measured using the Hilliard Circle technique on a number of photomicrographs and was 15 tam. The average unit size of the gamma-prime particles was estimated to be about 0.3 tam or less.Threaded-end tensile specimens of 12.5 m m gauge length and 3.2 m m diameter were machined from the extruded bars and then polished.
Equipment The test equipment consisted of a tensile testing machine equipped with a variable speed drive, capable of speeds equivalent to initial strain rates in the
333 TABLE 1 Chemical composition of P/M IN-100 superalloy Element
wt. %
Element
wt. %
Cr Co Mo AI Ti C B
9.20 14.90 2.90 5.80 4.56 0.02 0.02
Zr V Fe Mn Si S Ni
0.11 0.97 0.55 0.13 0.10 0.006 Bal.
Fig. 1. Microstrueture of P/M IN-100. range 1 0 - ~ - - 1 0 -~ s -~. The applied load was m e a s u r e d b y a load cell having a linear response over the a n t i c i p a t e d load range. T h e signal f r o m the load cell was fed t o an X- Y recorder. The cross-head speed was m e a s u r e d with an a t t a c h e d m i c r o m e t e r c l o c k gauge. T h e e q u i p m e n t was fitted with a f u r n a c e having a m a x i m u m t e m p e r a t u r e o f 1200°C, and i n s t r u m e n t a t i o n f o r temperature c o n t r o l to within -+ 2°C. Provision was m a d e also f o r the circulation o f
334 pre-heated purified argon. A device was incorporated in the equipment to determine when the specimen fractures.
Test procedure Tensile tests were carried out on the specimens at constant cross-head speeds in the temperature range 900--1150°C. The strain and the strain rate were determined from measurements of the average cross-head speed over the test period. Tests were carried out in either an atmosphere of purified argon, or in air, and over a wide range of initial strain rates between 10 -3 and 10 -6 s-'. All the tests except those at 1050°C were stopped after a strain of about 25% had been achieved. The specimens were then quenched and the gauge dimensions measured to cross-check the strain values derived from the measurement of the cross-head travel. The tests at 1050°C were continued until the specimens fractured in order to determine the m a x i m u m total elongation. From Kobayashi's results [5] and also the results of preliminary tests carried out in this investigation, it appears t h a t m a x i m u m elongation occurs at around this temperature. A typical fractured specimen is shown in Fig. 2.
Fig. 2. Specimen tested in purified argon at 1050°C and at a strain rate of 2.4 x 10 -5 s-'. Discussion o f results
Dependence o f strain-rate sensitivity on strain Stress-strain curves were plotted for the various temperatures and strain rates investigated. Typical plots are shown in Fig. 3. Clearly, the strain rate has
335
200 _
o ^. =~3.TxlO-#sec~1 • ~: =4~xlu-~ × $ =2.~x10-6o ,,
~
,' ~,~ ~ :~x~O- 7 ,, l
"~:
E
5J" x 10-
"
o
Z
x
100
lb
2o
3b
4o
strain (%)
Fig. 3. Stress--strain curves for specimens tested at 1050°C and at a various strain rates ~.
o 10% strain • 20% ,, o 40% " 1% "
c~
o----''° j o / ~ ''~'~
o/___
z Z
0
. . . . . . . .
-7
#
-5
. . . . . .
~
-5
log strain
. . . . . .
r
-4 rate (sec -1)
. . . . . . . . .
-3
Fig. 4. T h e e f f e c t o f s t r a i n o n t h e s h a p e o f t h e log s t r e s s - - l o g s t r a i n - r a t e c u r v e .
a significant effect on both the deformation stress and the maximum elongation achievable. Plots of log stress versus log strain-rate were made from data obtained from the stress-strain curves for various values of strain; a typical plot is shown in Fig. 4. The shape of the log stress--log strain-rate curve is clearly a function of the strain. The S-shape of the curve at 1% strain is similar to those obtained by Holt and Backofen [7] and also Rai and Grant [8] for AI-Cu eutectic superplastic alloy. Geckinli and Barrett [9] also obtained a similar S-shape with Sn-Pb alloy. The figure shows that at least 10% strain is required to obtain a linear shape characteristic of superplastic behaviour, hence subsequent plots were based on 15% strain. The strain rate sensitivity was determined from measurements of the slope of the log stress--log strain-rate curve.
The effect o f temperature and strain rate on strain-rate sensitivity Figure 5 shows the log stress--log strain-rate plots for various temperatures.
336
_5
% z
z 4 ~y
~
c c (a,r) C (argon) C x : 1150°C
2 log strain rate (sec-I)
Fig. 5. Log stress--log strain-rate curves.
Some of the tests at 1050°C were carried o u t both in purified argon and in air under otherwise identical conditions in order to determine the effect of oxidation on the strain-rate-sensitivity index. There was no significant effect, as can be seen from Fig. 5. This is in agreement with the results of Kobayashi [5]. The strain-rate sensitivity varies between 0.35 and 0.8 for the temperature range 1000°C--1150°C, the maximum being at ll00°C. At strain rates above 10 -4 s-', there is an abrupt change in the slope of all the curves, indicating a change in the mechanism of superplastic deformation. At low strain rates there is no change in the slope of the curves even at values of the order of 10 -6 s-', contrary to the results of Holt and Backofen [7] which showed a decrease in slope. The S-shape obtained b y the latter researchers is apparently not a characteristic of the deformation, but a result of the method of determination of the strain-rate sensitivity.
The effect o f strain rate on elongation The effect of strain rate on the total elongation before fracture is shown in Fig. 6. The figure shows also the effect of the test environment. The strain at failure is considerably higher for the tests carried o u t in purified argon when compared with those carried o u t in air, due apparently to oxidation. Further, the total elongation increases as the strain rate decreases. This is in agreement with the results of numerous investigations on superplastic materials [9, 10]. However, at strain rates below a b o u t 10 -5 s-', there is a sharp fall in elongation even for tests carried o u t in purified argon. The gamma and gamma-prime grain sizes of the specimen tested at 10 -6 s -1 were determined and found to have increased to a b o u t 25 pm. Kobayashi [5] found a similar drop in elongation b u t with respect to temperature, and also evidence of grain coarsening and evidence of intergranular and surface oxidation in the microstructure of samples tested at the higher temperatures and lower strain rates. He also observed that part of the gamma-prime phase was re-dissolved under these conditions. However, Rai and Grant [8] found no such reversal in elongation with increasing temperature or decreasing strain rates, even at strain rates as low as 10 -~ s -1 and with significant grain growth. This discrepancy is due probably to differences in the oxidation characteristics of P/M IN-100 and AI-Cu eutectic alloy.
337 8C0
j
o~
/\
00j
o argon
ioo / .
-6
.
-5 log strain rate (sec-1)
-4
-3
Fig. 6. The effect of strain rate and te~t atmosphere on the total elongation.
The mechanism o f plastic deformation The activation energies for deformation at constant stress were determined for various stress levels from plots of log strain-rate versus the reciprocal of absolute temperature from data taken from Fig. 5, in accordance with the relationship =A exp(-Q/RT) where ~ is the strain rate, Q the activation energy, T the absolute temperature, R the universal gas constant, and A a constant. The curves are shown in Fig. 7. The average value of Q for the five stress levels is 470 kJ/mol. This compares with the value obtained by Reichman and Smythe [2]. The high value of the strain rate sensitivity (0.8) suggests that one or more of the diffusional creep modes -- bulk, grain boundary or dislocation diffusion -- are the dominant mechanisms in the superplastic region; the activation energy of 470 kJ/mol is considerably higher, however, than that of bulk diffusion of aluminium in nickel, which was estimated to be 250 kJ/mol [11]. The activation energies for grain boundary and dislocation diffusion would be even lower. Further, Kobayashi [5], in his study of the post-deformation microstructure of P/M IN-100 superplastic alloy, found that the gamma matrix grains remained reasonably equiaxed, which is consistent with either Nabarro--Herring or Coble creep. The absence of grain elongation suggests that grain boundary sliding is the rate-controlling mechanism, but neither this nor diffusional creep can account for the large strains achieved. Kobayashi [5] studied the microstructural changes which occurred in P/M IN-100 superplastic alloy after varying degrees of straining and found substantial gamma-prime grain growth from the initial size of 0.5 pm or less, to about 20/~m, and also large precipitate-free areas in the gamma matrix.
338 o s t r e s s = 200 M N / m 2 • -2
~
= 100
n
.
=
50
',
•
,,
=
20
"
.c -4
_8,-5 -6
7"0 7.1
7:3
IO00/T
¢5
7:7
7.9
Fig. 7. Strain rate ~ versus reciprocal o f the absolute t e m p e r a t u r e T for c o n s t a n t initial flow-stresses.
He also observed that the gamma-prime grains elongated in the direction of the stress axis at the lower strains,the directionality disappearing as the strain increased. It appears, therefore, that diffusional creep is the dominant ratecontrolling mechanism in the initialstages of superplastic deformation. In the absence of diffusion data it is difficultto determine whether it occurs in the gamma-prime lattice,or at the grain boundary, or both. The fact that the directionality in the gamma-prime phase is lost as strain increases, suggests that grain-boundary sliding becomes the predominant mechanism. The grain growth in the gamma-prime phase eventually reduces the propensity of the alloy to continue to deform superplastically and since it is both temperature and strain-rate dependent, it is probably partially responsible for the reduction in elongation at the higher temperatures and lower strain rates, even when the specimens were tested in purified argon. Although there was no evidence of internal oxidation of the specimens, there was substantial surface oxidation and this m a y have caused surface microcracks and, consequently, premature failure.
Conclusions The foregoing results show that ultra-fine P/M IN-100 super alloy can deform superplastically between 1000°C and 1150°C and at strain rates below 10 -4 s-~; the latter m a y change of course with chemical composition, grain size, etc. There is no evidence of the S-shape of the log stress-logstrain rate curve, even at very low strain rates, except at very low strains when the deformation is inhomogeneous. The total strain before fracture depends on the test conditions. High temperature and low strain rate cause rapid grain growth of the gamma-prime precipitate which eventually stops superplastic deformation and causes premature fracture. Oxidation under these conditions may also have a negative dominant or contributory effect on elongation. The results of this investigation and previous studies suggest that the mechanism of
339
superplastic deformation is a complex interaction between grain-boundary sliding and volume or interface diffusional creep, occurring mainly in the gamma-prime phase. It is possible, however, that other mechanisms may be involved.
Acknowledgements Part of the investigation reported in this paper was carried out while the author was a Fulbright-Hays Visiting Scientist in the Department of Materials Science and Engineering, Massachusetts Institute of Technology. He would like to express his profound gratitude to the Institute and especially to Professor Nicholas Grant, the Director of the Center for Materials Science and Engineering, for providing the equipment and materials, and also to the Council for International Exchange of Scholars for providing the funding.
References 1 S.H. Reichman, B.W. Castledine and J.W. Smythe, SAE Congress, Detroit, Michigan, 1970. 2 S.H. Reichman and J.W. Smythe, Int. J. Powder Met., 6 (1970) 65. 3 J.B. Moore and R.L. Athey, 18th Sagamore Army Materials Conference, Syracuse Univ. Press, 1972, p. 281. 4 L.N. Moskowitz, R.M. Pelloux and N.J. Grant, Proc. 2nd Int. Conf. on Superalloys, Champion, Penn., 1972, Z-1. 5 Y. Kobayashi, M.S. Thesis, Dept. of Matl. Sc. and Eng., Mass. Inst. of Tech., 1974. 6 A.R. Kaufman, U.S. Patent No. 3,099,041, (1963). 7 D.L. Holt and W.A. Baekofen, Trans. ASM, 59 (1966) 755. 8 G. Rai and N.J. Grant, Met. Trans. A, 6A (1975) 385. 9 A.E. Geckinli and G.R. Barrett, Scripta Met., 8 (1974) 115. 10 D.A. Woodford, Trans Quart. ASM, 62 (1969) 291. 11 R. Swalin and A. Martin, Trans. AIME, 206 (1956) 567.
331
T H E E F F E C T O F T E M P E R A T U R E AND S T R A I N R A T E ON T H E S U P E R P L A S T I C B E H A V I O U R O F P/M IN-100 S U P E R A L L O Y
A.A. AFONJA* Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, Mass. (U.S.A.) (Received February 14, 1979; accepted in revised form June 15, 1979)
Industrial S u m m a r y IN-100 is a very important superalloy, used in such industrial high-temperature, highstrength, applications as gas turbine blades, nozzles and wheels. Such parts are normally produced from cast alloy and complex heat-treatment is required to obtain the desirable fine-grained structure. An alloy of superior microstructure can be produced by powder metallurgy techniques. This greatly simplifies subsequent processing and enhances the hot ductility of the alloy within a certain range of temperature and strain rate. This paper reports the results of an investigation designed to establish the optimum range of conditions within which a low-carbon P/M IN-100 superalloy is superplastic. Enhanced hot ductility of the alloy is of considerable practical importance since it extends the range of possible manufacturing techniques to such precision processes as pressing, extrusion and forging.
Introduction IN-100 is a nickel-base p r e c i p i t a t i o n - h a r d e n a b l e superalloy d e v e l o p e d basically as a cast alloy. T h e high h o t strength and specific strength m a k e it particularly suitable f o r such h i g h - t e m p e r a t u r e applications as gas t u r b i n e blades, nozzles and wheels. H o w e v e r , the c o m p l e x i t y o f the alloy c o m p o s i t i o n is such t h a t severe segregation occurs during casting, resulting in p o o r h o t ductility. Castings o f the alloy are usually coarse-grained and n u m e r o u s heat t r e a t m e n t and working cycles are required to h o m o g e n i s e and refine the microstructure. However, a u n i f o r m , ultra-fine-grained alloy with i m p r o v e d h o t d u c t i l i t y can be p r o d u c e d b y p o w d e r m e t a l l u r g y techniques. T h e P/M alloy is also super-ductile if w o r k e d u n d e r certain c o n d i t i o n s o f t e m p e r a t u r e and strain rate. This has immense practical implications since it widens the scope o f possible m a n u f a c t u r i n g processes t o such precision processes as forging and extrusion. Smith et al. [1, 2] studied the h o t d e f o r m a b i l i t y o f P/M IN-100 m a d e b y e x t r u d i n g pre-alloyed argon gas a t o m i z e d p o w d e r at 1177°C. T h e alloy was superplastic b e t w e e n 927°C and 1093°C and strain rates o f 10 -5 and 10 -3 s -~. Present address: Metallurgical and Materials Engineering, University of Ife, Nigeria.
332
Strains in excess of 500% were achieved. Athey and Moore [3] reported superplastic strains of over 1300% and described a method of production of full-scale turbine parts. Moscowitz et al. [4] studied the microstructure of IN-100 alloys produced by hot extrusion of pre~alloyed powders produced by three techniques: inert-gas atomization, vacuum atomization, and the rotating electrode process. The product had ultra-fine structure of gamma-prime particles of the order of 0.2 tam uniformly dispersed in a gamma matrix of grain size 8 tam. On the basis of the room- and elevated-temperature mechanical properties, the alloy made from powder produced by the rotating electrode process showed the greatest propensity for superplastic deformation in the range of temperature and strain rate, 982°C--1149°C and 10-3--10 s- ~, respectively. There was also a substantial growth in the average grain size of the gamma-prime particles. Similar results were obtained by Kobayashi [5] in his study of the post-superplastic deformation microstructure. In addition to gamma-prime grain coarsening, Kobayashi observed large gamma-prime free areas in the gamma matrix, mainly at grain boundaries normal to the axis of the applied stress. These morphological changes were attributed to stressinduced diffusional creep. The investigations carried out so far on the superplastic behaviour of P/M IN-100 have been on materials of varying chemical and microstructural properties under varying experimental conditions. Also, different methods have been used in determining the strain-rate sensitivity index. There is no c o m m o n basis therefore for comparing and correlating the results of these investigations. The object of this study was to determine the limits of temperature and strain rate within which a P/M IN-100 superalloy may be expected to deform superplastically. Experimental procedure Material Pre-alloyed IN-100 powder was prepared by the N M Powder Rotating Electrode Process [6] and extruded in an evacuated stainless steel container at 1150°C and at an extrusion ratio of 30:1. The chemical composition of the alloy is presented in Table 1. The as-extruded alloy had a fine-grained, equiaxed, uniform structure of extremely fine gamma-prime particles dispersed in a g a m m a matrix (Fig. I). The grain size of the g a m m a matrix was measured using the Hilliard Circle technique on a number of photomicrographs and was 15 tam. The average unit size of the gamma-prime particles was estimated to be about 0.3 tam or less.Threaded-end tensile specimens of 12.5 m m gauge length and 3.2 m m diameter were machined from the extruded bars and then polished.
Equipment The test equipment consisted of a tensile testing machine equipped with a variable speed drive, capable of speeds equivalent to initial strain rates in the
333 TABLE 1 Chemical composition of P/M IN-100 superalloy Element
wt. %
Element
wt. %
Cr Co Mo AI Ti C B
9.20 14.90 2.90 5.80 4.56 0.02 0.02
Zr V Fe Mn Si S Ni
0.11 0.97 0.55 0.13 0.10 0.006 Bal.
Fig. 1. Microstrueture of P/M IN-100. range 1 0 - ~ - - 1 0 -~ s -~. The applied load was m e a s u r e d b y a load cell having a linear response over the a n t i c i p a t e d load range. T h e signal f r o m the load cell was fed t o an X- Y recorder. The cross-head speed was m e a s u r e d with an a t t a c h e d m i c r o m e t e r c l o c k gauge. T h e e q u i p m e n t was fitted with a f u r n a c e having a m a x i m u m t e m p e r a t u r e o f 1200°C, and i n s t r u m e n t a t i o n f o r temperature c o n t r o l to within -+ 2°C. Provision was m a d e also f o r the circulation o f
334 pre-heated purified argon. A device was incorporated in the equipment to determine when the specimen fractures.
Test procedure Tensile tests were carried out on the specimens at constant cross-head speeds in the temperature range 900--1150°C. The strain and the strain rate were determined from measurements of the average cross-head speed over the test period. Tests were carried out in either an atmosphere of purified argon, or in air, and over a wide range of initial strain rates between 10 -3 and 10 -6 s-'. All the tests except those at 1050°C were stopped after a strain of about 25% had been achieved. The specimens were then quenched and the gauge dimensions measured to cross-check the strain values derived from the measurement of the cross-head travel. The tests at 1050°C were continued until the specimens fractured in order to determine the m a x i m u m total elongation. From Kobayashi's results [5] and also the results of preliminary tests carried out in this investigation, it appears t h a t m a x i m u m elongation occurs at around this temperature. A typical fractured specimen is shown in Fig. 2.
Fig. 2. Specimen tested in purified argon at 1050°C and at a strain rate of 2.4 x 10 -5 s-'. Discussion o f results
Dependence o f strain-rate sensitivity on strain Stress-strain curves were plotted for the various temperatures and strain rates investigated. Typical plots are shown in Fig. 3. Clearly, the strain rate has
335
200 _
o ^. =~3.TxlO-#sec~1 • ~: =4~xlu-~ × $ =2.~x10-6o ,,
~
,' ~,~ ~ :~x~O- 7 ,, l
"~:
E
5J" x 10-
"
o
Z
x
100
lb
2o
3b
4o
strain (%)
Fig. 3. Stress--strain curves for specimens tested at 1050°C and at a various strain rates ~.
o 10% strain • 20% ,, o 40% " 1% "
c~
o----''° j o / ~ ''~'~
o/___
z Z
0
. . . . . . . .
-7
#
-5
. . . . . .
~
-5
log strain
. . . . . .
r
-4 rate (sec -1)
. . . . . . . . .
-3
Fig. 4. T h e e f f e c t o f s t r a i n o n t h e s h a p e o f t h e log s t r e s s - - l o g s t r a i n - r a t e c u r v e .
a significant effect on both the deformation stress and the maximum elongation achievable. Plots of log stress versus log strain-rate were made from data obtained from the stress-strain curves for various values of strain; a typical plot is shown in Fig. 4. The shape of the log stress--log strain-rate curve is clearly a function of the strain. The S-shape of the curve at 1% strain is similar to those obtained by Holt and Backofen [7] and also Rai and Grant [8] for AI-Cu eutectic superplastic alloy. Geckinli and Barrett [9] also obtained a similar S-shape with Sn-Pb alloy. The figure shows that at least 10% strain is required to obtain a linear shape characteristic of superplastic behaviour, hence subsequent plots were based on 15% strain. The strain rate sensitivity was determined from measurements of the slope of the log stress--log strain-rate curve.
The effect o f temperature and strain rate on strain-rate sensitivity Figure 5 shows the log stress--log strain-rate plots for various temperatures.
336
_5
% z
z 4 ~y
~
c c (a,r) C (argon) C x : 1150°C
2 log strain rate (sec-I)
Fig. 5. Log stress--log strain-rate curves.
Some of the tests at 1050°C were carried o u t both in purified argon and in air under otherwise identical conditions in order to determine the effect of oxidation on the strain-rate-sensitivity index. There was no significant effect, as can be seen from Fig. 5. This is in agreement with the results of Kobayashi [5]. The strain-rate sensitivity varies between 0.35 and 0.8 for the temperature range 1000°C--1150°C, the maximum being at ll00°C. At strain rates above 10 -4 s-', there is an abrupt change in the slope of all the curves, indicating a change in the mechanism of superplastic deformation. At low strain rates there is no change in the slope of the curves even at values of the order of 10 -6 s-', contrary to the results of Holt and Backofen [7] which showed a decrease in slope. The S-shape obtained b y the latter researchers is apparently not a characteristic of the deformation, but a result of the method of determination of the strain-rate sensitivity.
The effect o f strain rate on elongation The effect of strain rate on the total elongation before fracture is shown in Fig. 6. The figure shows also the effect of the test environment. The strain at failure is considerably higher for the tests carried o u t in purified argon when compared with those carried o u t in air, due apparently to oxidation. Further, the total elongation increases as the strain rate decreases. This is in agreement with the results of numerous investigations on superplastic materials [9, 10]. However, at strain rates below a b o u t 10 -5 s-', there is a sharp fall in elongation even for tests carried o u t in purified argon. The gamma and gamma-prime grain sizes of the specimen tested at 10 -6 s -1 were determined and found to have increased to a b o u t 25 pm. Kobayashi [5] found a similar drop in elongation b u t with respect to temperature, and also evidence of grain coarsening and evidence of intergranular and surface oxidation in the microstructure of samples tested at the higher temperatures and lower strain rates. He also observed that part of the gamma-prime phase was re-dissolved under these conditions. However, Rai and Grant [8] found no such reversal in elongation with increasing temperature or decreasing strain rates, even at strain rates as low as 10 -~ s -1 and with significant grain growth. This discrepancy is due probably to differences in the oxidation characteristics of P/M IN-100 and AI-Cu eutectic alloy.
337 8C0
j
o~
/\
00j
o argon
ioo / .
-6
.
-5 log strain rate (sec-1)
-4
-3
Fig. 6. The effect of strain rate and te~t atmosphere on the total elongation.
The mechanism o f plastic deformation The activation energies for deformation at constant stress were determined for various stress levels from plots of log strain-rate versus the reciprocal of absolute temperature from data taken from Fig. 5, in accordance with the relationship =A exp(-Q/RT) where ~ is the strain rate, Q the activation energy, T the absolute temperature, R the universal gas constant, and A a constant. The curves are shown in Fig. 7. The average value of Q for the five stress levels is 470 kJ/mol. This compares with the value obtained by Reichman and Smythe [2]. The high value of the strain rate sensitivity (0.8) suggests that one or more of the diffusional creep modes -- bulk, grain boundary or dislocation diffusion -- are the dominant mechanisms in the superplastic region; the activation energy of 470 kJ/mol is considerably higher, however, than that of bulk diffusion of aluminium in nickel, which was estimated to be 250 kJ/mol [11]. The activation energies for grain boundary and dislocation diffusion would be even lower. Further, Kobayashi [5], in his study of the post-deformation microstructure of P/M IN-100 superplastic alloy, found that the gamma matrix grains remained reasonably equiaxed, which is consistent with either Nabarro--Herring or Coble creep. The absence of grain elongation suggests that grain boundary sliding is the rate-controlling mechanism, but neither this nor diffusional creep can account for the large strains achieved. Kobayashi [5] studied the microstructural changes which occurred in P/M IN-100 superplastic alloy after varying degrees of straining and found substantial gamma-prime grain growth from the initial size of 0.5 pm or less, to about 20/~m, and also large precipitate-free areas in the gamma matrix.
338 o s t r e s s = 200 M N / m 2 • -2
~
= 100
n
.
=
50
',
•
,,
=
20
"
.c -4
_8,-5 -6
7"0 7.1
7:3
IO00/T
¢5
7:7
7.9
Fig. 7. Strain rate ~ versus reciprocal o f the absolute t e m p e r a t u r e T for c o n s t a n t initial flow-stresses.
He also observed that the gamma-prime grains elongated in the direction of the stress axis at the lower strains,the directionality disappearing as the strain increased. It appears, therefore, that diffusional creep is the dominant ratecontrolling mechanism in the initialstages of superplastic deformation. In the absence of diffusion data it is difficultto determine whether it occurs in the gamma-prime lattice,or at the grain boundary, or both. The fact that the directionality in the gamma-prime phase is lost as strain increases, suggests that grain-boundary sliding becomes the predominant mechanism. The grain growth in the gamma-prime phase eventually reduces the propensity of the alloy to continue to deform superplastically and since it is both temperature and strain-rate dependent, it is probably partially responsible for the reduction in elongation at the higher temperatures and lower strain rates, even when the specimens were tested in purified argon. Although there was no evidence of internal oxidation of the specimens, there was substantial surface oxidation and this m a y have caused surface microcracks and, consequently, premature failure.
Conclusions The foregoing results show that ultra-fine P/M IN-100 super alloy can deform superplastically between 1000°C and 1150°C and at strain rates below 10 -4 s-~; the latter m a y change of course with chemical composition, grain size, etc. There is no evidence of the S-shape of the log stress-logstrain rate curve, even at very low strain rates, except at very low strains when the deformation is inhomogeneous. The total strain before fracture depends on the test conditions. High temperature and low strain rate cause rapid grain growth of the gamma-prime precipitate which eventually stops superplastic deformation and causes premature fracture. Oxidation under these conditions may also have a negative dominant or contributory effect on elongation. The results of this investigation and previous studies suggest that the mechanism of
339
superplastic deformation is a complex interaction between grain-boundary sliding and volume or interface diffusional creep, occurring mainly in the gamma-prime phase. It is possible, however, that other mechanisms may be involved.
Acknowledgements Part of the investigation reported in this paper was carried out while the author was a Fulbright-Hays Visiting Scientist in the Department of Materials Science and Engineering, Massachusetts Institute of Technology. He would like to express his profound gratitude to the Institute and especially to Professor Nicholas Grant, the Director of the Center for Materials Science and Engineering, for providing the equipment and materials, and also to the Council for International Exchange of Scholars for providing the funding.
References 1 S.H. Reichman, B.W. Castledine and J.W. Smythe, SAE Congress, Detroit, Michigan, 1970. 2 S.H. Reichman and J.W. Smythe, Int. J. Powder Met., 6 (1970) 65. 3 J.B. Moore and R.L. Athey, 18th Sagamore Army Materials Conference, Syracuse Univ. Press, 1972, p. 281. 4 L.N. Moskowitz, R.M. Pelloux and N.J. Grant, Proc. 2nd Int. Conf. on Superalloys, Champion, Penn., 1972, Z-1. 5 Y. Kobayashi, M.S. Thesis, Dept. of Matl. Sc. and Eng., Mass. Inst. of Tech., 1974. 6 A.R. Kaufman, U.S. Patent No. 3,099,041, (1963). 7 D.L. Holt and W.A. Baekofen, Trans. ASM, 59 (1966) 755. 8 G. Rai and N.J. Grant, Met. Trans. A, 6A (1975) 385. 9 A.E. Geckinli and G.R. Barrett, Scripta Met., 8 (1974) 115. 10 D.A. Woodford, Trans Quart. ASM, 62 (1969) 291. 11 R. Swalin and A. Martin, Trans. AIME, 206 (1956) 567.