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Enhanced superplasticity in a Ti-6Al-4V alloy processed by severe plastic deformation
Scripta mater. 43 (2000) 819 – 824 www.elsevier.com/locate/scriptamat
ENHANCED SUPERPLASTICITY IN A Ti-6Al-4V ALLOY PROCESSED BY SEVERE PLASTIC DEFORMATION A.V. Sergueevaa,b, V.V. Stolyarova, R.Z. Valieva and A.K. Mukherjeeb a
University of California, Chemical Engineering and Material Science Department, Davis, CA 95616, USA b Ufa State Aviation Technical University, Institute of Physics of Advanced Materials, 12 Marks str., Ufa 450000, Russia (Received April 19, 2000) (Accepted in revised form May 29, 2000) Keywords: Ultrafine-grained titanium alloy; Severe plastic deformation; High strain rate superplasticity Introduction Titanium alloys are well known for their potential as superplastic materials for applications in different areas including the aerospace, chemical, transportation, naval and biomedical fields. Generally, superplastic flow in these materials is achieved at intermediate temperatures (usually higher than 800°C) and at low strain rates (10⫺4 . . . 10⫺3 s⫺1) [1–3]. Decrease in superplastic deformation temperature and/or increase in strain rate are important from a technological viewpoint. Use of faster strain rates allows significant decrease in the time of forming, which is of great importance for industrial application. It enlarges the opportunity of using superplastic deformation for mass processing of products. Recent experiments on superplastic materials have demonstrated that a reduction in grain size has the potential of both decreasing the temperature and increasing the strain rate associated with optimum superplastic flow. In particular, it can be achieved through grain refinement by severe plastic deformation (SPD), i.e. high pressure torsion straining, equal channel angular pressing and multiple forging [4,5]. For example, it has been revealed that titanium alloys can exhibit superplasticity at lower temperatures when they possess ultrafine-grained (UFG) microstructure [5]. At the same time, several recent publications report the occurrence of high strain rate superplasticity (HSRS) where the optimum strain rate is higher than 10⫺2 s⫺1 in UFG aluminum alloys after SPD-processing [6 – 8]. In present work, we report the first results on HSRS in a Ti-based alloy processed by high pressure torsion (HPT) straining. This method of severe plastic deformation (SPD) allows the production of samples with no residual porosity and with grain sizes as small as 100 nm or less [9]. This paper deals with a grain refinement and enhanced superplasticity in the well-known commercial Ti-6Al-4V alloy. Experimental The specimens of 15 mm in diameter and 0.5 mm in thickness were cut out of Ti-6Al-4V hot rolled rods, 20 mm in diameter, with a mean grain size of 10 m. In initial state this alloy had about 5% -phase. These specimens were subjected to HPT with true logarithmic strain of about 7 at room temperature. Details of the HPT processing were reported earlier [4,9]. Two sets of Ti-6Al-4V alloy 1359-6462/00/$–see front matter. © 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S1359-6462(00)00496-6
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Figure 1. TEM micrographs (with SAED as an insert) of (a) as-processed SPD Ti-6Al-4V alloy after HPT by route 1 and (b) after HPT by route 2 and heating up to 725°C.
specimens were studied, hereafter to be referred to as processed by HPT under applied pressure of 2.2 GPa (route 1) and by HPT under applied pressure of 5 GPa (route 2). The microstructures of the processed specimens were investigated by transmission electron microscopy (TEM). Thin foils for TEM were prepared by an electrolitic thinning using a twin jet polisher with an electrolyte solution of 6% perchloric acid, 35% butanol and 59% methanol at ⫺25°C. The tensile tests were performed using a custom-built computer controlled constant strain rate tensile test machine with a displacement resolution of 5 m and a load resolution of 0.1 N. Tensile specimens with a 1-mm gage length ⫻ 1 mm width were electro discharge machined from disks. In order to estimate the strain rate sensitivity (m-value) the strain rate change tests were carried out by instantaneously changing the cross-head speed by factor of 2. Results TEM micrographs and selected area electron diffraction patterns (SAED), obtained from region with a diameter of 2 m, of as-processed alloy by routes 1 is shown in Fig. 1a. It is seen that in the as-deformed state the microstructure of the Ti-6Al-4V alloy is highly strained with complicated non-uniform contrast because of the presence high density of different type of lattice defects. The dark field images indicate that the size of separate fragments of structure is within 100 –200 nm. The diffraction pattern (inset) with numerous spots arranged along circles indicates the presence of crystallites separated by high angle boundaries. The appearance of significant streaking of diffraction spots indicates the presence of high internal stresses and elastic distortions of crystal lattice. Such microstructure was observed in samples processed by both routes but the streaking of spots and their number on SAED patterns were slightly higher for the alloy after HPT under the higher pressure (route 2), suggesting higher density of lattice defects. Observed diffraction patterns have corresponded only to ␣-phase and did not reveal any presence of -phase, indicating its dissolution during SPD-processing. The heating of deformed specimens up to 650°C does not lead to changes of microstructure. However there is evidence of appearance of some areas (grains) that are free of dislocations with size of about 200 –300 nm after heating to 725°C. But the fraction of such areas is small and it does not exceed 15 and 25% for the alloy processed by either route 1 and route 2. The microstructure of specimen processed by route 2 and after heating to 725°C is shown on Fig. 1b. Tensile straining at the elevated temperatures (650 and 725°C) leads to further structural changes and to the formation of UFG structures with grain size of about 400 –500 nm, as shown below.
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Figure 2. Typical stress-strain curves for the alloy after HPT by route 1 at 650°C and different strain rates.
The tensile testing of the as-processed Ti alloy showed that a view of true stress-strain curves depends significantly on the strain rate. Typical true stress-strain curves at 650°C and different strain rates for the alloy processed by route 1 are shown on Fig. 2. It is seen that the flow curve at strain rate of 10⫺4 s⫺1 has a shape typical for superplastic deformation but at higher strain rates the alloy shows an extensive hardening and lower tensile elongations. The increasing of the pressure during HPT (route 2) leads to significant decrease in the flow stress and an increasing of elongation (Fig. 3). The results obtained on the mechanical behavior of the HPT-processed Ti-6Al-4V alloy at 650° and 725°C are presented in the Table 1. The analysis of tensile tests data (Table 1) shows that the alloy exhibits superplasticity with tensile elongations more than 500% at 650°C and at strain rates 10⫺4 s⫺1(route 1) and 10⫺3 s⫺1 (route 2). Moreover, high elongations (more than 200%) were obtained at 650 and 725°C and at a strain rate 10⫺2 s⫺1 that is two orders of magnitude higher than the strain rate for superplastic deformation of this alloy in the microcrystalline state [2,3]. This alloy processed by route 2 has shown an elongation more than 200% even at a strain rate 10⫺1 s⫺1 at 725°C. A view of several specimens processed by HPT before and after deformation is shown on Fig. 4. The m-values were obtained from the slope of the log-log⑀ curves and from the strain rate change tests. It was determined to be within 0.38 – 0.46 at 650° and 725°C in the strain rate range 10⫺4-10⫺3 s⫺1 for the alloy processed by route 1 and the m-values decrease down to 0.18 – 0.28 at higher rates. But this value is approximately the same (0.32– 0.39) in a wide strain rate range (10⫺3-10⫺1 s⫺1) at the same temperatures for the alloy processed by route 2. The m-values obtained from jump-tests are also given in the Table 1.
Figure 3. Stress-strain curves for the alloy after HPT by different routes at 650°C and strain rate of 10⫺3 s⫺1.
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TABLE 1 Superplastic Characteristics of Ti-6Al-4V after HPT State e 1 2 3 4 5 6 7 8 9
Process
Temperature, T, °C
Route 1
650
725 Route 2
650 725
Strain Rate ⑀, s⫺1
Elong. %
Flow Stress 0.1, MPa
10⫺2 10⫺3 10⫺4 10⫺1 10⫺2 10⫺2 10⫺3 10⫺1 10⫺2
262 336 530 130 400 305 575 215 364
300 190 65 300 155 185 85 335 135
Strain Rate Sensitivity, m 0.19 0.38 0.18 0.37 0.32
Special TEM investigations were carried out on the samples after superplastic deformation. The thin foils were prepared from the gage of deformed specimens. The typical microstructure of the HPTprocessed alloy (route 1) after tensile deformation of 100% at 725°C and at a strain rate 10⫺2 s⫺1 is shown on Fig. 5. It is seen that UFG structure with mean grain size of approximately 400 nm was formed during the tensile straining. Similar structure was observed after tensile deformation in samples produced by HPT under the higher pressure (route 2) but the after deformation grain size was smaller (mean grain size of about 300 nm). Despite the high elongation of samples the grains remained an equaxied that is typical for microstructure after superplastic deformation [10,11]. Dislocations within grains and in grain boundaries were observed. It is necessary to note that the appearance of particles of -phase with size of 100 –200 nm were revealed by using of Energy Dispersive Spectrometry (EDS) analysis and they were located mostly on the grain boundaries and in triple points. But the volume of these particles was very small (less than 5%). The details of this investigation will be reported later. Another feature of the alloy after superplastic deformation is the absence of any voids and we could not detect any cavitation after high elongation even before fracture. Discussion Superplasticity of Ti-6Al-4V alloy with microcrystalline structure with grain size of few microns was investigated in detail in [1–3,10]. As a rule, superplasic deformation in such alloy was observed at
Figure 4. A view of samples before and after deformation.
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Figure 5. TEM micrograph (with SAED as an insert) of Ti-6Al-4V alloy after HPT by route 2 and tensile deformation at 725°C and a strain rate of 10⫺2 s⫺1.
800[ellipsis]900°C and a strain rate of about 10⫺4 s⫺1. Typical characteristics of superplasticity in microcrystalline alloy are strain hardening at high temperatures and/or low strain rates and strain softening at lower temperatures and higher strain rates, which were observed in tensile tests [3]. Such difference in mechanical behavior is associated with some grain growth and changes in the ␣/ phase ratio during heating and deformation, e.g., the content of the -phase can be up to 70% at 900°C. Moreover, this alloy with microcrystalline structure develops cavities during superplastic deformation, which grow and coalesce, leading to cavitation damage [12]. The present results show that superplasticity in this alloy can be significantly enhanced by SPDprocessing. After HPT this alloy exhibits superplastic behavior at lower temperatures (650 . . . 725°C) and high strain rates (even at 10⫺1 s⫺1) and this is a result of formation of ultra-fine grains during heating and deformation of the alloy subjected to severe plastic deformation. Decrease in the grain size of the alloy due to higher pressure during HPT (route 2) leads to further decrease in flow stresses, which emphasizes the role of the grain size in the processed UFG microstructure. In both cases, no cavities were revealed after high elongations even before fracture. Probably, a reduction in grain size reduced cavity nucleation by either relaxing stress concentration rapidly or developing conditions not favorable for nucleation of cavities at triple points or at grain boundary ledges. Our investigations revealed very low volume fraction of -phase in the HPT-processed Ti-6Al-4V alloy at testing temperatures and after deformation (less than 5%). Hence the ␣-phase dominated the deformation behavior. The decreased amount of -phase in the as-processed alloy is obviously due to the severe plastic deformation, which has potential to dissolve second phases because of increase in the limit of solubility [9]. There was intensive strain hardening at all testing temperatures and strain rates higher than 10⫺3 s⫺1. Such behavior has been observed already in several alloys subjected to SPD [13,14]. They exhibited enhanced superplastic properties, but they also showed significant strain hardening that is caused by a change in deformation mechanism, probably due to difficulties of dislocation accommodation of grain boundary sliding in small grains [15]. Conclusions 1. High elongations in Ti-6Al-4V alloy subjected to high pressure torsion were observed during tensile tests at relatively low temperatures, 650° and 725°C. 2. The as-processed alloy exhibited high elongations even at high strain rates of 10⫺2 and 10⫺1 s⫺1, indicating the appearance of high strain rate superplasticity.
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3. Enhanced superplasticity in the alloy subjected to high pressure torsion is associated with the formation of ultra-fine grains during heating after SPD-processing. An increase in applied pressure during HPT leads to formation of yet smaller grain sizes. Acknowledgments This investigation was supported by a grant from the US National Science Foundation (NSF-DMR9903321). References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.
C. H. Hamilton and A. K. Ghosh, in Titanium’80: Science and Technology, Proceedings of Fourth International Conference on Titanium, ed. H. Kamura and O. Trumi, p. 1001 (1980). O. A. Kaibyshev, Superplasticity of Alloys, Intermetallics, and Ceramics, p. 317, Springer-Verlag, Berlin (1992). H. S. Yang, G. Gurewitz, and A. K. Mukherjee. Mater. Trans. JIM. 32(5), 465 (1991). R. Z. Valiev, A. V. Korznikov, and R. R. Mulyukov, Mater. Sci. Eng. 168A, 141 (1993). G. A. Salishchev, R. M. Galeyev, S. P. Malysheva, and O. R. Valiakhmetov, Mater. Sci. Forum. 243–245, 585 (1997). R. Z. Valiev, D. A. Salimonenko, N. K. Tsenev, P. B. Berbon, and T. G. Langdon, Scripta Mater. 37(12), 1945 (1997). D. W. Kum, W. J. Kim, and G. Frommeyer, Scripta Mater. 40(2), 223 (1999). P. B. Berbon, M. Furukawa, Z. Horita, M. Nemoto, N. K. Tsenev, R. Z. Valiev, and T. G. Langdon, Phil. Mag. Lett. 78(4), 313 (1998). R. Z. Valiev, R. K. Islamgaliev, and I. V. Alexandrov, Prog. Mater. Sci. 45, 103 (2000). A. K. Mukherjee, Material Science and Technology, vol. 6 —Plastic Deformation and Fraction of Materials, p. 407, VCH Verlagsgessellschaft, Weinheim, Germany (1993). T. G. Nie, J. Wadworth, and O. D. Sherby, Superplasticity in Metals and Ceramics, Cambridge University Press (1997). S. L. Semiatin, et al., Mater. Sci. Eng. A256, 92 (1998). R. S. Mishra, R. Z. Valiev, S. X. McFadden, and A. K. Mukherjee, Mater. Sci. Eng. A252, 174 (1998). R. S. Mishra, R. Z. Valiev, S. X. McFadden, R. K. Islamgaliev, and A. K. Mukherjee, Phil. Mag. to be published. R. S. Mishra and A. K. Mukherjee, in Superplasticity and Superplastic Forming, ed. A. K., Ghosh and T. R. Bieler, p. 109, TMS, Warrendale, PA (1998).
ENHANCED SUPERPLASTICITY IN A Ti-6Al-4V ALLOY PROCESSED BY SEVERE PLASTIC DEFORMATION A.V. Sergueevaa,b, V.V. Stolyarova, R.Z. Valieva and A.K. Mukherjeeb a
University of California, Chemical Engineering and Material Science Department, Davis, CA 95616, USA b Ufa State Aviation Technical University, Institute of Physics of Advanced Materials, 12 Marks str., Ufa 450000, Russia (Received April 19, 2000) (Accepted in revised form May 29, 2000) Keywords: Ultrafine-grained titanium alloy; Severe plastic deformation; High strain rate superplasticity Introduction Titanium alloys are well known for their potential as superplastic materials for applications in different areas including the aerospace, chemical, transportation, naval and biomedical fields. Generally, superplastic flow in these materials is achieved at intermediate temperatures (usually higher than 800°C) and at low strain rates (10⫺4 . . . 10⫺3 s⫺1) [1–3]. Decrease in superplastic deformation temperature and/or increase in strain rate are important from a technological viewpoint. Use of faster strain rates allows significant decrease in the time of forming, which is of great importance for industrial application. It enlarges the opportunity of using superplastic deformation for mass processing of products. Recent experiments on superplastic materials have demonstrated that a reduction in grain size has the potential of both decreasing the temperature and increasing the strain rate associated with optimum superplastic flow. In particular, it can be achieved through grain refinement by severe plastic deformation (SPD), i.e. high pressure torsion straining, equal channel angular pressing and multiple forging [4,5]. For example, it has been revealed that titanium alloys can exhibit superplasticity at lower temperatures when they possess ultrafine-grained (UFG) microstructure [5]. At the same time, several recent publications report the occurrence of high strain rate superplasticity (HSRS) where the optimum strain rate is higher than 10⫺2 s⫺1 in UFG aluminum alloys after SPD-processing [6 – 8]. In present work, we report the first results on HSRS in a Ti-based alloy processed by high pressure torsion (HPT) straining. This method of severe plastic deformation (SPD) allows the production of samples with no residual porosity and with grain sizes as small as 100 nm or less [9]. This paper deals with a grain refinement and enhanced superplasticity in the well-known commercial Ti-6Al-4V alloy. Experimental The specimens of 15 mm in diameter and 0.5 mm in thickness were cut out of Ti-6Al-4V hot rolled rods, 20 mm in diameter, with a mean grain size of 10 m. In initial state this alloy had about 5% -phase. These specimens were subjected to HPT with true logarithmic strain of about 7 at room temperature. Details of the HPT processing were reported earlier [4,9]. Two sets of Ti-6Al-4V alloy 1359-6462/00/$–see front matter. © 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S1359-6462(00)00496-6
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Figure 1. TEM micrographs (with SAED as an insert) of (a) as-processed SPD Ti-6Al-4V alloy after HPT by route 1 and (b) after HPT by route 2 and heating up to 725°C.
specimens were studied, hereafter to be referred to as processed by HPT under applied pressure of 2.2 GPa (route 1) and by HPT under applied pressure of 5 GPa (route 2). The microstructures of the processed specimens were investigated by transmission electron microscopy (TEM). Thin foils for TEM were prepared by an electrolitic thinning using a twin jet polisher with an electrolyte solution of 6% perchloric acid, 35% butanol and 59% methanol at ⫺25°C. The tensile tests were performed using a custom-built computer controlled constant strain rate tensile test machine with a displacement resolution of 5 m and a load resolution of 0.1 N. Tensile specimens with a 1-mm gage length ⫻ 1 mm width were electro discharge machined from disks. In order to estimate the strain rate sensitivity (m-value) the strain rate change tests were carried out by instantaneously changing the cross-head speed by factor of 2. Results TEM micrographs and selected area electron diffraction patterns (SAED), obtained from region with a diameter of 2 m, of as-processed alloy by routes 1 is shown in Fig. 1a. It is seen that in the as-deformed state the microstructure of the Ti-6Al-4V alloy is highly strained with complicated non-uniform contrast because of the presence high density of different type of lattice defects. The dark field images indicate that the size of separate fragments of structure is within 100 –200 nm. The diffraction pattern (inset) with numerous spots arranged along circles indicates the presence of crystallites separated by high angle boundaries. The appearance of significant streaking of diffraction spots indicates the presence of high internal stresses and elastic distortions of crystal lattice. Such microstructure was observed in samples processed by both routes but the streaking of spots and their number on SAED patterns were slightly higher for the alloy after HPT under the higher pressure (route 2), suggesting higher density of lattice defects. Observed diffraction patterns have corresponded only to ␣-phase and did not reveal any presence of -phase, indicating its dissolution during SPD-processing. The heating of deformed specimens up to 650°C does not lead to changes of microstructure. However there is evidence of appearance of some areas (grains) that are free of dislocations with size of about 200 –300 nm after heating to 725°C. But the fraction of such areas is small and it does not exceed 15 and 25% for the alloy processed by either route 1 and route 2. The microstructure of specimen processed by route 2 and after heating to 725°C is shown on Fig. 1b. Tensile straining at the elevated temperatures (650 and 725°C) leads to further structural changes and to the formation of UFG structures with grain size of about 400 –500 nm, as shown below.
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Figure 2. Typical stress-strain curves for the alloy after HPT by route 1 at 650°C and different strain rates.
The tensile testing of the as-processed Ti alloy showed that a view of true stress-strain curves depends significantly on the strain rate. Typical true stress-strain curves at 650°C and different strain rates for the alloy processed by route 1 are shown on Fig. 2. It is seen that the flow curve at strain rate of 10⫺4 s⫺1 has a shape typical for superplastic deformation but at higher strain rates the alloy shows an extensive hardening and lower tensile elongations. The increasing of the pressure during HPT (route 2) leads to significant decrease in the flow stress and an increasing of elongation (Fig. 3). The results obtained on the mechanical behavior of the HPT-processed Ti-6Al-4V alloy at 650° and 725°C are presented in the Table 1. The analysis of tensile tests data (Table 1) shows that the alloy exhibits superplasticity with tensile elongations more than 500% at 650°C and at strain rates 10⫺4 s⫺1(route 1) and 10⫺3 s⫺1 (route 2). Moreover, high elongations (more than 200%) were obtained at 650 and 725°C and at a strain rate 10⫺2 s⫺1 that is two orders of magnitude higher than the strain rate for superplastic deformation of this alloy in the microcrystalline state [2,3]. This alloy processed by route 2 has shown an elongation more than 200% even at a strain rate 10⫺1 s⫺1 at 725°C. A view of several specimens processed by HPT before and after deformation is shown on Fig. 4. The m-values were obtained from the slope of the log-log⑀ curves and from the strain rate change tests. It was determined to be within 0.38 – 0.46 at 650° and 725°C in the strain rate range 10⫺4-10⫺3 s⫺1 for the alloy processed by route 1 and the m-values decrease down to 0.18 – 0.28 at higher rates. But this value is approximately the same (0.32– 0.39) in a wide strain rate range (10⫺3-10⫺1 s⫺1) at the same temperatures for the alloy processed by route 2. The m-values obtained from jump-tests are also given in the Table 1.
Figure 3. Stress-strain curves for the alloy after HPT by different routes at 650°C and strain rate of 10⫺3 s⫺1.
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TABLE 1 Superplastic Characteristics of Ti-6Al-4V after HPT State e 1 2 3 4 5 6 7 8 9
Process
Temperature, T, °C
Route 1
650
725 Route 2
650 725
Strain Rate ⑀, s⫺1
Elong. %
Flow Stress 0.1, MPa
10⫺2 10⫺3 10⫺4 10⫺1 10⫺2 10⫺2 10⫺3 10⫺1 10⫺2
262 336 530 130 400 305 575 215 364
300 190 65 300 155 185 85 335 135
Strain Rate Sensitivity, m 0.19 0.38 0.18 0.37 0.32
Special TEM investigations were carried out on the samples after superplastic deformation. The thin foils were prepared from the gage of deformed specimens. The typical microstructure of the HPTprocessed alloy (route 1) after tensile deformation of 100% at 725°C and at a strain rate 10⫺2 s⫺1 is shown on Fig. 5. It is seen that UFG structure with mean grain size of approximately 400 nm was formed during the tensile straining. Similar structure was observed after tensile deformation in samples produced by HPT under the higher pressure (route 2) but the after deformation grain size was smaller (mean grain size of about 300 nm). Despite the high elongation of samples the grains remained an equaxied that is typical for microstructure after superplastic deformation [10,11]. Dislocations within grains and in grain boundaries were observed. It is necessary to note that the appearance of particles of -phase with size of 100 –200 nm were revealed by using of Energy Dispersive Spectrometry (EDS) analysis and they were located mostly on the grain boundaries and in triple points. But the volume of these particles was very small (less than 5%). The details of this investigation will be reported later. Another feature of the alloy after superplastic deformation is the absence of any voids and we could not detect any cavitation after high elongation even before fracture. Discussion Superplasticity of Ti-6Al-4V alloy with microcrystalline structure with grain size of few microns was investigated in detail in [1–3,10]. As a rule, superplasic deformation in such alloy was observed at
Figure 4. A view of samples before and after deformation.
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Figure 5. TEM micrograph (with SAED as an insert) of Ti-6Al-4V alloy after HPT by route 2 and tensile deformation at 725°C and a strain rate of 10⫺2 s⫺1.
800[ellipsis]900°C and a strain rate of about 10⫺4 s⫺1. Typical characteristics of superplasticity in microcrystalline alloy are strain hardening at high temperatures and/or low strain rates and strain softening at lower temperatures and higher strain rates, which were observed in tensile tests [3]. Such difference in mechanical behavior is associated with some grain growth and changes in the ␣/ phase ratio during heating and deformation, e.g., the content of the -phase can be up to 70% at 900°C. Moreover, this alloy with microcrystalline structure develops cavities during superplastic deformation, which grow and coalesce, leading to cavitation damage [12]. The present results show that superplasticity in this alloy can be significantly enhanced by SPDprocessing. After HPT this alloy exhibits superplastic behavior at lower temperatures (650 . . . 725°C) and high strain rates (even at 10⫺1 s⫺1) and this is a result of formation of ultra-fine grains during heating and deformation of the alloy subjected to severe plastic deformation. Decrease in the grain size of the alloy due to higher pressure during HPT (route 2) leads to further decrease in flow stresses, which emphasizes the role of the grain size in the processed UFG microstructure. In both cases, no cavities were revealed after high elongations even before fracture. Probably, a reduction in grain size reduced cavity nucleation by either relaxing stress concentration rapidly or developing conditions not favorable for nucleation of cavities at triple points or at grain boundary ledges. Our investigations revealed very low volume fraction of -phase in the HPT-processed Ti-6Al-4V alloy at testing temperatures and after deformation (less than 5%). Hence the ␣-phase dominated the deformation behavior. The decreased amount of -phase in the as-processed alloy is obviously due to the severe plastic deformation, which has potential to dissolve second phases because of increase in the limit of solubility [9]. There was intensive strain hardening at all testing temperatures and strain rates higher than 10⫺3 s⫺1. Such behavior has been observed already in several alloys subjected to SPD [13,14]. They exhibited enhanced superplastic properties, but they also showed significant strain hardening that is caused by a change in deformation mechanism, probably due to difficulties of dislocation accommodation of grain boundary sliding in small grains [15]. Conclusions 1. High elongations in Ti-6Al-4V alloy subjected to high pressure torsion were observed during tensile tests at relatively low temperatures, 650° and 725°C. 2. The as-processed alloy exhibited high elongations even at high strain rates of 10⫺2 and 10⫺1 s⫺1, indicating the appearance of high strain rate superplasticity.
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3. Enhanced superplasticity in the alloy subjected to high pressure torsion is associated with the formation of ultra-fine grains during heating after SPD-processing. An increase in applied pressure during HPT leads to formation of yet smaller grain sizes. Acknowledgments This investigation was supported by a grant from the US National Science Foundation (NSF-DMR9903321). References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.
C. H. Hamilton and A. K. Ghosh, in Titanium’80: Science and Technology, Proceedings of Fourth International Conference on Titanium, ed. H. Kamura and O. Trumi, p. 1001 (1980). O. A. Kaibyshev, Superplasticity of Alloys, Intermetallics, and Ceramics, p. 317, Springer-Verlag, Berlin (1992). H. S. Yang, G. Gurewitz, and A. K. Mukherjee. Mater. Trans. JIM. 32(5), 465 (1991). R. Z. Valiev, A. V. Korznikov, and R. R. Mulyukov, Mater. Sci. Eng. 168A, 141 (1993). G. A. Salishchev, R. M. Galeyev, S. P. Malysheva, and O. R. Valiakhmetov, Mater. Sci. Forum. 243–245, 585 (1997). R. Z. Valiev, D. A. Salimonenko, N. K. Tsenev, P. B. Berbon, and T. G. Langdon, Scripta Mater. 37(12), 1945 (1997). D. W. Kum, W. J. Kim, and G. Frommeyer, Scripta Mater. 40(2), 223 (1999). P. B. Berbon, M. Furukawa, Z. Horita, M. Nemoto, N. K. Tsenev, R. Z. Valiev, and T. G. Langdon, Phil. Mag. Lett. 78(4), 313 (1998). R. Z. Valiev, R. K. Islamgaliev, and I. V. Alexandrov, Prog. Mater. Sci. 45, 103 (2000). A. K. Mukherjee, Material Science and Technology, vol. 6 —Plastic Deformation and Fraction of Materials, p. 407, VCH Verlagsgessellschaft, Weinheim, Germany (1993). T. G. Nie, J. Wadworth, and O. D. Sherby, Superplasticity in Metals and Ceramics, Cambridge University Press (1997). S. L. Semiatin, et al., Mater. Sci. Eng. A256, 92 (1998). R. S. Mishra, R. Z. Valiev, S. X. McFadden, and A. K. Mukherjee, Mater. Sci. Eng. A252, 174 (1998). R. S. Mishra, R. Z. Valiev, S. X. McFadden, R. K. Islamgaliev, and A. K. Mukherjee, Phil. Mag. to be published. R. S. Mishra and A. K. Mukherjee, in Superplasticity and Superplastic Forming, ed. A. K., Ghosh and T. R. Bieler, p. 109, TMS, Warrendale, PA (1998).