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The superelasticity of TiPdNi high temperature shape memory alloy
Intermetallics 11 (2003) 773–778 www.elsevier.com/locate/intermet
The superelasticity of TiPdNi high temperature shape memory alloy Jiansheng Wua, Qingchao Tianb,* a
Key Laboratory of the Ministry of Education for High Temperature Materials and Tests, Shanghai Jiaotong University, Shanghai 200030, PR China b R&D Testing Center, Technology Center, Baoshan Iron & Steel Co. Ltd. Shanghai 201900, PR China Accepted 1 April 2003
Abstract In the high temperature shape memory alloy (HTSMA) Ti51Pd30Ni19, superelasticity is found for the first time and the superelasticity is quite different from that of NiTi or CuZnAl. The recoverable superelastic strain is 7% without failure of the specimen. Shape memory effect (SME) at room temperature is evaluated as 7.2% with the recovery rate of 100%. The phase transformation temperatures were determined using differential scanning calorimeter (DSC). It has been found that the temperatures are changed during the process of alloy preparation. XRD analyses indicate that the macro- residual stress is eliminated and precipitation occurred after recovery treatment. The characteristics of both mechanical and transformation behavior are associated with the changes in the recovery process. Accordingly, an effective way to obtain the superelasticity is pointed out. # 2003 Elsevier Ltd. All rights reserved. Keywords: A. Intermetallics; B. Mechanical properties at high temperatures; B. Phase identification; B. Martensitic transformations; B. Shapememory effects (including superelasticity)
1. Introduction In Ti50PdxNi(50x) alloys, it has been found that the TMT (temperature of martensitic transformation) increases with the increase of Pd content from 20 to 50 at.%. Among these alloys, Ti50Pd30Ni20 has received much more attention because of its high TMT, which is higher than 473 K. This temperature is high enough for engineering application as a high temperature shape memory alloy (HTSMA) [1–7]. It is known that shape memory effect and superelasticity (or pseudoelasticty) are two important characteristics of shape memory alloys. However, the latter property for HTSMA is rarely obtained because the critical stress for slip (CSS) is low [8]. To solve this problem, alloying, aging induced precipitation, and thermomechanical treatment may be effective ways [5,8]. Using the first two methods, superelasticity has not yet been found in TiPd-based HTSMA. However, Golberg et al. [2] found the tendency of superelasticity in Ti50Pd30Ni20 specimens prepared by hot-rolling, cold-rolling, and annealing at 673 K, in sequence. In such a sequence, large numbers of * Corresponding author. E-mail address: [email protected] (Q. Tian). 0966-9795/03/$ - see front matter # 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0966-9795(03)00075-X
dislocations introduced by rolling, after experiencing thermal rearrangement, lead to a high value of CSS. So it is possible that stress induced transformation occurs before slip deformation, resulting in superelasticity. Shimizu et al. [5] improved the shape memory effect of TiNiPd alloys by precipitation hardening. They changed the alloy composition and obtained homogeneously distributed fine precipitates, which increases the CSS. In order to obtain perfect superelastic cycling of HTSMA, the composition of Ti50Pd30Ni20 was adjusted with reference to the methods mentioned above.
2. Experimental procedure Ti51Pd30Ni19 alloy was made by arc melting 99% Ti, 99% Ni and 99.9% Pd on a water-cooled copper mould under a controlled argon atmosphere. The ingot was melted four times for homogenization. Then, it was homogenized in vacuum at 1273 K for 5 h (I). Finally, it was hot-rolled into plate of 1.1 mm thickness at 1073 K, and subsequently cold-rolled into 1 mm thick sheets (II). Tensile specimens were spark cut along the rolling direction, as shown in Fig. 1. The specimens were mechanically polished and then annealed at 673 K in quartz capsule filled with argon for 1 h. After annealing,
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the specimens were quenched into ice water by crushing the capsule (III). In order to find out whether the transformation temperature is changed during the preparation process, specimens for differential scanning calorimetry (DSC) were also prepared. X-ray diffraction (XRD) measurements were performed using ‘D/max IIIA’ automatic diffractometer. Experiments were carried out at room temperature first; then the second measurement was performed as soon as the specimen was slowly heated to 673 K; the diffractometer ran the third scan after the specimens were isothermal for 1 h. With such a specific procedure imitating the annealing process, we can determine whether there were precipitates formed in the recovery process at high austenite temperature. As we know, martensite may exhibit a complicated XRD spectrum of many peaks because of the lower crystal symmetry and the texture of the specimen, therefore, it is difficult to confirm whether the precipitation occurred or not. XRD profile of another specimen in the recovered-state was also measured for comparison in order to find the changes of the residual stress. Tensile tests were carried out using ‘Shimadzu Autograph’ tensile machine at room temperature, 503 and 513 K, respectively. The strain rate was 1104/s. The displacement of the pin, which was inserted into the hole of the specimen, was recorded. The deformation outside the gauge section was ignored in the experiments. At first the specimen was tested at room temperature, in which the alloy is in the martensitic state, and then the specimens were heated to (Af +100) K, thus the martensite transformed to parent phase, in this way, the recovery rate was determined.
tioned previously. The experiments were run under argon atmosphere. Fig. 2 shows the DSC results. In Fig. 2, the numbers 1, 5 and 10 refer to the heating rates of 1, 5 and 10 K/min, respectively. It can be found that the temperatures corresponding to the peak of heat flow curves are the same for the different heating rates (Fig. 2a); the peaks for II correspond well to that for I (Fig. 2b). Fig. 2c represents the transformation temperatures of the alloy in different states during the process of alloy
3. Results and discussion 3.1. Phase transformation behavior Three kinds of specimens for DSC measurements were cut from the alloy in the solution treated state (I), cold-rolled state (II) and recovered state (III) as men-
Fig. 1. The dimensions of the specimen (units, mm).
Fig. 2. The transformation temperatures determined by DSC, the heat flow curves of specimen (a) in solution treated state, (b) under different heat treatment conditions, and (c) the changes of transformation temperatures during the process of alloy preparation, the numbers 1, 5 and 10 in the figure refer to the heating rate.
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preparation. It can be seen that the temperatures for II are similar to that for I, while the temperatures for III are much lower. It may be deduced that the reason why the transformation temperature is changed is that precipitation happens during the annealing process. Because the composition of the matrix may be changed due to precipitation, the transformation temperatures may in turn be affected. The changes may be closely related to the complicated effect of high-density dislocations and the internal stress produced during cold rolling as well. The alloy is basically perfect crystal with few defects after solution treatment at high temperature, and the transformation temperatures measured are close to the intrinsic values. The onset or finish temperatures of phase transformation may vary owing to the different heating rates, but the values corresponding to the peaks remain unchanged. For cold-rolled alloy, on the one hand, high dislocation density provides favorable sites for nucleation during solid-state transformation. On the other hand, the internal compressive stress impedes the phase transformation. The so two effects may maintain equilibrium in the cold-rolled alloy, so the transformation temperatures and the peak temperature do not change a lot. So long as the alloy experiences recovery treatment, the internal stress would be eliminated while the high-density dislocations just adjust appropriately, as indicated by the following X-ray analyses. Thus, the equilibrium was upset, which results in the decrease of the transformation temperatures. 3.2. XRD analyses The structures of TiPdNi alloys were studied by Enami et al. [9] and Lo et al. [10], etc. The structure of the martensite in Ti50PdxNi(50x) (x!15 at.%) alloys was determined as orthorhombic B19 (2H) type while the structure of austenite is of B2 (CsCl type). The XRD spectrum can be indexed with reference to the lattice parameters of parent phase a=0.31 nm (B2), of martensite a=0.278 nm, b=0.489 nm, c=0.459 nm [11], and of the precipitates (Ti2Pd) a=0.309 nm, c=1.0054 nm [12], respectively. Fig. 3 represents the XRD results. The profiles of martensitic phase in both cold-rolled state and recovered state exist only as two clear peaks, which can be identified as (120)B19 and (111)B19 (Fig. 3a). The crystal indices of austenitic phase were identified in Fig. 3b. It can be found from the figure that the specimens exhibit apparent texture, and a strong (110)B2 growth texture, which is similar to the observation of Mathews et al. [13]. Fig. 3a shows that after recovery treatment, the peaks of both (120)B19 and (110)B19 in the cold-rolled alloy move to higher angles, resulting from the increase of spacing between the crystal planes. This means that the macro-residual stress introduced by cold-rolling was
775
compressive, and was eliminated after annealing. It can be also seen that the peaks in recovery state are higher than those in the cold-rolled state, indicating that the lattice distortion in the former is not as serious as that in the latter, or more martensite transformed in the recovered specimen. Fig. 3b shows that after the annealing process was completed, (11-1)Ti2Pd(Ni) peak appears. As we know, low volume fraction of precipitates is difficult to detect by XRD. So it can be deduced from this observation that much precipitation has occurred in the alloy. In fact, precipitates may form during alloy preparation, e.g. during cooling after homogenization, or during hot rolling. These homogeneously distributed precipitates are very fine, they will congregate and grow in the recovery process, and were detected by XRD. It can also be seen that the peaks of the austenitic phase grow a bit, and the reason is similar to that for martensite. 3.3. Tensile deformation Fig. 4a shows stress–strain curves obtained at room temperature. The arrows in the figure indicate recovery strains after heating the specimens to (Af +100) K followed by cooling to room temperature. Three virgin specimens were loaded to 450, 750 and 940 MPa, respectively. The total strain of the first specimen is 7.2% with a recovery rate of 100%, and that for the second specimen is much higher, 11%, and the SME became incomplete (95%). The third specimen was deformed to a total strain of 12.9% without fracture; however, the SME became very poor (55%). The residual strain at room temperature in the three experiments is 3, 6, and 6.8% with increasing applied stress. Fig. 4b represents the experimental results at 503 K. A virgin specimen was repeatedly deformed five times. As shown in Fig. 4b, in the second experiment, the specimen was loaded to the same load as that in the first experiment. However, the width between the loading curve and the unloading curve is narrower, and the stress–strain curve tends to be a cycle. In the followed experiments, the loads were gradually increased to a higher level, the specimen yields at a higher stress correspondingly, and the elastic strength limit increases. The phenomenon is the well-known Bauschinger effect, which is caused by the movement of dislocations. The experiment indicates that Bauschinger effect may promote the CSS to be high enough so that martensite can be induced by stress. 3.4. Superelastic cycling Inspired by the last experiment, we carried out another experiment at a higher temperature, that is, a virgin specimen was tested in the austenitic state. The evolution of superelastic cycling is represented in Fig. 5.
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It can be seen that the stress–strain cycle becomes stable after the specimen was trained five times under stress up to 530 MPa. Then, the cycle stress was elevated to 680 MPa, the cycle becomes stable after cycling proceeds three times. Fig. 5b shows selected cycles of the experiment. It can be clearly seen that the loading curves warp up with the increase of superelastic cycling. A further consideration is shown in Fig. 5c.
The superelastic stress–strain curve may be divided into several stages, as denoted in Fig. 5c, according to the apparently different characteristics of the slope. It is considered that the difference corresponds to different deformation mechanisms. The elastic deformation of austenite proceeds with the initial loading at stage I. Another straight slope is regarded as the elastic deformation of martensite, defined as stage III. The bend connecting the two stages is denoted as stage II, which is
Fig. 3. XRD profiles of Ti51Pd30Ni19, (a) at room temperature and (b) at austenitic state.
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Fig. 4. Stress–strain curves of Ti51Pd30Ni19 tested (a) at room temperature and (b) 503 K.
associated with the stress-induced martensitic transformation and the reorientation of martensitic variants. The unloading curve is observed to be nonlinear, and does not look like its loading part, so it is wholly denoted as stage IV. In this stage, the elastic recovery of austenite and martensite, and the superelastic recovery of the stress-induced transformation are achieved. It can be clearly seen that the superelastic cycle is greatly different to that of TiNi or CuZnAl, where the cycles usually exhibit a platform in stage II. Two aspects may contribute this difference, i.e. the kind of martensite and the precipitation introduced during the recovery treatment. It is known that the martensite in TiPd–Ni (B19) is of higher symmetry and hence has less variants than that in TiNi (B190 ). So it cannot accommodate the applied stress field as effectively as the B190 kind of martensite can. As for the latter reason, the stress introduced martensitic transformation is compressed owing to the stress field produced by precipitation. It can be concluded that large numbers of dislocations and precipitates, which increase the CSS, are two necessary conditions for obtaining superelasticity in TiPdNi HTSMA, while training make the ideal realizable.
Fig. 5. Superelastic cycling of Ti51Pd30Ni19, (a) tested at 513 K, (b) warp of the loading curves and (c) the definition of different deformation stages.
4. Conclusion
1. The phase transformation temperatures of the Ti51Pd30Ni19 alloy are different for the different states in the alloy preparation process. This is attributed to complicated reasons including internal defects and residual stress introduced
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during cold rolling, and the precipitates produced during annealing. 2. Superelasticity was obtained for the first time in TiPdNi alloys. A recoverable superelastic strain of 7% at high temperature, and a total strain of 7.2% with a recovery rate of 100% at room temperature, were obtained. The superelastic cycle of the alloy is different from that in NiTi or CuZnAl, probably due to the relatively higher symmetry of martensite and large amount of precipitates. 3. Superelasticity is obtained by introducing large numbers of dislocations and precipitates in the matrix of TiPdNi alloys and by training as well.
Acknowledgements This work is sponsored by the Science and Technology Commission of the Shanghai Municipal Government, No. 00JC14055.
References [1] Golberg D, Xu Y, Murakami Y, Morito S, Otsuka K, Ueki T. Horikawa J Materials Letters 1995;22:241–8. [2] Golberg D, Xu Y, Murakami Y, Morito S, Otsuka K. Intermetallics 1995;3:35–46. [3] Golberg D, Xu Y, Murakami Y, Morito S, Otsuka K, Ueki T. Horikawa J Script Metallurgica Et Materialia 1994;30(10):1349–54. [4] Cai W, Otsuka K, Asai M. Materials Transactions, JIM 1999; 40(9):895–8. [5] Shimizu S, Xu Y, Okunishi E, Tanaka S, Otsuka K, Mitose K. Materials Letters 1998;34:23. [6] Xu Y, Shimizu S, Suzuki Y, Otsuka K, Ueki T, Mitose K. Acta Mater 1997;45(4):1503–11. [7] Lo YC, Wu SK, Wayman CM. Script Metallurgica Et Materialia 1990;24:1571–6. [8] Otsuka K, Oda K, Ueno Y, Piao M, Ueki T, Horikawa H. Script Metallurgica Et Materialia 1993;29:1355–8. [9] Enami K, Horri K, Takakura J. ISIJ International 1989;29:430. [10] Lo YC, Wu SK. Script Metallurgica Et Materialia 1992;27:1097– 102. [11] Quandt E, Halene C, Holleck H, Feit K, Kohl M, Schlossmacher P, Skokan A, Skrobanek KD. Sensors and Actuators A (Switzerland) 1996;A53:434–9. [12] Schlossmacher P. Materials Letters 1997;31:119–25. [13] Mathews S, Li J, Su Q, Wuttig M. Philosophical Magazine Letters 1999;79(5):265–72.
The superelasticity of TiPdNi high temperature shape memory alloy Jiansheng Wua, Qingchao Tianb,* a
Key Laboratory of the Ministry of Education for High Temperature Materials and Tests, Shanghai Jiaotong University, Shanghai 200030, PR China b R&D Testing Center, Technology Center, Baoshan Iron & Steel Co. Ltd. Shanghai 201900, PR China Accepted 1 April 2003
Abstract In the high temperature shape memory alloy (HTSMA) Ti51Pd30Ni19, superelasticity is found for the first time and the superelasticity is quite different from that of NiTi or CuZnAl. The recoverable superelastic strain is 7% without failure of the specimen. Shape memory effect (SME) at room temperature is evaluated as 7.2% with the recovery rate of 100%. The phase transformation temperatures were determined using differential scanning calorimeter (DSC). It has been found that the temperatures are changed during the process of alloy preparation. XRD analyses indicate that the macro- residual stress is eliminated and precipitation occurred after recovery treatment. The characteristics of both mechanical and transformation behavior are associated with the changes in the recovery process. Accordingly, an effective way to obtain the superelasticity is pointed out. # 2003 Elsevier Ltd. All rights reserved. Keywords: A. Intermetallics; B. Mechanical properties at high temperatures; B. Phase identification; B. Martensitic transformations; B. Shapememory effects (including superelasticity)
1. Introduction In Ti50PdxNi(50x) alloys, it has been found that the TMT (temperature of martensitic transformation) increases with the increase of Pd content from 20 to 50 at.%. Among these alloys, Ti50Pd30Ni20 has received much more attention because of its high TMT, which is higher than 473 K. This temperature is high enough for engineering application as a high temperature shape memory alloy (HTSMA) [1–7]. It is known that shape memory effect and superelasticity (or pseudoelasticty) are two important characteristics of shape memory alloys. However, the latter property for HTSMA is rarely obtained because the critical stress for slip (CSS) is low [8]. To solve this problem, alloying, aging induced precipitation, and thermomechanical treatment may be effective ways [5,8]. Using the first two methods, superelasticity has not yet been found in TiPd-based HTSMA. However, Golberg et al. [2] found the tendency of superelasticity in Ti50Pd30Ni20 specimens prepared by hot-rolling, cold-rolling, and annealing at 673 K, in sequence. In such a sequence, large numbers of * Corresponding author. E-mail address: [email protected] (Q. Tian). 0966-9795/03/$ - see front matter # 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0966-9795(03)00075-X
dislocations introduced by rolling, after experiencing thermal rearrangement, lead to a high value of CSS. So it is possible that stress induced transformation occurs before slip deformation, resulting in superelasticity. Shimizu et al. [5] improved the shape memory effect of TiNiPd alloys by precipitation hardening. They changed the alloy composition and obtained homogeneously distributed fine precipitates, which increases the CSS. In order to obtain perfect superelastic cycling of HTSMA, the composition of Ti50Pd30Ni20 was adjusted with reference to the methods mentioned above.
2. Experimental procedure Ti51Pd30Ni19 alloy was made by arc melting 99% Ti, 99% Ni and 99.9% Pd on a water-cooled copper mould under a controlled argon atmosphere. The ingot was melted four times for homogenization. Then, it was homogenized in vacuum at 1273 K for 5 h (I). Finally, it was hot-rolled into plate of 1.1 mm thickness at 1073 K, and subsequently cold-rolled into 1 mm thick sheets (II). Tensile specimens were spark cut along the rolling direction, as shown in Fig. 1. The specimens were mechanically polished and then annealed at 673 K in quartz capsule filled with argon for 1 h. After annealing,
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the specimens were quenched into ice water by crushing the capsule (III). In order to find out whether the transformation temperature is changed during the preparation process, specimens for differential scanning calorimetry (DSC) were also prepared. X-ray diffraction (XRD) measurements were performed using ‘D/max IIIA’ automatic diffractometer. Experiments were carried out at room temperature first; then the second measurement was performed as soon as the specimen was slowly heated to 673 K; the diffractometer ran the third scan after the specimens were isothermal for 1 h. With such a specific procedure imitating the annealing process, we can determine whether there were precipitates formed in the recovery process at high austenite temperature. As we know, martensite may exhibit a complicated XRD spectrum of many peaks because of the lower crystal symmetry and the texture of the specimen, therefore, it is difficult to confirm whether the precipitation occurred or not. XRD profile of another specimen in the recovered-state was also measured for comparison in order to find the changes of the residual stress. Tensile tests were carried out using ‘Shimadzu Autograph’ tensile machine at room temperature, 503 and 513 K, respectively. The strain rate was 1104/s. The displacement of the pin, which was inserted into the hole of the specimen, was recorded. The deformation outside the gauge section was ignored in the experiments. At first the specimen was tested at room temperature, in which the alloy is in the martensitic state, and then the specimens were heated to (Af +100) K, thus the martensite transformed to parent phase, in this way, the recovery rate was determined.
tioned previously. The experiments were run under argon atmosphere. Fig. 2 shows the DSC results. In Fig. 2, the numbers 1, 5 and 10 refer to the heating rates of 1, 5 and 10 K/min, respectively. It can be found that the temperatures corresponding to the peak of heat flow curves are the same for the different heating rates (Fig. 2a); the peaks for II correspond well to that for I (Fig. 2b). Fig. 2c represents the transformation temperatures of the alloy in different states during the process of alloy
3. Results and discussion 3.1. Phase transformation behavior Three kinds of specimens for DSC measurements were cut from the alloy in the solution treated state (I), cold-rolled state (II) and recovered state (III) as men-
Fig. 1. The dimensions of the specimen (units, mm).
Fig. 2. The transformation temperatures determined by DSC, the heat flow curves of specimen (a) in solution treated state, (b) under different heat treatment conditions, and (c) the changes of transformation temperatures during the process of alloy preparation, the numbers 1, 5 and 10 in the figure refer to the heating rate.
J. Wu, Q. Tian / Intermetallics 11 (2003) 773–778
preparation. It can be seen that the temperatures for II are similar to that for I, while the temperatures for III are much lower. It may be deduced that the reason why the transformation temperature is changed is that precipitation happens during the annealing process. Because the composition of the matrix may be changed due to precipitation, the transformation temperatures may in turn be affected. The changes may be closely related to the complicated effect of high-density dislocations and the internal stress produced during cold rolling as well. The alloy is basically perfect crystal with few defects after solution treatment at high temperature, and the transformation temperatures measured are close to the intrinsic values. The onset or finish temperatures of phase transformation may vary owing to the different heating rates, but the values corresponding to the peaks remain unchanged. For cold-rolled alloy, on the one hand, high dislocation density provides favorable sites for nucleation during solid-state transformation. On the other hand, the internal compressive stress impedes the phase transformation. The so two effects may maintain equilibrium in the cold-rolled alloy, so the transformation temperatures and the peak temperature do not change a lot. So long as the alloy experiences recovery treatment, the internal stress would be eliminated while the high-density dislocations just adjust appropriately, as indicated by the following X-ray analyses. Thus, the equilibrium was upset, which results in the decrease of the transformation temperatures. 3.2. XRD analyses The structures of TiPdNi alloys were studied by Enami et al. [9] and Lo et al. [10], etc. The structure of the martensite in Ti50PdxNi(50x) (x!15 at.%) alloys was determined as orthorhombic B19 (2H) type while the structure of austenite is of B2 (CsCl type). The XRD spectrum can be indexed with reference to the lattice parameters of parent phase a=0.31 nm (B2), of martensite a=0.278 nm, b=0.489 nm, c=0.459 nm [11], and of the precipitates (Ti2Pd) a=0.309 nm, c=1.0054 nm [12], respectively. Fig. 3 represents the XRD results. The profiles of martensitic phase in both cold-rolled state and recovered state exist only as two clear peaks, which can be identified as (120)B19 and (111)B19 (Fig. 3a). The crystal indices of austenitic phase were identified in Fig. 3b. It can be found from the figure that the specimens exhibit apparent texture, and a strong (110)B2 growth texture, which is similar to the observation of Mathews et al. [13]. Fig. 3a shows that after recovery treatment, the peaks of both (120)B19 and (110)B19 in the cold-rolled alloy move to higher angles, resulting from the increase of spacing between the crystal planes. This means that the macro-residual stress introduced by cold-rolling was
775
compressive, and was eliminated after annealing. It can be also seen that the peaks in recovery state are higher than those in the cold-rolled state, indicating that the lattice distortion in the former is not as serious as that in the latter, or more martensite transformed in the recovered specimen. Fig. 3b shows that after the annealing process was completed, (11-1)Ti2Pd(Ni) peak appears. As we know, low volume fraction of precipitates is difficult to detect by XRD. So it can be deduced from this observation that much precipitation has occurred in the alloy. In fact, precipitates may form during alloy preparation, e.g. during cooling after homogenization, or during hot rolling. These homogeneously distributed precipitates are very fine, they will congregate and grow in the recovery process, and were detected by XRD. It can also be seen that the peaks of the austenitic phase grow a bit, and the reason is similar to that for martensite. 3.3. Tensile deformation Fig. 4a shows stress–strain curves obtained at room temperature. The arrows in the figure indicate recovery strains after heating the specimens to (Af +100) K followed by cooling to room temperature. Three virgin specimens were loaded to 450, 750 and 940 MPa, respectively. The total strain of the first specimen is 7.2% with a recovery rate of 100%, and that for the second specimen is much higher, 11%, and the SME became incomplete (95%). The third specimen was deformed to a total strain of 12.9% without fracture; however, the SME became very poor (55%). The residual strain at room temperature in the three experiments is 3, 6, and 6.8% with increasing applied stress. Fig. 4b represents the experimental results at 503 K. A virgin specimen was repeatedly deformed five times. As shown in Fig. 4b, in the second experiment, the specimen was loaded to the same load as that in the first experiment. However, the width between the loading curve and the unloading curve is narrower, and the stress–strain curve tends to be a cycle. In the followed experiments, the loads were gradually increased to a higher level, the specimen yields at a higher stress correspondingly, and the elastic strength limit increases. The phenomenon is the well-known Bauschinger effect, which is caused by the movement of dislocations. The experiment indicates that Bauschinger effect may promote the CSS to be high enough so that martensite can be induced by stress. 3.4. Superelastic cycling Inspired by the last experiment, we carried out another experiment at a higher temperature, that is, a virgin specimen was tested in the austenitic state. The evolution of superelastic cycling is represented in Fig. 5.
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It can be seen that the stress–strain cycle becomes stable after the specimen was trained five times under stress up to 530 MPa. Then, the cycle stress was elevated to 680 MPa, the cycle becomes stable after cycling proceeds three times. Fig. 5b shows selected cycles of the experiment. It can be clearly seen that the loading curves warp up with the increase of superelastic cycling. A further consideration is shown in Fig. 5c.
The superelastic stress–strain curve may be divided into several stages, as denoted in Fig. 5c, according to the apparently different characteristics of the slope. It is considered that the difference corresponds to different deformation mechanisms. The elastic deformation of austenite proceeds with the initial loading at stage I. Another straight slope is regarded as the elastic deformation of martensite, defined as stage III. The bend connecting the two stages is denoted as stage II, which is
Fig. 3. XRD profiles of Ti51Pd30Ni19, (a) at room temperature and (b) at austenitic state.
J. Wu, Q. Tian / Intermetallics 11 (2003) 773–778
777
Fig. 4. Stress–strain curves of Ti51Pd30Ni19 tested (a) at room temperature and (b) 503 K.
associated with the stress-induced martensitic transformation and the reorientation of martensitic variants. The unloading curve is observed to be nonlinear, and does not look like its loading part, so it is wholly denoted as stage IV. In this stage, the elastic recovery of austenite and martensite, and the superelastic recovery of the stress-induced transformation are achieved. It can be clearly seen that the superelastic cycle is greatly different to that of TiNi or CuZnAl, where the cycles usually exhibit a platform in stage II. Two aspects may contribute this difference, i.e. the kind of martensite and the precipitation introduced during the recovery treatment. It is known that the martensite in TiPd–Ni (B19) is of higher symmetry and hence has less variants than that in TiNi (B190 ). So it cannot accommodate the applied stress field as effectively as the B190 kind of martensite can. As for the latter reason, the stress introduced martensitic transformation is compressed owing to the stress field produced by precipitation. It can be concluded that large numbers of dislocations and precipitates, which increase the CSS, are two necessary conditions for obtaining superelasticity in TiPdNi HTSMA, while training make the ideal realizable.
Fig. 5. Superelastic cycling of Ti51Pd30Ni19, (a) tested at 513 K, (b) warp of the loading curves and (c) the definition of different deformation stages.
4. Conclusion
1. The phase transformation temperatures of the Ti51Pd30Ni19 alloy are different for the different states in the alloy preparation process. This is attributed to complicated reasons including internal defects and residual stress introduced
778
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during cold rolling, and the precipitates produced during annealing. 2. Superelasticity was obtained for the first time in TiPdNi alloys. A recoverable superelastic strain of 7% at high temperature, and a total strain of 7.2% with a recovery rate of 100% at room temperature, were obtained. The superelastic cycle of the alloy is different from that in NiTi or CuZnAl, probably due to the relatively higher symmetry of martensite and large amount of precipitates. 3. Superelasticity is obtained by introducing large numbers of dislocations and precipitates in the matrix of TiPdNi alloys and by training as well.
Acknowledgements This work is sponsored by the Science and Technology Commission of the Shanghai Municipal Government, No. 00JC14055.
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