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The two way memory effect in TiNi alloys
Scripta Materialia, Vol. 35, No. 3, pp. 349-354, 1996 Elsevier Science Ltd Copyright 8 1996 Acta Metallurgica Inc. Printed in the USA. All rights reserved 1359-6462/96 $12.00 + .OO
Pergaimon PI1 S1359-6462(96)001625
THIE TWO WAY MEMORY EFFECT IN TiNi ALLOYS Peter Filip and Karel Mazanec Institute of Materials Engineering, Technical University of Ostrava, 708 33 Ostrava-Poruba, Czech Republic (Received October 3, 1995) (Accepted January 16, 1996) Introduction The two way memory effect (TWME) in shape memory alloys represents a reversible spontaneous shape changes during cooling and heating processes. This is a consequence of reversible phase transformations observed without application of any external stresses. The TWME is usually obtained after a thermomechanical treatment often called training. Different training routes were described (l-5) but physical origin of TWME is still unclear.The two main mechanisms for the TWME discussed in literature are either based on residual stresses induced in matrix or on retained (stabilized) martensite (6-9). Based on the observations of mechanical behaviour, the TWME has been attributed to the oriented residual stresses accompanying the dislocation arrangements (6,7). The oriented residual stresses only favour the nucleation and growth of preferential variants of martensite since the residual stresses are relaxed by the accompanying shape change (6). The preferentially formed variants of martensite grow without any external stress asistance during cooling, which is connected with shape change, and these variants transform back to the parent phase during heating whereby shape recovery occurs. This process can be repeated in the next thermal cycles. Based on the assumption that the generation of dislocations is linked with the “true plasticity”, a residual plastic strain is connected with the TWME. A second idea of TWME mechanism is b,ased on the localized stabilization of preferentially oriented martensitic variants which are retained after the heating above the original A, temperature (1,6,9). The cause of the observed martensite stabilization is the dislocation generation. These dislocations do not allow the martensite to shrink and disappear completely at heating. During cooling, these stabilized small martensite plates grow and hence modify the arrangement of further variants formed during subsequent cooling. This mechanism requires an incomplete reverse transformation and the residual deformation is also considered as a prerequisite of the TWME. Stahnans et al. (3,s) distinguish between macro-stress influence (responsible for one way memory effect) and “microstructure asymetry” (important for TWME) and on the basis of the interpretation of mechanical properties they try to disprove previous explanations. Following their studies, the most important influence of the dislocations which are generated during training is not the development of internal oriented stresses but rather a microstructural anisotropy inducing a thermodynamic anisotropy. Manach and Favier, however, clearly demonstrated the presence of internal stresses in trained TiNi specimens recently (2). According to their study the most important role play the dislocations and their influence on the mobility of martensite/high-temperature interfaces. These ideas are in agreement with the results obtained in (lo), according which the dislocations play an important role at cycling and repeated transforrnations~ under load. Perkins and Muesing (11) described the presence of “ghost martensite” in
349
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trained specimens in which the TWME occurs. This structural feature has been mentioned with different chemical composition in (512) too, but no relation to the dislocation arrays has been found. The different opinions were published concerning the influence of R-phase formation in TiNi based alloys on the stability of TWME. Stachowiak and McCormick (13), have stated that the formation of R-phase diminishes the extent of obtained reversible strain as well as the stability of TWME. In contrast to this conclusion, the authors (14,15) observed enhanced stability of TWME in the cases when the R-phase formation preceds the martensitic transformation B2 - B 19’. Repeating the heating cycles at loading of the memory materials, the complex degradation mechanisms of TiNi alloys occurs (10). The dislocation generation in the matrix of TiNi alloys (work hardening) strongly influences the maximum level of generated stresses and the extent of reversible strain (16- 18). As a contribution to the explanation of physical principles of TWME, the present paper is devoted to the study of substructure after so called “hard training” (4,5) at which the specimens are deformed (E = 4%) and cyclically heated and cooled down at constrained conditions. Exuerimental The experimental material used in this study was a commercial Ti - 50.1 at% Ni alloy supplied by FIBRA Ltd. (Czech Republic) in the form of wire with diameter d = 2Smm produced by several cold drawing steps with intermediate annealings in vacuum at 800”C/30min/water. The wires were annealed at 9OO”C/2h/waterand subsequently deformed with the constant bending strain of 4% at room temperature (as received conditions). The thermomechanical training has consisted of 20 sequences each of them composed of heating the material to 100°C and subsequent cooling to the temperature of liquid nitrogen (-196°C). The bent specimens (e = 4%) were thermally cycled in constrained conditions without a possibility to change their shape during heating and cooling routine (4,5). The transformation temperatures T, (the temperature at which the rhombohedral R-phase starts to form), M, (the temperature of the B 19’martensite start) and Ar (the temperature at which austenitic B2 phase finishes) of used materials correspond to its application object in human medicine and were established from the measurements of electric resistance p versus temperature T dependences. The microstructure of specimens was studied by a TEM Jeol 200CX. Thin foils were finished by a twin jet polisher in an electrolyte of HClO, and CH,COOH at U = 15V and temperature T F:0°C and some of them were heated above Ar temperature to eliminate the effects of deep undercooling.
Results and Discussion The transformation temperatures of both specimen types were established from p vs T measurements as
shown in Figure 1 and are given in Table 1. It is apparent that the training process is connected with decrease of the M, temperature and increase of the TR and Af temperatures, respectively. We suppose that this phenomena correspond to the work hardening during training procedure. It has been found that the optimal conditions for R-phase formation are given by higher internal stresses in the matrix (6,14,15,17). The increase of internal stresses as a consequence of enhanced dislocation density well corresponds to the observed behaviour. The mobility of B2/B19’ interfaces is lowered and the decrease of M, as well as the increase of Ar can be explained on this basis similar as in the case of work hardened specimens without training (6,16,17). Moreover, Piao (19), Otsuka (20) have shown that by reorientation of multivariant martensite into a single one, the stored elastic energy is eliminated and the high-temperature phase starts to form and finishes at higher temperatures. The reorientation process (preferential formation of martensite) is typical for training process as will be shown later. The substructure of the material in as received conditions observed at room temperature is shown in Figure 2. It is evident that the substructure is dominantly formed of internally twinned martensite B19’.
Vol. 35, No. 3
SHAPE MEMORY EFFECT
Figure 1. Electrical kstance
351
p vs temperature T dependencies obtained in as received conditions (a) and after 20 training cycles (b).
The R-phase wa.snot detected by using of TEM and most probably the R - B 19’ transformation occurred during cooling to T = 0°C (preparation of thin foils). The substructure of “hard trained” specimens changes dramatically as can be seen from Figure 3, depicting the typical substructure observed after 20 training sequences, martensite plates have a preferential orientation. We did not observe typical internal twinning any more. Instead of it, the individual martensite plates seem to be single variants with similar relationship to the matrix. It means that during repeating the working cycles, the preferentially oriented martensite variants are formed directly and the reorientation based on the movement of twin boundaries is lacking. The increased dislocation density after training procedure and heating of thin foils above Ar is shown in Figure 4. The dislocation arrays favour probably the repeated formation of particular martensite variants formed during training even when the material is transformed under no stress. The residual stresses are relaxed by the accompanying shape change (6). Based on the assumption that the generation of dislocations is linked with the “true plasticity”, a residual plastic strain is connected with the TWME. After heating of thin foils to 100°C (above original A,= 60°C) and cooling down to room temperature, the regions of stabilized martensite B 19’ were observed (Figure 5). In more works (1,6,9), the explanation of TWME mechanism is based on the localized stabilization of preferentially oriented martensitic variants which are retained after the heating above the original Af temperature. These dislocations do not allow the martensite to shrink and disappear completely at heating. According (1,6,9), these stabilized small martensite plates should start to grow first during cooling and hence modify the arrangement of further variants. This mechanism requires an incomplete reverse transformation and the residual deformation is also considered as a prerequisite of TWME. The cause of the observed martensite stabilisation is given by increase of dislocation density in matrix (20). The shrinkage of these martensite plates is limited by interaction with dislocations. Similar mechanism will play a role at cooling and we do not suppose that these stabilized variants can grow preferentially because the frictional stresses are too high. But it is evident that the stabilized martensite introduces a structural anisotropy and hence may play an important role at TWME.
Transformation
in
Temperatures
TABLE 1 of Investigated
specimen type as received conditions after traiuing
Alloy at Different Conditions T&‘C] 13 31
MJ’C] 9 0
A#C] 40 60
Figure 2. Substructure of the specimens in as received conditions and cooling to T = 0-C (preparation procedure).
Figure 3. Substructure after 20 training cycles and cooling to T r; O’C (preparation procedure).
The very important characteristic of the trained specimens is a special contrast observed by using of TEM. It concerns the so called “ghost martensite” depicted in Figure 6a. The structure in such regions is formed of high temperature phase B2 as follows from the diffraction pattern (Figure 6b). The alternate bands have not any direct correlation to the dislocations observed in the structure. The physical nature of “ghost martensite” is not clear yet. The “ghost bands” have earlier observed Olson et al. (2 1) and Lee (22) in steels and they have shown that this feature results from microsegregation of alloying elements. These “ghost bands” had dimensions of impurity segregation profiles (some hundred urn) but in the case of TiNi alloys the width of bans is substantially smaller. We do not expect the similar mechanisms and we suppose that the diffraction contrast could be connected with changes on the atomic level. This substructure feature probably represents a premartensitic stage. Some authors have described the lattice instability in TiNi based alloys (23,24). Even if the diffraction pattern (Figure 6b) does not allow us to analyse reliably the lattice modulations, we could suppose that the “ghost martensite” represents the matrix instability preceeding the B2 - B19’ transformation. During TWME, the martensite plates are probably formed preferentially in orientation related to the alternate bands of “ghost martensite” (see Figure 6a) as has been proposed in Cu-based shape memory alloys (11). These variants are favoured with respect to the other crystallographical variants. The presence of the “ghost martensite” could be held for an important characteristic for TWME in more shape memory systems and may lead to the thermodynamic anisotropy as proposed in (3,s).
Figure 4. Increased dislocation density, observed after training process (thin foils were heated above A, and cooled down to the room temperature).
Figure 5. Stabilized B19’ martensite (thin foils were heated above Ar and cooled down to the room temperature).
Figure 6a. “Ghost Inartensite” observed after training process.
Figure 6b. Diffraction pattern corresponding the area of “ghost martensite”.
Conclusions It follows horn the results obtained after so called “hard training” procedure that the TWME is influenced by:
-
the presence of stabilized B 19’martensite which produces a structural anisotropy, the presence of dislocations which are formed and arranged in the course of training process; these dislocations influence the internal stresses distributicn linked with the preferential growth ofmartensite variants, modify the transformation temperatures and occurrence of R-phase as well as the nucleation of martensit’e, the formation of the “ghost martensite” which was observed after the training process; the B19’ martensite forms preferentially in orientation related to the alternate bands during cooling without applied stress.
-
The further study is devoted to the determination of hierarchy of above mentioned individual parameters. Acknowledgments The authors are grateful to Grant Agency of Czech Republic (GACR) for financial support (project 106/93/0736). References J.Perkins and R.O.Sponholz, Metall.Trans. 15A, 313 (1984). P.Y.Manach and D.Favier, Scr.MetaJl.Mater. 28, 1417 (1993). RStalmans, J. Van Humbeeck and L.Delaey, Acta Metall.Mater. 40, 501 (1992). P.Pacholek, P.Filip and K.Mazanec, Metallic Materials 33, 129 (1995) (in Czech). P.Filip and KMazanec, Int.Jnl.of Materand Product Technology, accepted for publication (1995). P.Filip, “Physical Metallurgy Parameters of Shape Memory Effects in TiNi-type Shape Memory Alloys and Potential Practical Uses of this Effect”, Ph.D. thesis, TU Ostrava, (1988) (in Czech). 7. J. Perkins. and D.Hodgson, “The Two-Way Shape Memory Effect”, Engineering Aspects of Shape Memory Alloys, T.W.Duerig, K..N.Melton, DStoeckel and C.M.Wayman (eds.), Butterworth-Heinemann, 195 (1990). 8. R.Stalmans, J.‘van Humbeeck and L.Delaey, Acta Metall.Mater. 40,292l (1992) 9. TSaburi and S.Nenno, Scr.Metall. 8, 1363 (1974). 10. P.Filip and KMazanec, Scripta Metall.et Mater. 30,67 (1994). 11. J.Perkins and W.E.Muesing, Metall.Trans. 14A, 33 (1983) 12. P.Filip and A.C.Kneissl, PraktMetallographie 26, 201 (1995). 1. 2. 3. 4. 5. 6.
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13. G.B.Stachowiak and P.McCormick, Scripta Metall. 21,403 (1987). 14. Ch.Yiuying, Ch.Jinfang, L.Rundong, Y.Liping and Z.Ming, “Effect of Long Term Bending Deformation on the Shape Memory Characteristics ofTiNi Allos”, Shape Memory Materials’94, Y.Chu and H.Tu (eds.), Intemat.Academic Publishers, Beijing, 181 (1994). 15. Z. Jinbo, S.Xu, W.Fusheng and H.Yongjian, “Effect of Annealing on Characters of TiNi SMA”, Shape Memory Materials’W, Y.Chu and H.Tu (eds.), Intemat.Academic Publishers, Beijing, 176 (1994). 16. P.Filip, V.Matjsek and K.Mazanec, Zeitschr.Metallkde. 83, 877 (1992). 17. P.Filip, J.Rusek and K.Mazanec, Zeitschr.Metallkde, 82,488 (1991). 18. M.Piao, K.Otsuka S.Miyazaki and H.Horikawa, Materials Trans. JIM, 34,919 (1993). 19. K.Otsuka, “Recent Developments of Ti-Ni and Ti-Ni Based Ternary Shape Memory Alloys” Shape Memory Materials’94, Y.Chu and H.Tu (eds.), InternatAcademic Publishers, Beijing, 129 (1994). 20. P.Filip and K.Mazanec, Mater.Sci.Eng. A127, L19 (1990). 21. G.B.Olson, A.A.Anctil, T.S.DeSisto and E.B.Kula, Metall.Trans. 14A, 1661 (1983). 22. E.U.Lee, MetalLand Mater.Trans. 26A, 1313 (1995). 23. C.M.Hwang, M.Meichle, M.B.Salamon and C.M.Wayman, Phil.Mag. 47A, 177 (1983). 24. P.Moine, G.M.Michal and R.Sinclair, Acta.Metall. 30, 109 and 125 (1982).
Pergaimon PI1 S1359-6462(96)001625
THIE TWO WAY MEMORY EFFECT IN TiNi ALLOYS Peter Filip and Karel Mazanec Institute of Materials Engineering, Technical University of Ostrava, 708 33 Ostrava-Poruba, Czech Republic (Received October 3, 1995) (Accepted January 16, 1996) Introduction The two way memory effect (TWME) in shape memory alloys represents a reversible spontaneous shape changes during cooling and heating processes. This is a consequence of reversible phase transformations observed without application of any external stresses. The TWME is usually obtained after a thermomechanical treatment often called training. Different training routes were described (l-5) but physical origin of TWME is still unclear.The two main mechanisms for the TWME discussed in literature are either based on residual stresses induced in matrix or on retained (stabilized) martensite (6-9). Based on the observations of mechanical behaviour, the TWME has been attributed to the oriented residual stresses accompanying the dislocation arrangements (6,7). The oriented residual stresses only favour the nucleation and growth of preferential variants of martensite since the residual stresses are relaxed by the accompanying shape change (6). The preferentially formed variants of martensite grow without any external stress asistance during cooling, which is connected with shape change, and these variants transform back to the parent phase during heating whereby shape recovery occurs. This process can be repeated in the next thermal cycles. Based on the assumption that the generation of dislocations is linked with the “true plasticity”, a residual plastic strain is connected with the TWME. A second idea of TWME mechanism is b,ased on the localized stabilization of preferentially oriented martensitic variants which are retained after the heating above the original A, temperature (1,6,9). The cause of the observed martensite stabilization is the dislocation generation. These dislocations do not allow the martensite to shrink and disappear completely at heating. During cooling, these stabilized small martensite plates grow and hence modify the arrangement of further variants formed during subsequent cooling. This mechanism requires an incomplete reverse transformation and the residual deformation is also considered as a prerequisite of the TWME. Stahnans et al. (3,s) distinguish between macro-stress influence (responsible for one way memory effect) and “microstructure asymetry” (important for TWME) and on the basis of the interpretation of mechanical properties they try to disprove previous explanations. Following their studies, the most important influence of the dislocations which are generated during training is not the development of internal oriented stresses but rather a microstructural anisotropy inducing a thermodynamic anisotropy. Manach and Favier, however, clearly demonstrated the presence of internal stresses in trained TiNi specimens recently (2). According to their study the most important role play the dislocations and their influence on the mobility of martensite/high-temperature interfaces. These ideas are in agreement with the results obtained in (lo), according which the dislocations play an important role at cycling and repeated transforrnations~ under load. Perkins and Muesing (11) described the presence of “ghost martensite” in
349
350
SHAPE MEMORY EFFECT
Vol. 35, No. 3
trained specimens in which the TWME occurs. This structural feature has been mentioned with different chemical composition in (512) too, but no relation to the dislocation arrays has been found. The different opinions were published concerning the influence of R-phase formation in TiNi based alloys on the stability of TWME. Stachowiak and McCormick (13), have stated that the formation of R-phase diminishes the extent of obtained reversible strain as well as the stability of TWME. In contrast to this conclusion, the authors (14,15) observed enhanced stability of TWME in the cases when the R-phase formation preceds the martensitic transformation B2 - B 19’. Repeating the heating cycles at loading of the memory materials, the complex degradation mechanisms of TiNi alloys occurs (10). The dislocation generation in the matrix of TiNi alloys (work hardening) strongly influences the maximum level of generated stresses and the extent of reversible strain (16- 18). As a contribution to the explanation of physical principles of TWME, the present paper is devoted to the study of substructure after so called “hard training” (4,5) at which the specimens are deformed (E = 4%) and cyclically heated and cooled down at constrained conditions. Exuerimental The experimental material used in this study was a commercial Ti - 50.1 at% Ni alloy supplied by FIBRA Ltd. (Czech Republic) in the form of wire with diameter d = 2Smm produced by several cold drawing steps with intermediate annealings in vacuum at 800”C/30min/water. The wires were annealed at 9OO”C/2h/waterand subsequently deformed with the constant bending strain of 4% at room temperature (as received conditions). The thermomechanical training has consisted of 20 sequences each of them composed of heating the material to 100°C and subsequent cooling to the temperature of liquid nitrogen (-196°C). The bent specimens (e = 4%) were thermally cycled in constrained conditions without a possibility to change their shape during heating and cooling routine (4,5). The transformation temperatures T, (the temperature at which the rhombohedral R-phase starts to form), M, (the temperature of the B 19’martensite start) and Ar (the temperature at which austenitic B2 phase finishes) of used materials correspond to its application object in human medicine and were established from the measurements of electric resistance p versus temperature T dependences. The microstructure of specimens was studied by a TEM Jeol 200CX. Thin foils were finished by a twin jet polisher in an electrolyte of HClO, and CH,COOH at U = 15V and temperature T F:0°C and some of them were heated above Ar temperature to eliminate the effects of deep undercooling.
Results and Discussion The transformation temperatures of both specimen types were established from p vs T measurements as
shown in Figure 1 and are given in Table 1. It is apparent that the training process is connected with decrease of the M, temperature and increase of the TR and Af temperatures, respectively. We suppose that this phenomena correspond to the work hardening during training procedure. It has been found that the optimal conditions for R-phase formation are given by higher internal stresses in the matrix (6,14,15,17). The increase of internal stresses as a consequence of enhanced dislocation density well corresponds to the observed behaviour. The mobility of B2/B19’ interfaces is lowered and the decrease of M, as well as the increase of Ar can be explained on this basis similar as in the case of work hardened specimens without training (6,16,17). Moreover, Piao (19), Otsuka (20) have shown that by reorientation of multivariant martensite into a single one, the stored elastic energy is eliminated and the high-temperature phase starts to form and finishes at higher temperatures. The reorientation process (preferential formation of martensite) is typical for training process as will be shown later. The substructure of the material in as received conditions observed at room temperature is shown in Figure 2. It is evident that the substructure is dominantly formed of internally twinned martensite B19’.
Vol. 35, No. 3
SHAPE MEMORY EFFECT
Figure 1. Electrical kstance
351
p vs temperature T dependencies obtained in as received conditions (a) and after 20 training cycles (b).
The R-phase wa.snot detected by using of TEM and most probably the R - B 19’ transformation occurred during cooling to T = 0°C (preparation of thin foils). The substructure of “hard trained” specimens changes dramatically as can be seen from Figure 3, depicting the typical substructure observed after 20 training sequences, martensite plates have a preferential orientation. We did not observe typical internal twinning any more. Instead of it, the individual martensite plates seem to be single variants with similar relationship to the matrix. It means that during repeating the working cycles, the preferentially oriented martensite variants are formed directly and the reorientation based on the movement of twin boundaries is lacking. The increased dislocation density after training procedure and heating of thin foils above Ar is shown in Figure 4. The dislocation arrays favour probably the repeated formation of particular martensite variants formed during training even when the material is transformed under no stress. The residual stresses are relaxed by the accompanying shape change (6). Based on the assumption that the generation of dislocations is linked with the “true plasticity”, a residual plastic strain is connected with the TWME. After heating of thin foils to 100°C (above original A,= 60°C) and cooling down to room temperature, the regions of stabilized martensite B 19’ were observed (Figure 5). In more works (1,6,9), the explanation of TWME mechanism is based on the localized stabilization of preferentially oriented martensitic variants which are retained after the heating above the original Af temperature. These dislocations do not allow the martensite to shrink and disappear completely at heating. According (1,6,9), these stabilized small martensite plates should start to grow first during cooling and hence modify the arrangement of further variants. This mechanism requires an incomplete reverse transformation and the residual deformation is also considered as a prerequisite of TWME. The cause of the observed martensite stabilisation is given by increase of dislocation density in matrix (20). The shrinkage of these martensite plates is limited by interaction with dislocations. Similar mechanism will play a role at cooling and we do not suppose that these stabilized variants can grow preferentially because the frictional stresses are too high. But it is evident that the stabilized martensite introduces a structural anisotropy and hence may play an important role at TWME.
Transformation
in
Temperatures
TABLE 1 of Investigated
specimen type as received conditions after traiuing
Alloy at Different Conditions T&‘C] 13 31
MJ’C] 9 0
A#C] 40 60
Figure 2. Substructure of the specimens in as received conditions and cooling to T = 0-C (preparation procedure).
Figure 3. Substructure after 20 training cycles and cooling to T r; O’C (preparation procedure).
The very important characteristic of the trained specimens is a special contrast observed by using of TEM. It concerns the so called “ghost martensite” depicted in Figure 6a. The structure in such regions is formed of high temperature phase B2 as follows from the diffraction pattern (Figure 6b). The alternate bands have not any direct correlation to the dislocations observed in the structure. The physical nature of “ghost martensite” is not clear yet. The “ghost bands” have earlier observed Olson et al. (2 1) and Lee (22) in steels and they have shown that this feature results from microsegregation of alloying elements. These “ghost bands” had dimensions of impurity segregation profiles (some hundred urn) but in the case of TiNi alloys the width of bans is substantially smaller. We do not expect the similar mechanisms and we suppose that the diffraction contrast could be connected with changes on the atomic level. This substructure feature probably represents a premartensitic stage. Some authors have described the lattice instability in TiNi based alloys (23,24). Even if the diffraction pattern (Figure 6b) does not allow us to analyse reliably the lattice modulations, we could suppose that the “ghost martensite” represents the matrix instability preceeding the B2 - B19’ transformation. During TWME, the martensite plates are probably formed preferentially in orientation related to the alternate bands of “ghost martensite” (see Figure 6a) as has been proposed in Cu-based shape memory alloys (11). These variants are favoured with respect to the other crystallographical variants. The presence of the “ghost martensite” could be held for an important characteristic for TWME in more shape memory systems and may lead to the thermodynamic anisotropy as proposed in (3,s).
Figure 4. Increased dislocation density, observed after training process (thin foils were heated above A, and cooled down to the room temperature).
Figure 5. Stabilized B19’ martensite (thin foils were heated above Ar and cooled down to the room temperature).
Figure 6a. “Ghost Inartensite” observed after training process.
Figure 6b. Diffraction pattern corresponding the area of “ghost martensite”.
Conclusions It follows horn the results obtained after so called “hard training” procedure that the TWME is influenced by:
-
the presence of stabilized B 19’martensite which produces a structural anisotropy, the presence of dislocations which are formed and arranged in the course of training process; these dislocations influence the internal stresses distributicn linked with the preferential growth ofmartensite variants, modify the transformation temperatures and occurrence of R-phase as well as the nucleation of martensit’e, the formation of the “ghost martensite” which was observed after the training process; the B19’ martensite forms preferentially in orientation related to the alternate bands during cooling without applied stress.
-
The further study is devoted to the determination of hierarchy of above mentioned individual parameters. Acknowledgments The authors are grateful to Grant Agency of Czech Republic (GACR) for financial support (project 106/93/0736). References J.Perkins and R.O.Sponholz, Metall.Trans. 15A, 313 (1984). P.Y.Manach and D.Favier, Scr.MetaJl.Mater. 28, 1417 (1993). RStalmans, J. Van Humbeeck and L.Delaey, Acta Metall.Mater. 40, 501 (1992). P.Pacholek, P.Filip and K.Mazanec, Metallic Materials 33, 129 (1995) (in Czech). P.Filip and KMazanec, Int.Jnl.of Materand Product Technology, accepted for publication (1995). P.Filip, “Physical Metallurgy Parameters of Shape Memory Effects in TiNi-type Shape Memory Alloys and Potential Practical Uses of this Effect”, Ph.D. thesis, TU Ostrava, (1988) (in Czech). 7. J. Perkins. and D.Hodgson, “The Two-Way Shape Memory Effect”, Engineering Aspects of Shape Memory Alloys, T.W.Duerig, K..N.Melton, DStoeckel and C.M.Wayman (eds.), Butterworth-Heinemann, 195 (1990). 8. R.Stalmans, J.‘van Humbeeck and L.Delaey, Acta Metall.Mater. 40,292l (1992) 9. TSaburi and S.Nenno, Scr.Metall. 8, 1363 (1974). 10. P.Filip and KMazanec, Scripta Metall.et Mater. 30,67 (1994). 11. J.Perkins and W.E.Muesing, Metall.Trans. 14A, 33 (1983) 12. P.Filip and A.C.Kneissl, PraktMetallographie 26, 201 (1995). 1. 2. 3. 4. 5. 6.
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13. G.B.Stachowiak and P.McCormick, Scripta Metall. 21,403 (1987). 14. Ch.Yiuying, Ch.Jinfang, L.Rundong, Y.Liping and Z.Ming, “Effect of Long Term Bending Deformation on the Shape Memory Characteristics ofTiNi Allos”, Shape Memory Materials’94, Y.Chu and H.Tu (eds.), Intemat.Academic Publishers, Beijing, 181 (1994). 15. Z. Jinbo, S.Xu, W.Fusheng and H.Yongjian, “Effect of Annealing on Characters of TiNi SMA”, Shape Memory Materials’W, Y.Chu and H.Tu (eds.), Intemat.Academic Publishers, Beijing, 176 (1994). 16. P.Filip, V.Matjsek and K.Mazanec, Zeitschr.Metallkde. 83, 877 (1992). 17. P.Filip, J.Rusek and K.Mazanec, Zeitschr.Metallkde, 82,488 (1991). 18. M.Piao, K.Otsuka S.Miyazaki and H.Horikawa, Materials Trans. JIM, 34,919 (1993). 19. K.Otsuka, “Recent Developments of Ti-Ni and Ti-Ni Based Ternary Shape Memory Alloys” Shape Memory Materials’94, Y.Chu and H.Tu (eds.), InternatAcademic Publishers, Beijing, 129 (1994). 20. P.Filip and K.Mazanec, Mater.Sci.Eng. A127, L19 (1990). 21. G.B.Olson, A.A.Anctil, T.S.DeSisto and E.B.Kula, Metall.Trans. 14A, 1661 (1983). 22. E.U.Lee, MetalLand Mater.Trans. 26A, 1313 (1995). 23. C.M.Hwang, M.Meichle, M.B.Salamon and C.M.Wayman, Phil.Mag. 47A, 177 (1983). 24. P.Moine, G.M.Michal and R.Sinclair, Acta.Metall. 30, 109 and 125 (1982).