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Thermal cycling effects in high temperature Cu–Al–Ni–Mn–B shape memory alloys J. Font and J. Muntasell Departament de F´ısica i Enginyeria Nuclear, ETSII, Universitat Politècnica de Catalunya, Avda. Diagonal, 647 E-08028 Barcelona, Spain

J. Pons and E. Cesari Departament de F´ısica, Universitat de les Illes Balears, Ctra. de Valldemossa, km 7.5, E-07071 Palma de Mallorca, Spain (Received 3 June 1995; accepted 18 June 1996)

The effects of thermal cycling through the martensitic transformation have been studied in three Cu–Al–Ni –Mn –B high temperature shape memory alloys. An increase of the martensitic transformation temperatures with the number of cycles (up to , 7 K after 60 cycles) has been generally observed by DSC measurements. The microstructure of these alloys is rather complicated, with the presence of big manganese or aluminum boride particles and small boron precipitates, as well as the formation of dislocations during thermal cycling. By means of aging experiments, it has been shown that the evolution of transformation temperatures during cycling is mainly due to the step-by-step aging in parent phase accompanying the thermal cycling, and that the dislocations formed during cycling have only a very small effect, at least up to 60 cycles.

I. INTRODUCTION

In the last few years, a number of Cu–Al –Ni-based shape memory alloys exhibiting a thermoelastic martensitic transformation have been investigated for high transformation temperature applications, above 100 ±C, due to the relatively good stability of the Cu–Al –Ni alloys at these temperatures.1 However, the poor ductility of these alloys causes important problems for the machining of the material, mainly due to the grain coarsening and the g phase precipitation at the grain boundaries. To avoid these problems, several alloys with 4–5 components based on the Cu–Al–Ni system have been developed. For instance, the substitution of Al by Mn has been observed to prevent the g phase precipitation. Also, additions of B or Ti result in a significant grain refinement. As a result, alloys such as Cu–Al –Ni–Ti, Cu–Al –Ni–Mn –Ti, or Cu–Al –Ni –Mn–B present an improved ductility and also a good thermal stability even at temperatures between 100 ±C and 200 ±C.2–7 This temperature range cannot be achieved by the Ni–Ti alloys (the most extensively used for the shape memory applications). Then, the complete development of the above-mentioned alloys can open a new domain of shape memory applications. One of the most serious problems of Cu-based alloys in regard to commercial applications is the phenomenon called stabilization of martensite. This term is used to designate the increase of the reverse transformation temperatures observed after a quench and/or aging in the martensitic condition.8,9 The Cu–Zn –Al alloys are much affected by this phenomenon and, in some 2288

J. Mater. Res., Vol. 12, No. 9, Sep 1997

cases, increases of the reverse transformation temperatures as high as hundreds of K have been observed.10 The Cu–Al –Ni alloys are, however, less prone to stabilization.1 This problem has also been studied in the high temperature Cu–Al–Ni-based alloys and, in some cases, it can be reduced to a large extent.5,11–13 Another point to be considered is the stability of the transformation characteristics when the forward and reverse transformation, either thermal or stress-induced, are successively repeated. This procedure is known as thermal or pseudoelastic cycling, respectively. A large amount of research work has been devoted to the study of this topic in Cu-based shape memory alloys, mainly in Cu–Zn –Al14–17 and Cu–Al –Ni.18,19 To our knowledge, only a few works have been devoted to the study of thermal cycling in the high temperature Cu–Al–Nibased alloys, and they were restricted only to shortterm cycling.7,19,20 In the present paper we will show some results of a more extended thermal cycling in some Cu–Al –Ni –Mn–B alloys, followed by differential scanning calorimetry (DSC) and the corresponding microstructural changes studied by transmission electron microscopy (TEM). II. EXPERIMENTAL PROCEDURE

Three alloys with the following compositions (in at. %) have been used: alloy A, 70.2 Cu–24.9 Al– 2.9 Ni –2.0 Mn–0.15 B; alloy B, 70.0 Cu–23.9 Al– 2.9 Ni –3.1 Mn–0.15 B; and alloy C (not containing Mn), 69.7 Cu–26.0 Al–4.1 Ni–0.2 B. The nominal austenite start temperature for alloys A and B was about  1997 Materials Research Society

J. Font et al.: Thermal cycling effects in high temperature Cu – Al – Ni – Mn – B shape memory alloys

370 K, and for alloy C it was about 430 K, all of them being fully martensitic at room temperature. The alloys were hot extruded to rods of , 5 mm diameter. Different samples taken from each alloy were submitted to the following thermal treatments: (1) 20 min annealing at 1070 K and quench to water at room temperature; (2) the same as 1 followed by an aging at 570 K for 15 min in a melted salt bath and quench to water at room temperature; (3) 20 min annealing at 1070 K and air cooling to room temperature. The grain sizes of these materials were observed to be independent of the thermal treatment, ranging between 120 and 180 mm for alloy A, 250 and 300 mm for alloy B, and 80 and 100 mm for alloy C. Several cylindrical samples having , 5 mm diameter and ,2 mm thickness were cut from each alloy with a low-speed diamond saw. The different specimens (which will be denoted by a letter and a number, the letter standing for the alloy and the number standing for the thermal treatment; for example, A1 is a sample of alloy A submitted to treatment 1) were thermally cycled up to 60 times in the range 300 –440 K, where the complete direct and reverse martensitic transformation between the b-L21 and martensitic 18R structures took place. Due to their higher transformation temperatures, the samples taken from alloy C were cycled between 300 and 500 K. The cycling was performed in a Setaram DSC92 differential scanning calorimeter at a rate 2 Kymin. The microstructure of the samples was investigated by means of optical (Olympus BH) and transmission electron microscopy (Hitachi H600 at 100 kV). The specimen preparation for the OM consisted of an electropolishing in a mixture of 30% HNO3 in methanol at 230 K. The TEM samples were obtained by mechanical grinding up to a thickness of 0.1 mm and double-jet electropolishing in the above-mentioned electrolite until perforation.

III. EXPERIMENTAL RESULTS A. Evolution of transformation temperatures

The evolution of the transformation temperatures during thermal cycling is shown in Figs. 1, 2, and 3, corresponding, respectively, to the alloys A, B, and C. The values plotted in these figures are the temperatures corresponding to the peak (maximum or minimum) of the thermogram, which are denoted as M and A for the direct and reverse transformation, respectively. Taking into account that, at least up to 60 thermal cycles, the major change obtained in the thermograms is a shift in temperature, hardly changing the shape of the DSC curve (i.e., for each sample the differences Ms-Mf and Af-As are constant during cycling; see Figs. 4–6), the evolution of the peak temperatures can

FIG. 1. Plots of A and M values for samples A1 (w), A2 (h), and A3 (s).

be considered as representative of the transformation temperature evolution during cycling. In some cases it can be observed that the As values for the first reverse transformation just after the thermal treatment are very different than for the following cycles. This phenomenon was already known to occur in Cubased alloys and it corresponds to stabilization (when the first value of As is higher than the following ones) or it is called suppression (the opposite case, due to quenchedin disorder).21 A detailed study of these phenomena in the same alloys used for the present work was reported by Segui and Cesari.22,23 The present results show that in alloys A and B, the specimens quenched from 1070 K (thermal treatment 1) are stabilized by about 10 K, the specimens quenched and aged at 570 K (treatment 2) are suppressed again by about 10 K (as a consequence of the final quench from 570 K), and the specimens air cooled from 1070 K (treatment 3) are practically unaffected (sample A3 is stabilized only by 1 K and sample B3 is suppressed by 1 K). In the case of alloy C (not containing manganese), all the samples are suppressed by about 10 K (samples C1 and C2) or 5 K (sample C3). In the present work we will concentrate on the cycling effects, so the evolutions of the transformation


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FIG. 2. Plots of A and M values for samples B1 (w), B2 (h), and B3 (s).

FIG. 3. Plots of A and M values for samples C1 (w), C2 (h), and C3 (s).

temperatures during cycling will be considered from the third thermal cycle on (when the plots of Figs. 1, 2, and 3 show a regular evolution). For alloys A and B, all the specimens show an increase of the transformation temperatures of , 7 K after 60 thermal cycles, independently of the thermal treatment. For alloy C, the sample quenched from 1070 K (C1) has the same increase of , 7 K; the sample heated 15 min at 570 K shows little evolution with cycling. Finally, for the sample air-cooled from 1070 K there is a slight increase of the reverse transformation temperatures, leading to a widening of hysteresis. B. Microstructure of the alloys

One example of the microstructure of these alloys observed at room temperature in the optical microscope is shown in Fig. 7. It can be seen that the microstructure is formed by very thin martensite needles. In the majority of grains, most of the area is covered by martensite plates belonging to one self-accommodating group of variants and the remaining zones are filled by much smaller plates from other groups (fillings). Some particles with a lighter contrast can also be observed inside the grains in the optical micrographs. 2290

FIG. 4. Thermograms corresponding to the reverse transformation of sample A1 for cycles 1, 2, and 60.

The size of the particles (on the order of mm) is approximately the same in alloys A and B and slightly smaller in alloy C. The distribution of these particles is rather independent of the thermal treatment performed on the samples. In very similar alloys containing Mn,



FIG. 5. Thermograms corresponding to the direct transformation of sample B2 for cycles 1, 2, and 60.

FIG. 6. Thermograms corresponding to the reverse transformation of sample C1 for cycles 1, 2, and 60.

Morris observed the same kind of particles which, after EDS microanalysis, were identified as manganese or aluminum borides. The presence of these particles is probably due to difficulties of dissolution of boron.24 From TEM observations, the selected area electron diffraction patterns (SADP) confirm that the martensite has an 18R structure. All the specimens had a marked texture: the orientation of practically all the grains was very close to martensite zone axes which are derived from a h110jb plane of the parent phase; i.e., the direction of extrusion (perpendicular to the TEM foil) goes along h110jb . In all specimens the SADP show that the martensite has a next-nearest neighbor type of order; i.e., it is derived from a D03 or L21 ordered parent phase. The specimens taken from alloys A and B contain some holes after the double-jet polishing. These holes would be a consequence of a preferential polishing of the

FIG. 7. Optical micrograph showing the microstructure of alloy C air-cooled from 1120 K.

boride particles shown in Fig. 7. In the case of alloy C the same type of holes was again observed. However, in this alloy some particles with sizes of the order of mm were still visible in the thin foil [see Fig. 8(a)]. The majority of these precipitates were thicker than the surrounding matrix and were not transparent to the electron beam, but in some cases it was possible to obtain diffraction patterns from them. The patterns indicate that they are g precipitates [Fig. 8(b)]. This kind of precipitates was observed only in specimens from alloy C and not in alloys A and B. Thus, we can conclude that the particles observed in alloys A and B are manganese or aluminum borides. In alloy C (not containing manganese), some of the particles can be aluminum borides and the others are g precipitates. Additionally, another type of precipitation was observed in the three alloys A, B, and C, consisting of dispersions of very small precipitates (sizes of tens of nm) inside the martensite plates (as an example, see Fig. 9). The existence of the same type of particles was also reported by Morris24 and was identified as undissolved boron precipitates. Although the diffraction patterns do not show additional spots which would confirm that identification, the fact that the contrast of this precipitate is always very weak can support the nature of boron particles, due to the very different scattering factor of this element compared with the other components of the alloy. The matrix surrounding the small precipitates does not show any strain contrast (probably they are incoherent). These particles do not seem to interact in a significant way with the martensite plates. The microstructure of sample C3 presents some peculiarities. After the thermal treatment (air cooling from 1070 K) the TEM images show a very finely mottled contrast [Fig. 10(a)], which could represent the start of a matrix decomposition process. By aging the air-cooled


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(a) FIG. 9. TEM micrograph showing small boron particles, taken from alloy A.

(b) FIG. 8. (a) TEM micrograph showing a g-phase particle embedded in martensite, taken from sample C3. (b) Selected area diffraction pattern from the precipitate, k011lg zone axis.

samples at 520 K the decomposition of the b-matrix continues. A dense dispersion of very fine precipitates (, 5 nm) appears in the TEM micrographs after tens of hours at 520 K [Fig. 10(b)]. These precipitates are more dense at the matrix APB, in a similar way as g phase precipitates in Cu–Zn–Al single crystals.25 A more prolonged aging (up to 28 days at 520 K) finally leads to bainite formation [Fig. 10(c)]. The DSC measurements 2292

of aged samples also reflect the decomposition of the matrix. At the first stages of aging the martensitic transformation temperatures (Ms and Mf) increase and the reverse transformation temperatures (As, Af) also increase, but in a higher amount, leading to a broadening of hysteresis. For instance, after 92 h at 520 K, the Ms is raised by 40 K and the As by 70 K, which brings about an increase of hysteresis of 30 K. For a longer aging the effect of decomposition is stronger, with a considerable broadening of DSC peaks and reduction of transformation heat. In all the alloys, another microstructural aspect of the thermally cycled specimens is the presence of dislocations inside the martensite plates. The dislocation contrast is strongly dependent on the orientation of the foil (see Fig. 11), and in many cases it is mixed with the stacking fault contrast, which makes observation of the dislocations very difficult. Some specimens were in situ heated until the back-transformation to the parent b phase. In some areas which at room temperature (martensitic state) seemed to be dislocation-free, a dispersion of dislocations was visible after heating to the b phase (Fig. 12). It is already known that in Cu-based shape memory alloys, dislocations form after thermal cycling through the martensitic transformation. In the case of Cu–Zn–Al single crystals with subambient transformation temperatures, the dislocations formed by thermal cycling have been studied in detail14,15,26 (having an austenitic matrix makes the observation and analysis of the defects easier). In that case, the dislocations are of mixed type, with Burgers vector b k100lb and line direction u k111lb . These dislocations are probably formed by plastic deformation of martensite due to the



grain refined Cu–Al–Ni –Mn –B, the analysis of the dislocations is more difficult (at room temperature, the material is martensitic and the martensite plates are very small), but, in principle, it is reasonable to admit that the nature and origin of the dislocations observed would be the same as in the Cu–Zn–Al single crystals. The density of dislocations formed in the present alloys is comparable with that observed in Cu–Zn–Al for the same number of cycles,26 although a precise counting of the density has not been performed, due to the unclear contrast of the dislocations in martensite. IV. DISCUSSION (a)

(b)

(c) FIG. 10. TEM micrographs showing the evolution of microstructure of alloy C air-cooled from 1120 K and aged at 520 K. (a) Mottled contrast in the air-cooled alloy without aging at 520 K; (b) dark-field image of small precipitates obtained after 13 days aging; (c) bainite needles in the alloy aged for 28 days.

internal stresses originated during the transformation, specially the stresses coming from the difference in the volume of both phases.27 In the present case of

The effect of thermal cycling in Cu-based shape memory alloys has been the object of several studies in recent years. In Cu–Zn–Al alloys, the influence of thermal cycling up to 50–60 cycles is very small (1–2 K). After a few hundreds of cycles, the Ms and Af increase by 2–6 K with similar decreases of Mf and As, producing a widening of the transformation range Ms-Mf (or Af-As).17,28 These evolutions were interpreted as energetic effects of the dislocations formed during cycling.26 In another research, an increase of 10 K in Ms was observed in a B2-ordered Cu–Zn–Al alloy, whereas for a D03 -ordered alloy (with different composition), a decrease of 12 K in Ms was obtained. Changes in the parent phase ordering were considered as the mechanism responsible for these modifications during cycling.16 In Cu–Al –Ni alloys exhibiting a b –g′ martensitic transformation, an increase in Ms and Af of 10 –15 K was obtained after 20 cycles, while the Mf and As remained practically constant.19 These effects were interpreted in terms of the defects created during cycling which may assist the nucleation of the martensite and increase the friction. In a recent work on grain refined Cu–Al –Ni– Mn –B alloys with compositions similar to ours, an increase of transformation temperatures with the number of cycles has also been obtained. In this case, after only 8 thermal cycles an increase of transformation temperatures up to 25 K for quenched specimens (which is notably higher than the present results) and less than 10 K for specimens heated to 570 K for 30 min have been reported.20 In the paper, the behavior of the quenched specimens is related to the accommodation of the elastic strain energy which is made difficult by the incomplete order obtained after quenching. When the alloys have high transformation temperatures, as the present Cu–Al –Ni–Mn –B, the material is at a temperature of about 400–500 K after the reverse transformation has taken place, which could be high enough to activate the diffusional processes in the b phase. The process most likely to occur is an improvement of the degree of next-nearest neighbor ordering by reorientation of “wrong” pairs of atoms.21,29 The


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(a)

(a)

(b)

(b) FIG. 12. (a) Image taken from sample A1 at room temperature (martensitic state). (b) Same zone at 470 K (b phase, in situ heated).

(c) FIG. 11. Micrographs showing different contrasts of dislocations formed after 60 thermal cycles: (a) sample A1, (b) sample B1, and (c) sample B2. 2294

exchanged pairs of atoms are retained by the thermal treatment performed. In such a case, after the completion of one thermal direct and reverse transformation cycle, at the moment when the following direct transformation will start, the parent phase would have a slightly better degree of order than in the previous cycle, leading to a slightly higher transformation temperature. For a repetitive thermal cycling, this effect is cumulative, leading to a step-by-step-aging of the parent phase which could be responsible for the gradual increase of transformation temperatures. This mechanism could also explain the results of Morris and Gunter.20 The aging effect is overimposed to those related to microstructural changes produced by cycling, in particular the formation of dislocations.



In order to ascertain whether this aging effect is actually operating or not, additional experiments were performed. A new set of specimens was prepared with the same thermal treatments as before. Nevertheless, previously to the thermal cycling the specimens were aged 24 h at 420 K (in b phase) in a melted salt bath. After that, a single transformation cycle was recorded in the DSC followed by a new aging of 24 h at 420 K in the same DSC. Finally the samples were thermally cycled in the DSC as before. The additional aging at 420 K was estimated to be approximately equivalent to the cumulative aging in b phase performed during the thermal cycling up to 60 cycles. If the above-mentioned mechanism is actually operating, an increase of the transformation temperatures in relation to the values of the old samples would be expected. Additionally, after the aging at 420 K, the degree of next-nearest neighbors order would be improved to a considerable extent. Thus, for the subsequent step-by-step aging above Af produced by the thermal cycling routine, the cumulative reordering would be smaller than in the previous case; i.e., the increase of transformation temperatures during the thermal cycling of the new set of samples is expected to be smaller than the former.

FIG. 13. Plots of A and M temperatures for samples aged at 420 K: sample A5 (h); sample B4 (w); sample C6 (s).

The new set of samples will be labeled as follows: (4) treatment 1 plus aging at 420 K, (5) treatment 2 plus aging at 420 K, and (6) treatment 3 plus aging at 420 K. The results obtained for these specimens are summarized in Fig. 13. For alloys A and B aged at 420 K, all the transformation temperatures are higher than the previous noncycled samples and the evolution during cycling is largely reduced (slight increase of , 1 K), thus confirming the mechanism described above. In this case, the evolution of the transformation temperatures during cycling can be due to some “residual” reordering occurring in the parent phase, but it can also be an effect of the dislocations generated during cycling, in a similar way as in Cu–Zn –Al.26 For the new set of samples A and B submitted to the additional aging at 420 K, being the evolution of transformation temperatures and the distributions of dislocations comparable to those obtained in Cu–Zn–Al with the same number of cycles (only 1–2 K),17,26 it can be concluded that the role played by the dislocations is, in principle, equivalent in both alloys. The role of such dislocations, described in Ref. 26, has been mainly associated with relative chemical free energy changes of parent and martensite phases due to the different energies of the dislocations when they are embedded in each phase and to the energy of additional faults created in martensite by some of these dislocations. The behavior of the specimens taken from alloy C is different. For the samples quenched from 1070 K (sample C4), the aging at 420 K does not modify the transformation temperatures or the evolution during cycling (an increase by about 6 K is still found). The effect of the aging at 420 K is only the recovering of the suppression caused by the quench from 1070 K, which can be seen when comparing the As values of the 1st and 2nd cycles. Due to the higher transformation temperatures of this alloy, the thermal cycling is performed between 300 and 500 K. The subsequent evolution of transformation temperatures during cycling must be related to the step-by-step aging at 500 K caused by the thermal cycling, which brings the material to the equilibrium state of long-range order at 500 K. In fact, after some experiments of quenching these alloys from and aging to different temperatures, it has been obtained that the final state of order reached by the material has a notable influence on the transformation temperatures. This effect is particularly strong for alloy C, in which the permanent martensitic transformation temperatures of a specimen quenched from 600 K differs by about 100 K from those of a specimen quenched from 1070 K. For alloys A and B, the maximum changes of transformation temperatures were only 40 K.23 Concerning sample C5 (quenched from 1070 K and heated 15 min at 570 K), again the aging at 420 K does not alter the evolution during thermal cycling, which is nearly nonexistent. In


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this case, after heating to 570 K the disorder retained by the first quench from 1070 K is eliminated and the alloy acquires a better degree of long-range order, which is not modified by the subsequent step-by-step aging at 500 K brought about by thermal cycling. Again, the effect of the aging at 420 K is the elimination of the suppression caused by the last quench from 570 K. It is important to notice that, for alloys A and B the recovering of the suppression caused by the last quench from 570 K is slower than in alloy C (contrary to sample C2, in which there is still an evolution of the transformation temperatures of samples A2 and B2 during the subsequent thermal cycling). The behavior of alloy C air cooled from 1070 K (samples C3 and C6) is affected by the b phase decomposition explained in the last section. In this case, the evolution of transformation temperatures with thermal cycling is comparable to that found in equivalent samples aged at 520 K without thermal cycling, i.e., raise of transformation temperatures and increase of thermal hysteresis. So, the step-by-step aging at 500 K performed with the thermal cycling promotes a progressive decomposition process that was initiated during the slow cooling from 1070 K. At this point it is worth noting the very high sensitivity of the martensitic transformation temperatures to the microstructure of the material: after the thermal cycling or a short aging at 520 K, the TEM images do not show any significant difference with the slowly cooled samples (in both cases, a very fine unresolved mottled contrast was found), while the transformation temperatures indeed revealed an evolution. It is known that chemical and nonchemical effects are present when second phase particles act as additional obstacles for growth of the martensite plates. For example, the presence of dispersions of coherent or semicoherent g phase precipitates in b Cu–Zn –Al single crystals is known to influence the transformation characteristics. For relatively dense distributions of semicoherent g precipitates (size about , 1 mm) strong widening of hysteresis by about 30 –40 K and transformation temperature range (Ms-Mf or Af-As) have been obtained.30,31 Besides the slowly cooled samples of alloy C, as commented above, two types of precipitates are generally present in these Cu–Al–Ni –Mn –B alloys. The fine dispersions of small boron precipitates (as shown in Fig. 9) do not show a strong interaction with the martensite plates either in the cycled or in noncycled specimens. Then, these precipitates do not seem to have any noticeable influence in the evolution of the material during cycling. The other type of precipitates, i.e., the big aluminum or manganese borides, can certainly have some effect. From the chemical point of view, they can change, at least locally, the composition of the surrounding matrix and affect in this way the relative stability of the parent and martensite. Nevertheless, the density 2296

of the boride particle distribution is relatively small (for instance, it is much smaller than in the case of g precipitates in Cu–Zn –Al single crystals, which indeed considerably affect the martensitic transformation), so the possible effects of them are more localized and have a relative low influence on the global transformation. V. CONCLUSIONS

The thermal cycling effects on the martensitic transformation have been studied in several high temperature Cu–Al –Ni–Mn –B alloys. For the three thermal treatments studied (quench, quench plus flash heating, and slow cooling), a similar increase of the transformation temperatures with the number of cycles (up to 60 cycles) has been observed. Similarly with other Cu-based alloys, dislocations are formed during the thermal cycling. In spite of the complicated microstructure of these alloys (grain refined and containing two distinct distributions of precipitates: big manganese or aluminum boride particles and small boron precipitates), it has been proved that the evolution observed is mainly due to the step-bystep aging in the parent phase accompanying the thermal cycling, which produces a progressively improved nextnearest neighbor ordering. Once the reordering process is completed (or practically completed) by means of a prolonged aging in b phase, the evolution of transformation temperatures during the subsequent thermal cycling is strongly reduced. This fact indicates that, at least up to 60 cycles, the dislocations formed during cycling have only a small effect, comparable to what happens in other Cu-based alloys. ACKNOWLEDGMENT

Partial financial support from CICYT (Research project MAT93-0188) is gratefully acknowledged. REFERENCES 1. J. Van Humbeeck and L. Delaey, in The Martensitic Transformation in Science and Technology, edited by E. Hornbogen and N. Jost (DGM Informationsgesellschaft, Oberursel, Germany, 1989), p. 15. 2. K. Sugimoto, K. Kamei, H. Matsumoto, S. Komatsu, and T. Sugimoto, J. de Physique 43, C4-761 (1982). 3. G. N. Sure and L. C. Brown, Metall. Trans. 15A, 1613 (1984). 4. J. Van Humbeeck, M. Chandrasekaran, and L. Delaey, I.S.I.J. Int. 29, 388 (1989). 5. D. P. Dunne, J. Van Humbeeck, and L. Delaey, in Shape Memory Materials, edited by K. Otsuka and K. Shimizu (Mater. Res. Soc. Symp. Proc. 9, Pittsburgh, PA, 1989), p. 329. 6. S. Eucken, E. Kobus, and E. Hornbogen, Z. Metallk. 82, 640 (1991). 7. M. A. Morris and T. Lipe, Acta Metall. Mater. 42, 1583 (1994). 8. A. Abu-Arab, M. Chandrasekaran, and M. Ahlers, Scripta Metall. 18, 899 (1984). 9. M. Ahlers, in Proc. Int. Conf. on Martensitic Transformations, edited by I. Tamura (The Japan Institute of Metals, Sendai, Japan, 1987), p. 786.



10. A. Abu-Arab and M. Ahlers, J. de Physique 43, C4-709 (1982). 11. D. Dunne, J. Van Humbeeck, and M. Chandrasekaran, Mater. Sci. Forum 56 – 58, 463 (1990). 12. Y. Itsumi, Y. Miyamoto, T. Takashima, K. Kamei, and K. Sugimoto, Mater. Sci. Forum 56 – 58, 469 (1990). 13. C. Segu´ı and E. Cesari, Trans. Mater. Res. Soc. Jpn. 18B, 919 (1994). 14. D. R´ıos-Jara and G. Guénin, Acta Metall. 35, 109 (1987). 15. M. Sade, R. Rapacioli, and M. Ahlers, Acta Metall. 33, 487 (1985). 16. T. Tadaki, M. Takamori, and K. Shimizu, Transact. JIM 28, 120 (1987). 17. C. Auguet, E. Cesari, and Ll. Mañosa, J. Phys. D: Appl. Phys. 22, 1712 (1989). 18. Y. Nakata, T. Tadaki, and K. Shimizu, Transact. JIM 26, 646 (1985). 19. P. Fisher, D. Dunne, and N. Kennon, in Proc. Int. Conf. on Martensitic Transformations, edited by I. Tamura (The Japan Institute of Metals, Sendai, Japan, 1987), p. 946. 20. M. A. Morris and S. Gunter, Scripta Metall. Mater. 26, 1663 (1992).

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