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Characterization of low-temperature superplasticity in a thermomechanically processed TiAl based alloy
Materials Science and Engineering A329– 331 (2002) 885– 890 www.elsevier.com/locate/msea
Characterization of low-temperature superplasticity in a thermomechanically processed TiAl based alloy J. Sun *, Y.H. He, J.S. Wu Key Laboratory for High Temperature Materials and Tests of Ministry of Education, School of Materials Science and Engineering, Shanghai Jiao-tong Uni6ersity, Shanghai 200030, People’s Republic of China
Abstract Superplasticity of a TiAl based alloy at low temperature ranging from 750 to 900 °C and at strain rates from 2 ×10 − 5 to 2 ×10 − 4 s − 1 is characterized in this work. In order to refine the grains, two-step forging techniques were applied for materials processing without subsequent annealing after forging. The tensile elongations between 150 and 533% were obtained. An extensive strain hardening on the true stress-strain curves is linked to the high dense mobile dislocations during deformation. The activation energy of 220 kJ mol − 1 was measured which is close to the activation energy of dislocation pipe diffusion. Evolution of the microstructures after superplastic deformation were also performed by optical microscope and transmission electron microscope to correlate the mechanical properties. Based on these studies, it is suggested that the predominant mechanism for low temperature superplasticity is grain boundary sliding at low strain and dislocation glide creep at high strain rates, respectively. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Titanium aluminides; Superplastic behaviors; Thermomechanical treatment
1. Introduction Superplasticity of the ordered intermetallic alloys has received much attention during the past several years because superplastic forming is an attractive technology to produce the near-net-shape structural components for these alloys, which normally exhibit low ductility and toughness at intermediate and room temperatures [1,2]. Generally, the melting points of ordered intermetallic alloys are high and most of them keep the long-range-ordered structure up to the melting point due to their strong covalent inter-atomic bonding. Therefore, these alloys usually exhibit superplastic properties at relatively high temperatures. Titanium aluminide, one of the most important ordered intermetallic alloys, has been demonstrated superplasticity at high temperatures and the high temperature superplastic behaviors and mechanisms of TiAl based alloys have also been well investigated [3 – 8]. However, from the technological viewpoint, there is an obvious desire to reduce the deformation temperature at which the * Corresponding author. E-mail address: [email protected] (J. Sun).
alloys are still highly superplastic. Recently, Nieh et al. found superplasticity of a powder extruded TiAl based alloy with a metastable b phase at temperature as low as 800 °C, that is close to the ductile-to-brittle transition temperature of TiAl [9,10]. Low-temperature superplasticity was also achieved at 750 850 °C in submicro-crystalline TiAl alloy produced by multiple forging [11]. The activation energies for low-temperature superplasticity were reported to be 194 –200 kJ mol − 1, which are much lower than those for high temperature superplasticity. However, different mechanisms have been proposed to explain the superplastic deformation of those alloys by these authors, respectively [10,11]. The purpose of the present work is to investigate superplastic behaviors of a TiAl based alloy at low temperature ranging from 750 to 900 °C and at strain rates from 2× 10 − 5 to 2×10 − 4 s − 1. In order to refine the grain of the alloy, two-step forging techniques were applied for materials processing without subsequent annealing treatment after forging. Microstructural evolution after superplastic deformation is performed by optical microscope (OM) and transmission electron microscope (TEM) to correlate the mechanical properties.
0921-5093/02/$ - see front matter © 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 6 4 6 - X
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Based on these studies, possible low-temperature superplastic mechanisms of the TiAl alloys are finally discussed.
2. Experimental The alloy with a composition of Ti– 46.8Al – 2.2Cr– 0.2Mo (at.%) was prepared by arc-melting with a consumable electrode in an argon atmosphere. The ingot of 7 kg was given a hot iso-static pressing at 1250 °C for 4 h under a pressure of 170 MPa to remove porosity and homogenized at 1040 °C for 48 h in a flowing argon atmosphere. The ingot was firstly forged iso-thermally at temperature of 1040 °C and at strain rate of about 10 − 4 s − 1, and then subjected to the final forging with a compression strain of 60% at high strain rate of about 10 − 1 s − 1. Specimens for mechanical tests were wire-cut from the treated materials with a gauge section of 2×3× 6 mm. Tensile tests were conducted in air on a Shimaduz test machine equipped with a split and
three zone high temperature furnace. Test temperatures ranged from 750 to 900 °C and the approximate constant strain rates ranging from 2× 10 − 5 to 2×10 − 4 s − 1 were obtained through increasing of the cross-head speed of the machine. The gauge of specimens was coated with a thin special glass layer in order to prevent the specimen from early damage or failure due to oxidation during deformation at high temperature in air. Incremental strain rate tests were performed to determine the flow stress and the strain rate sensitivity value as a function of strain rates, during which the strain rates was increased until a reasonably steady flow stress was attained at each strain rate. Specimens were initially deformed at a strain rate of 2×10 − 4 s − 1 for a strain about 0.2 prior to the step strain rate tests. In such a case, the flow stress at each strain rate was obtained. Also, the dependence of the strain rate sensitivity values on the true strain during superplastic deformation was examined through the jump strain rate tests, during which a periodic increase in strain rate was imposed on a base strain rate. The initial microstructure and microstructure after superplastic deformation of the specimens were observed by optical microscope and transmission electron microscope (JEOL-200CX TEM). The Korll’s agent (10% HF+ 5% HNO3 + 85% H2O) was used to etch the specimens. TEM samples were first thinned mechanically, then finally prepared by twin-jet electro-polishing in a solution of 60 vol.% methanol, 35 vol.% butyl alcohol and 5 vol.% perchloric acid under 15 V and at −30 °C.
3. Results
3.1. Mechanical properties
Fig. 1. The initial microstructure (a); and sub-structure (b) of the thermo-mechanically processed TiAl based alloy examined by OM and TEM.
Fig. 1 shows the initial microstructure of the Ti– 46.8Al–2.2Cr–0.2Mo (at.%) alloy after thermomechanically processing. It is a relatively non-uniform microstructure consisting of partially recrystallized fine grain areas and small portion of coarse g-TiAl grain areas, where the minor a2-Ti3Al phase was present primarily as fine particles and occasionally as deformed lamellar plates. TEM analyses showed that g phase mean grain size was about 0.8 mm in recrystallized areas. High dense dislocations and sub-boundaries were observed within the g grains, especially in the coarse grains. No attempts were made to measure the volume fraction of the a2 phase of the TiAl based alloys in the present work. The tensile true stress-strain curves of the TiAl alloy loaded at different temperatures and strain rates are shown in Fig. 2. It is seen that strain hardening to large strains then followed by strain softening occurred on the true stress-strain curves. The strain corresponding
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Fig. 2. The true stress-strain curves of the TiAl based alloys loaded at different testing conditions.
Fig. 3. Plots of the flow stress versus strain rates of the TiAl alloy determined from the incremental strain rate tests.
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until a reasonably steady flow stress was attained at each strain rate. The plots of flow stresses of the TiAl alloys versus strain rates at different temperatures ranging from 800 to 900 °C are shown in Fig. 3. The strain rate sensitive values, m, were calculated from the slopes of curves. At low strain rate region, a strain rate sensitivity of 0.58 is found, which is close to 0.5, a typical value for many fine-grained superplastic alloys. At high strain rates, the m value of 0.36 was observed, which suggested the change of the deformation mechanism at higher strain rates. It is noted that transitions of strain rate sensitivity were influenced by the testing temperature. At high temperatures, the transition in m from a high value to a low level was shifted to high strain rates. In order to examine the dependence of the strain rate sensitivity on the true strain during superplastic deformation, jump strain rate test was conducted. The typical variation of the strain rate sensitivity, m, as a function of true strain at temperature of 800 °C and at strain rates of 2× 10 − 4 and 2× 10 − 5 s − 1 are shown in Fig. 4. The m value kept nearly constant at the level of about 0.33 at 2×10 − 4 s − 1, although the stress-strain curve exhibited an extended strain hardening stage followed by strain softening. At strain rate of 2× 10 − 5 s − 1, the m value was around 0.5 and decreased slightly with an increase of the true strain. To calculate the activation energy for the superplastic deformation, the strain rates corresponding to various stress levels, such as 60, 90 and 120 MPa, at different temperatures of 800900 °C were determined from Fig. 4. Using the usual power law creep equation, the average activation energy of 220 kJ mol − 1 was measured from the plot of the strain rates versus inverse of testing temperatures. This value is much lower than the activation energy for superplastic deformation at high temperatures (1000 °C and above) of the TiAl based alloys [3–8].
3.2. Microstructural e6olution Fig. 4. The dependence of the strain rate sensitivity on the true strain of the TiAl alloy deformed at 800 °C and at a strain rate of 2 ×10 − 4 and of 2 ×10 − 5 s − 1, respectively.
to the maximum stress increased with an increase of temperature or a decrease of strain rate. The superplastic elongation reached 150% at temperatures of 750 °C and at a strain rate of 2× 10 − 4 s − 1, and increased markedly to 425% at 900 °C and at 2×10 − 4 s − 1; Meanwhile, tensile elongation increased from 234 to 533% with decreasing strain rate from 2×10 − 4 to 2× 10 − 5 s − 1 at 800 °C. Incremental strain rate tests were also performed to determine the flow stress and the strain rate sensitivity of the alloy. Usually, the strain interval was about 0.07
The microstructure of the TiAl based alloys after superplastic deformation was analyzed to correlate the mechanical properties. Fig. 5 shows the OM microstructures of these alloys deformed at different testing conditions. The results of TEM analyses were shown in Figs. 6 and 7. It can be seen that the initial microstructure became uniform after superplastic deformation and the coarse g grains were refined (Figs 1a and 5). TEM observations reveal that the grain size reached 1.24 and 2 mm, respectively, when the sample tested at 800 °C, 2 × 10 − 4 and 2 × 10 − 5 s − 1. At 900 °C and 2×10 − 4 s − 1, the grain size of deformed alloy was about 4 mm. High dense wavy dislocations in grains were observed in the alloy deformed at 800 °C, 2× 10 − 4 s − 1 (Fig. 6b). Generally dislocation density in
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larger grains were higher than that in smaller grains. Small dislocation free grains surrounded by other grains being full of dislocations were also observed (Fig. 6c). With increasing of temperature or decreasing of strain rate, the density of dislocations in grains decreased obviously and dislocations emission from grain boundaries or the grain triple point were often
Fig. 5. Microstructural evolution of the TiAl based alloy after superplastic deformation at (a) 800 °C, 2× 10 − 4 s − 1; (b) 800 °C, 2 × 10 − 5 s − 1; and at (c) 900 °C, 2×10 − 4 s − 1.
observed (Fig. 7a, b), which suggested grain boundaries are the sources of dislocations.
4. Discussion The thermomechanically processed TiAl based alloy with grain size below 1 mm exhibits superplasticity at low-temperatures ranging from 750 to 900 °C. The tensile elongations are between 150 533%. The tensile results also show the following features: (1) an extended hardening stage on the true stress-strain curves; and (2) the low activation energy of 220 kJ mol − 1 for superplastic deformation. An extended hardening stage was observed in submicrocrystalline TiAl alloys during the superplastic deformation at low temperatures [11]. Also, this phenomenon was found in the tensile superplasticity at temperature as low as 650 °C in a nanocrystalline nickel aluminide prepared by severe plastic deformation [12,13]. In this work, TEM analyses show that the slightly grain coarsening occurred during deformation of the TiAl alloy. However, the slightly grain coarsening during deformation alone could not account for the extensive strain hardening on the true stress-strain curves. A possible interpretation for the extensive strain hardening could be related to the role of textures and dislocation activities generated during the deformation. As shown in Fig. 6b, high dense dislocations existed in the grains even after deformation at temperature of 800 °C. These high dense mobile dislocations might be responsible for the extensive strain hardening of the true stress-strain curves when the testing temperatures are below 850 °C. With increasing of temperature or decreasing of strain rate, the rate of strain hardening decreases due to low dense dislocations in grains in the process of deformation of the alloy, during which grain boundaries are considered to be the sources of dislocations. As revealed by TEM in Fig. 6c, small grains free of dislocations presented adjacent to other grains with high dense dislocations of the alloy deformed at 800 °C. This implies that dynamic re-crystallization might occur during the superplastic deformation, which results in the strain softening stage following the strain hardening on the true stress-strain curves at a certain true strain. However, detailed TEM studies are required to examine the microstructure and dislocations in the alloy deformed with different superplastic strains, especially with the strain at which the flow stress reached the maximum value. The value of activation energy of 220 kJ mol − 1 determined is close to 200 and 194 kJ mol − 1 for the low-temperature superplasticity obtained by Nieh et al. [9] and Imayev et al. [11], respectively. However, these value were much lower than the activation energy of 345–425 kJ mol − 1 for high temperature superplastic
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Fig. 6. TEM micrographs showing (a) the microstructure; (b) the typical dislocation morphology; and (c) the small dislocation free grain adjacent to other grains of the TiAl based alloy after superplastic deformation at 800 °C and at 2 ×10 − 4 s − 1.
deformation ( 1000 °C and above) of the TiAl based alloys with a typical duplex microstructure [3– 7]. It is well recognized that for fine-grained alloys, grain boundary sliding is the predominant deformation mechanism, that is governed by either lattice or grain boundary diffusion. The strain rate sensitivity is usually around a typical value of 0.5 [1]. Based on the analyses of activation energy for vacancy formation (154 kJ mol − 1), and for self-diffusion of g-TiAl alloys (291 kJ mol − 1) [14], Imayev et al. proposed that the lower activation energy of 194 kJ mol − 1 corresponds to the activation energy of grain boundary diffusion [11]. Thus, the dominant mechanism for low-temperature superplastic deformation is attributed to the grain boundary sliding controlled by grain boundary diffusion in TiAl alloys with submicrocrystallines. On the other hand, Nieh et al. considered the low activation energy of 210 kJ mol − 1 as the self-diffusion activation energy of b phase in a powder extruded TiAl-based alloy with a metastable b phase, which promotes grain boundary sliding during superplastic deformation [10]. However, those mechanisms could not explain the present experimental results, in which the deformation mechanism changes at different strain rates. As shown in Figs. 3 and 4, the value of m is close to 0.5 at low strain rates. TEM observations also reveal the presence of dislocations near grain boundaries or the triple junctions. In this case, superplastic deformation can be attributed to the grain boundary sliding [1]. However, when the alloy deformed at high strain rates, the m value drops to about 0.33, i.e. the stress exponent, n, is approximately 3. Such the stress exponent is often
noted in the power-law dislocation creep theory [15]. Pu et al. investigated the low-temperature superplasticity of an aluminum alloy subjected to thermomechanical processing. They found that the low-temperature superplasticity could be described by the power-law dislocation creep equation and the lower activation energy of 92 kJ mol − 1 is explained as the dislocation pipe diffusion [16]. As described above, a large amount of dislocations existed in the starting microstructure of the TiAl alloys after forging with a large compression strain and at high strain rate. In such cases the lower activation energy of 220 kJ mol − 1 for low-temperature superplastic deformation of TiAl alloy might be not only the possibility of grain boundary diffusion, but also the possibility of dislocation pipe diffusion. Therefore, the dislocation glide creep can be considered to control the superplastic deformation at high strain rates. This is also supported by TEM observations in which the high dense mobile dislocations were found in grains of the deformed alloy. It is note that since the present analyses of the activation energy were performed using the data from the incremental tests, where the TiAl alloys deformed with low strains, it might be possible that deformation mechanisms would be altered when the alloys deform to large strain. Further studies are required to better understand this question.
5. Conclusions The thermomechanically processed TiAl based alloys exhibits superplasticity at low-temperatures ranging
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ity at low strain rates could be explained by the grain boundary sliding mechanism and dislocation glide creep could account for superplasticity at higher strain rates.
Acknowledgements This work is supported by the National Natural Science Foundation of the People’s Republic of China, Project Number 59671006, and partially by the Science and Technology Commission of the ShangaiMunicipal Government.
References
Fig. 7. The microstructure (a); and dislocation morphology (b) of the TiAl alloy superplastically deformed at 900 °C and at 2 × 10 − 4 s − 1.
from 750 to 900 °C. The tensile elongations were between 150 and 533%. An extensive strain hardening of the true stress-strain curves is linked to the high dense mobile dislocations during deformation. The activation energy of 220 kJ mol − 1 for the low-temperature superplasticity is close to the activation energy for dislocation pipe diffusion. Based on m values, activation energy and the microstructural evolution, superplastic-
[1] A.K. Mukherjee, in: H. Mughrabi (Ed.), Plastic Deformation and Fracture of Materials, vol. 6, Materials Science and Technology, VCH, Weinheim, 1993, p. 407. [2] T.G. Nieh, J. Wadsworth, Inter. Mater. Rev. 44 (1999) 59. [3] R.M. Imavey, O.A. Kaibyshev, G.A. Salishchev, Acta Metall. Mater. 40 (1992) 581. [4] W.B. Lee, H.S. Yang, Y.-W. Kim, A.K. Mukherjee, Scripta Metall. Mater. 29 (1993) 1403. [5] S.C. Cheng, J. Wolfenstine, O.D. Sherby, Metall. Trans. 23A (1992) 1509. [6] J. Sun, J.S. Wu, G.X. Hu, Y.H. He, B.Y. Huang, in: High-Temperature Ordered Intermetallic Alloys, Mater. Res. Soc. Symp. Proc., Vol. 552, 1998, p. kk8.13.1. [7] J. Sun, J.S. Wu, Y.H. He, J. Mater. Sci. 35 (2000) 4919. [8] J. Sun, J.S. Wu, Y.H. He, Z. Metallkd. 91 (2000) 472. [9] T.G. Nieh, J.N. Wang, L.M. Hsiung, J. Wadsworth, V. Sikka, Scripta Mater. 37 (1997) 773. [10] T.G. Nieh, L.M. Hsiung, J. Wadsworth, Intermetallics 7 (1999) 163. [11] V.M. Imavey, G.A. Salishchev, M.R. Shagiev, A.V. Kuznetsov, R.M. Imavey, O.N. Senkov, F.H. Froes, Scripta Mater. 40 (1999) 183. [12] R.S. Mishra, R.Z. Valiev, S.X. McFadden, A.K. Mukherjee, Mater. Sci. Eng. A252 (1998) 174. [13] R.S. Mishra, S.X. McFadden, R.Z. Valiev, A.K. Mukherjee, JOM 51 (1999) 37. [14] G. Sattonnay, C. Dimitrov, O. Dimitrov, Structural Intermetallics, TMS, Warrendale, PA, 1997, p. 205. [15] J. Cadek, Creep in Metallic Materials, Elsevier Science Publishing Co. Inc., Amsterdam, 1988. [16] H.P. Pu, F.C. Liu, J.C. Huang, Metall. Trans. 26A (1995) 1153.
Characterization of low-temperature superplasticity in a thermomechanically processed TiAl based alloy J. Sun *, Y.H. He, J.S. Wu Key Laboratory for High Temperature Materials and Tests of Ministry of Education, School of Materials Science and Engineering, Shanghai Jiao-tong Uni6ersity, Shanghai 200030, People’s Republic of China
Abstract Superplasticity of a TiAl based alloy at low temperature ranging from 750 to 900 °C and at strain rates from 2 ×10 − 5 to 2 ×10 − 4 s − 1 is characterized in this work. In order to refine the grains, two-step forging techniques were applied for materials processing without subsequent annealing after forging. The tensile elongations between 150 and 533% were obtained. An extensive strain hardening on the true stress-strain curves is linked to the high dense mobile dislocations during deformation. The activation energy of 220 kJ mol − 1 was measured which is close to the activation energy of dislocation pipe diffusion. Evolution of the microstructures after superplastic deformation were also performed by optical microscope and transmission electron microscope to correlate the mechanical properties. Based on these studies, it is suggested that the predominant mechanism for low temperature superplasticity is grain boundary sliding at low strain and dislocation glide creep at high strain rates, respectively. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Titanium aluminides; Superplastic behaviors; Thermomechanical treatment
1. Introduction Superplasticity of the ordered intermetallic alloys has received much attention during the past several years because superplastic forming is an attractive technology to produce the near-net-shape structural components for these alloys, which normally exhibit low ductility and toughness at intermediate and room temperatures [1,2]. Generally, the melting points of ordered intermetallic alloys are high and most of them keep the long-range-ordered structure up to the melting point due to their strong covalent inter-atomic bonding. Therefore, these alloys usually exhibit superplastic properties at relatively high temperatures. Titanium aluminide, one of the most important ordered intermetallic alloys, has been demonstrated superplasticity at high temperatures and the high temperature superplastic behaviors and mechanisms of TiAl based alloys have also been well investigated [3 – 8]. However, from the technological viewpoint, there is an obvious desire to reduce the deformation temperature at which the * Corresponding author. E-mail address: [email protected] (J. Sun).
alloys are still highly superplastic. Recently, Nieh et al. found superplasticity of a powder extruded TiAl based alloy with a metastable b phase at temperature as low as 800 °C, that is close to the ductile-to-brittle transition temperature of TiAl [9,10]. Low-temperature superplasticity was also achieved at 750 850 °C in submicro-crystalline TiAl alloy produced by multiple forging [11]. The activation energies for low-temperature superplasticity were reported to be 194 –200 kJ mol − 1, which are much lower than those for high temperature superplasticity. However, different mechanisms have been proposed to explain the superplastic deformation of those alloys by these authors, respectively [10,11]. The purpose of the present work is to investigate superplastic behaviors of a TiAl based alloy at low temperature ranging from 750 to 900 °C and at strain rates from 2× 10 − 5 to 2×10 − 4 s − 1. In order to refine the grain of the alloy, two-step forging techniques were applied for materials processing without subsequent annealing treatment after forging. Microstructural evolution after superplastic deformation is performed by optical microscope (OM) and transmission electron microscope (TEM) to correlate the mechanical properties.
0921-5093/02/$ - see front matter © 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 6 4 6 - X
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Based on these studies, possible low-temperature superplastic mechanisms of the TiAl alloys are finally discussed.
2. Experimental The alloy with a composition of Ti– 46.8Al – 2.2Cr– 0.2Mo (at.%) was prepared by arc-melting with a consumable electrode in an argon atmosphere. The ingot of 7 kg was given a hot iso-static pressing at 1250 °C for 4 h under a pressure of 170 MPa to remove porosity and homogenized at 1040 °C for 48 h in a flowing argon atmosphere. The ingot was firstly forged iso-thermally at temperature of 1040 °C and at strain rate of about 10 − 4 s − 1, and then subjected to the final forging with a compression strain of 60% at high strain rate of about 10 − 1 s − 1. Specimens for mechanical tests were wire-cut from the treated materials with a gauge section of 2×3× 6 mm. Tensile tests were conducted in air on a Shimaduz test machine equipped with a split and
three zone high temperature furnace. Test temperatures ranged from 750 to 900 °C and the approximate constant strain rates ranging from 2× 10 − 5 to 2×10 − 4 s − 1 were obtained through increasing of the cross-head speed of the machine. The gauge of specimens was coated with a thin special glass layer in order to prevent the specimen from early damage or failure due to oxidation during deformation at high temperature in air. Incremental strain rate tests were performed to determine the flow stress and the strain rate sensitivity value as a function of strain rates, during which the strain rates was increased until a reasonably steady flow stress was attained at each strain rate. Specimens were initially deformed at a strain rate of 2×10 − 4 s − 1 for a strain about 0.2 prior to the step strain rate tests. In such a case, the flow stress at each strain rate was obtained. Also, the dependence of the strain rate sensitivity values on the true strain during superplastic deformation was examined through the jump strain rate tests, during which a periodic increase in strain rate was imposed on a base strain rate. The initial microstructure and microstructure after superplastic deformation of the specimens were observed by optical microscope and transmission electron microscope (JEOL-200CX TEM). The Korll’s agent (10% HF+ 5% HNO3 + 85% H2O) was used to etch the specimens. TEM samples were first thinned mechanically, then finally prepared by twin-jet electro-polishing in a solution of 60 vol.% methanol, 35 vol.% butyl alcohol and 5 vol.% perchloric acid under 15 V and at −30 °C.
3. Results
3.1. Mechanical properties
Fig. 1. The initial microstructure (a); and sub-structure (b) of the thermo-mechanically processed TiAl based alloy examined by OM and TEM.
Fig. 1 shows the initial microstructure of the Ti– 46.8Al–2.2Cr–0.2Mo (at.%) alloy after thermomechanically processing. It is a relatively non-uniform microstructure consisting of partially recrystallized fine grain areas and small portion of coarse g-TiAl grain areas, where the minor a2-Ti3Al phase was present primarily as fine particles and occasionally as deformed lamellar plates. TEM analyses showed that g phase mean grain size was about 0.8 mm in recrystallized areas. High dense dislocations and sub-boundaries were observed within the g grains, especially in the coarse grains. No attempts were made to measure the volume fraction of the a2 phase of the TiAl based alloys in the present work. The tensile true stress-strain curves of the TiAl alloy loaded at different temperatures and strain rates are shown in Fig. 2. It is seen that strain hardening to large strains then followed by strain softening occurred on the true stress-strain curves. The strain corresponding
J. Sun et al. / Materials Science and Engineering A329–331 (2002) 885–890
Fig. 2. The true stress-strain curves of the TiAl based alloys loaded at different testing conditions.
Fig. 3. Plots of the flow stress versus strain rates of the TiAl alloy determined from the incremental strain rate tests.
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until a reasonably steady flow stress was attained at each strain rate. The plots of flow stresses of the TiAl alloys versus strain rates at different temperatures ranging from 800 to 900 °C are shown in Fig. 3. The strain rate sensitive values, m, were calculated from the slopes of curves. At low strain rate region, a strain rate sensitivity of 0.58 is found, which is close to 0.5, a typical value for many fine-grained superplastic alloys. At high strain rates, the m value of 0.36 was observed, which suggested the change of the deformation mechanism at higher strain rates. It is noted that transitions of strain rate sensitivity were influenced by the testing temperature. At high temperatures, the transition in m from a high value to a low level was shifted to high strain rates. In order to examine the dependence of the strain rate sensitivity on the true strain during superplastic deformation, jump strain rate test was conducted. The typical variation of the strain rate sensitivity, m, as a function of true strain at temperature of 800 °C and at strain rates of 2× 10 − 4 and 2× 10 − 5 s − 1 are shown in Fig. 4. The m value kept nearly constant at the level of about 0.33 at 2×10 − 4 s − 1, although the stress-strain curve exhibited an extended strain hardening stage followed by strain softening. At strain rate of 2× 10 − 5 s − 1, the m value was around 0.5 and decreased slightly with an increase of the true strain. To calculate the activation energy for the superplastic deformation, the strain rates corresponding to various stress levels, such as 60, 90 and 120 MPa, at different temperatures of 800900 °C were determined from Fig. 4. Using the usual power law creep equation, the average activation energy of 220 kJ mol − 1 was measured from the plot of the strain rates versus inverse of testing temperatures. This value is much lower than the activation energy for superplastic deformation at high temperatures (1000 °C and above) of the TiAl based alloys [3–8].
3.2. Microstructural e6olution Fig. 4. The dependence of the strain rate sensitivity on the true strain of the TiAl alloy deformed at 800 °C and at a strain rate of 2 ×10 − 4 and of 2 ×10 − 5 s − 1, respectively.
to the maximum stress increased with an increase of temperature or a decrease of strain rate. The superplastic elongation reached 150% at temperatures of 750 °C and at a strain rate of 2× 10 − 4 s − 1, and increased markedly to 425% at 900 °C and at 2×10 − 4 s − 1; Meanwhile, tensile elongation increased from 234 to 533% with decreasing strain rate from 2×10 − 4 to 2× 10 − 5 s − 1 at 800 °C. Incremental strain rate tests were also performed to determine the flow stress and the strain rate sensitivity of the alloy. Usually, the strain interval was about 0.07
The microstructure of the TiAl based alloys after superplastic deformation was analyzed to correlate the mechanical properties. Fig. 5 shows the OM microstructures of these alloys deformed at different testing conditions. The results of TEM analyses were shown in Figs. 6 and 7. It can be seen that the initial microstructure became uniform after superplastic deformation and the coarse g grains were refined (Figs 1a and 5). TEM observations reveal that the grain size reached 1.24 and 2 mm, respectively, when the sample tested at 800 °C, 2 × 10 − 4 and 2 × 10 − 5 s − 1. At 900 °C and 2×10 − 4 s − 1, the grain size of deformed alloy was about 4 mm. High dense wavy dislocations in grains were observed in the alloy deformed at 800 °C, 2× 10 − 4 s − 1 (Fig. 6b). Generally dislocation density in
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larger grains were higher than that in smaller grains. Small dislocation free grains surrounded by other grains being full of dislocations were also observed (Fig. 6c). With increasing of temperature or decreasing of strain rate, the density of dislocations in grains decreased obviously and dislocations emission from grain boundaries or the grain triple point were often
Fig. 5. Microstructural evolution of the TiAl based alloy after superplastic deformation at (a) 800 °C, 2× 10 − 4 s − 1; (b) 800 °C, 2 × 10 − 5 s − 1; and at (c) 900 °C, 2×10 − 4 s − 1.
observed (Fig. 7a, b), which suggested grain boundaries are the sources of dislocations.
4. Discussion The thermomechanically processed TiAl based alloy with grain size below 1 mm exhibits superplasticity at low-temperatures ranging from 750 to 900 °C. The tensile elongations are between 150 533%. The tensile results also show the following features: (1) an extended hardening stage on the true stress-strain curves; and (2) the low activation energy of 220 kJ mol − 1 for superplastic deformation. An extended hardening stage was observed in submicrocrystalline TiAl alloys during the superplastic deformation at low temperatures [11]. Also, this phenomenon was found in the tensile superplasticity at temperature as low as 650 °C in a nanocrystalline nickel aluminide prepared by severe plastic deformation [12,13]. In this work, TEM analyses show that the slightly grain coarsening occurred during deformation of the TiAl alloy. However, the slightly grain coarsening during deformation alone could not account for the extensive strain hardening on the true stress-strain curves. A possible interpretation for the extensive strain hardening could be related to the role of textures and dislocation activities generated during the deformation. As shown in Fig. 6b, high dense dislocations existed in the grains even after deformation at temperature of 800 °C. These high dense mobile dislocations might be responsible for the extensive strain hardening of the true stress-strain curves when the testing temperatures are below 850 °C. With increasing of temperature or decreasing of strain rate, the rate of strain hardening decreases due to low dense dislocations in grains in the process of deformation of the alloy, during which grain boundaries are considered to be the sources of dislocations. As revealed by TEM in Fig. 6c, small grains free of dislocations presented adjacent to other grains with high dense dislocations of the alloy deformed at 800 °C. This implies that dynamic re-crystallization might occur during the superplastic deformation, which results in the strain softening stage following the strain hardening on the true stress-strain curves at a certain true strain. However, detailed TEM studies are required to examine the microstructure and dislocations in the alloy deformed with different superplastic strains, especially with the strain at which the flow stress reached the maximum value. The value of activation energy of 220 kJ mol − 1 determined is close to 200 and 194 kJ mol − 1 for the low-temperature superplasticity obtained by Nieh et al. [9] and Imayev et al. [11], respectively. However, these value were much lower than the activation energy of 345–425 kJ mol − 1 for high temperature superplastic
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Fig. 6. TEM micrographs showing (a) the microstructure; (b) the typical dislocation morphology; and (c) the small dislocation free grain adjacent to other grains of the TiAl based alloy after superplastic deformation at 800 °C and at 2 ×10 − 4 s − 1.
deformation ( 1000 °C and above) of the TiAl based alloys with a typical duplex microstructure [3– 7]. It is well recognized that for fine-grained alloys, grain boundary sliding is the predominant deformation mechanism, that is governed by either lattice or grain boundary diffusion. The strain rate sensitivity is usually around a typical value of 0.5 [1]. Based on the analyses of activation energy for vacancy formation (154 kJ mol − 1), and for self-diffusion of g-TiAl alloys (291 kJ mol − 1) [14], Imayev et al. proposed that the lower activation energy of 194 kJ mol − 1 corresponds to the activation energy of grain boundary diffusion [11]. Thus, the dominant mechanism for low-temperature superplastic deformation is attributed to the grain boundary sliding controlled by grain boundary diffusion in TiAl alloys with submicrocrystallines. On the other hand, Nieh et al. considered the low activation energy of 210 kJ mol − 1 as the self-diffusion activation energy of b phase in a powder extruded TiAl-based alloy with a metastable b phase, which promotes grain boundary sliding during superplastic deformation [10]. However, those mechanisms could not explain the present experimental results, in which the deformation mechanism changes at different strain rates. As shown in Figs. 3 and 4, the value of m is close to 0.5 at low strain rates. TEM observations also reveal the presence of dislocations near grain boundaries or the triple junctions. In this case, superplastic deformation can be attributed to the grain boundary sliding [1]. However, when the alloy deformed at high strain rates, the m value drops to about 0.33, i.e. the stress exponent, n, is approximately 3. Such the stress exponent is often
noted in the power-law dislocation creep theory [15]. Pu et al. investigated the low-temperature superplasticity of an aluminum alloy subjected to thermomechanical processing. They found that the low-temperature superplasticity could be described by the power-law dislocation creep equation and the lower activation energy of 92 kJ mol − 1 is explained as the dislocation pipe diffusion [16]. As described above, a large amount of dislocations existed in the starting microstructure of the TiAl alloys after forging with a large compression strain and at high strain rate. In such cases the lower activation energy of 220 kJ mol − 1 for low-temperature superplastic deformation of TiAl alloy might be not only the possibility of grain boundary diffusion, but also the possibility of dislocation pipe diffusion. Therefore, the dislocation glide creep can be considered to control the superplastic deformation at high strain rates. This is also supported by TEM observations in which the high dense mobile dislocations were found in grains of the deformed alloy. It is note that since the present analyses of the activation energy were performed using the data from the incremental tests, where the TiAl alloys deformed with low strains, it might be possible that deformation mechanisms would be altered when the alloys deform to large strain. Further studies are required to better understand this question.
5. Conclusions The thermomechanically processed TiAl based alloys exhibits superplasticity at low-temperatures ranging
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ity at low strain rates could be explained by the grain boundary sliding mechanism and dislocation glide creep could account for superplasticity at higher strain rates.
Acknowledgements This work is supported by the National Natural Science Foundation of the People’s Republic of China, Project Number 59671006, and partially by the Science and Technology Commission of the ShangaiMunicipal Government.
References
Fig. 7. The microstructure (a); and dislocation morphology (b) of the TiAl alloy superplastically deformed at 900 °C and at 2 × 10 − 4 s − 1.
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