Phase transformation characteristics of laser gas nitrided NiTi shape memory alloy

Surface & Coatings Technology 200 (2006) 5598 – 5605 www.elsevier.com/locate/surfcoat

Phase transformation characteristics of laser gas nitrided NiTi shape memory alloy H.C. Man a,*, N.Q. Zhao a,b a

Laser Processing Group, Advanced Manufacturing Technology Research Center, Department of Industrial and Systems Engineering, Hong Kong Polytechnic University, Hong Kong, China b School of Materials Science and Engineering, Tianjin University, China Received 18 March 2005; accepted in revised form 27 July 2005 Available online 22 September 2005

Abstract Laser gas nitriding (LGN) process is an efficient technique for modifying the surfaces of Ti and NiTi alloys. The process inevitably alters the composition, the microstructure and the properties of the alloyed surface. For components with thick cross sections, the changes of composition or microstructure at the surface may not affect the bulk properties of the materials. However, this may not be the same case if the cross-section thickness is small as compared to the thickness of the modified region. In this instance, when the cross section of the NiTi component is thin, the shape memory behaviour of NiTi alloy may be influenced as a result of the surface modification process. The present work aims to study the effect of the existence of the laser gas nitrided (LGN) layer of NiTi upon the phase transformation temperatures of the alloy. The changes in compositions in the LGN layer were chemically analysed using X-ray photoelectron spectroscopy and energy dispersive X-ray spectroscopy. The microstructure of the layer was analysed by X-ray diffraction. Differential Scanning Calorimetry (DSC) was used to measure the reversible B19Vto B2 transformation temperatures. The results showed that the matrix of LGN layer retains the same composition as the substrate alloy. The LGN process did not alter the phase transformation characteristics of the alloy. The LGN layer itself has the same transformation characteristics as the parent metal. D 2005 Elsevier B.V. All rights reserved. Keywords: NiTi SMA; TiN coating; Laser gas nitriding; Phase transformation

1. Introduction One of the unique characteristics of shape memory NiTi alloy is its biocompatibility due to the formation of a stable oxide layer on its surface [1]. However, NiTi exhibits poor resistance to localized corrosion in chloride-containing environments, and the healing of the passive film on NiTi has been reported to be a slow and difficult process. Moreover, the release of nickel, carcinogenicity and allergic hazards caused special concern [2]. In order to improve the biocompatibility of NiTi alloys, many surface modification techniques have been developed [3 –5]. Specifically, the formation of a titanium nitride (TiN)

* Corresponding author. Tel.: +852 2766 6629; fax: +852 2362 5267. E-mail address: [email protected] (H.C. Man). 0257-8972/$ - see front matter D 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2005.07.079

as a modified surface layer on NiTi substrate to improve the corrosion resistance has been reported by many researchers such as Fu et al. by pulsed high-energy density plasma (PHEDP) [6], Endo et al. by arc ion plating [7], Wu et al. by ion nitriding [8], Starosvetsdy and Gotman by power immersed reaction assistant coating (PIRAC) nitriding method [9], and Man et al. by laser gas nitriding (LGN)[10 – 12]. All these surface modification techniques inevitably alter the surface compositions, i.e., the ratio of Ni and Ti, and the microstructure of the alloy. The effect of surface modification upon the phase transition temperatures of the NiTi was, however, sparsely reported. It is well known that the shape memory behavior of the NiTi alloys, whose composition is close to the equiatomic ratio of nickel and titanium, is associated with a reversible transformation between the low-temperature martensitic NiTi phase (B19V) and the high-temperature

H.C. Man, N.Q. Zhao / Surface & Coatings Technology 200 (2006) 5598 – 5605

austenitic NiTi phase (B2), which occurs by a twinning mechanism [13]. The phase transition temperatures, which in simple systems are identified as Ms, Mf (start and finish of the formation of martensite on cooling) and As, Af (start and finish of the formation of austenite on heating), depend strongly on the composition proportion of Ni and Ti. In effect, the phase transition temperatures can be controlled by the Ni-content [14,15] as well as the thermal mechanical treatment [16]. Furthermore, the formation of an additional phase, termed the R-phase, is sometimes observed as an intermediate step between the transformation of B19V and B2 [17]. The purpose of this work is to investigate the effects of laser gas nitriding upon the surface chemical compositions and the phase transformation temperatures of NiTi alloys. The results are compared with the effects of conventional heat treatment of solution annealing and quenching followed by ageing.

2. Experimental 2.1. Materials The material used was a binary nickel-rich NiTi alloy with a nominal composition of 50.8 at.% Ni, Ti balance. The as-received hot rolled NiTi alloy in the form of a plate with a thickness of 3 mm was cut to 8  8  3 mm for the experiment. The specimen was ground clean using 120 grit SiC paper to remove any surface oxides or contaminants. A continuous wave 2 kW Nd –YAG laser was used to perform the laser gas nitriding process. Different process parameters such as laser power, beam size, scanning speed, and nitrogen gas flow rate were used to produce single melt tracks first. The cross-sections of these melt tracks were then examined by optical microscopy and the width and depth of these melt tracks were measured. The set of parameters that produced the track which has a hemispherical shape (for easy overlapping) with desired width and depth and free from crack and porosity was chosen. Based on this optimum set of parameters, a large nitrided surface was then produced by overlapping single tracks at 50% melt width interval. The optimum set of processing parameters selected in this work was laser beam power 500 W, 2 mm diameter defocused spot size, 150 mm focal length Zinc Selenide (ZnSe) focusing lens, high purity nitrogen gas (99.9%) N2 at 50 l/min discharging through a nozzle onto the laser melt pool, laser beam scanning speed 5 mm/s, and an overlapping interval of 0.5 mm. The specimens after the laser gas nitriding (LGN) treatment were solution treated at 800 -C for 4 h followed by water quenching. The quenched specimens were then aged at 480 -C for 1 h. The process of solution annealing and water quenching followed by aging is simply identified as heat treatment (HT) in this article.

5599

2.2. Characterization methods Scanning electron microscopy (SEM) with EDX (Leica Stereo Scan 440) was used to investigate the microstructure and the composition of the nitrided layer. A Bruker D8 Discover diffractometer was used for X-ray diffraction (XRD) study and the voltage –current settings of the equipment were 40 kV and 40 mA using CuKa radiation. The specimens were chemically analysed using a S-probe spectrometer (Perkin Elmer PHI 1600 ESCA System) operated in X-ray photoelectron spectroscopy (XPS) mode with monochromatized Al Ka radiation (hv = 1486.6 eV). The pass energies of 187.85 and 29.25 eV were used for the survey scan and narrow scan spectra, respectively. The experimental resolution of the narrow scan was 0.7 eV. The binding energy scale was calibrated to the C1s peak at 285.0 eV. The base pressures in the analyser and the preparation chambers were 10 8 Pa. Before the measurements, the surface of the specimen was argon etched with a sputtering rate of 5.6 nm/min for 0.5 min (EAr = 15 KeV, P = 300 W) in order to remove the natural protective titanium dioxide at the surface. For the Differential Scanning Calorimetry (DSC) experiments, a Perkin Elmer DSC 7 system was used. The specimens for DSC were sliced parallel to the surface of the LGN layer at a depth of 0.5 mm ( the LGN layer is about 300 – 400 Am). The specimens with masses between 10– 20 mg were initially cooled from room temperature to 50 -C, where they were held for 20 min to establish thermal equilibrium. Then the DSC measurements were started by heating to 100 -C at 10 -C/min. At 100 -C the specimen was then cooled to 50 -C with a cooling rate of 10 -C/min. For some specimens, the heating cycles were repeated for a second time. Dry nitrogen gas was used to flush the specimen chamber to prevent condensation of water vapour and oxidation of the NiTi. A second, empty aluminium pan served as an inert reference, and the apparatus was calibrated with an indium standard, as well as by the melting point of water.

3. Results and discussion 3.1. Microstructure Fig. 1a shows the cross-section microstructure of the LGN NiTi specimen. It can be seen that a continuous thin surface layer with thickness of about 2 Am was formed on the outmost surface of the specimen. A higher magnification of this TiN continuous surface is shown in Fig. 1b. Fig. 1c shows the cross-section microstructure at a location 40 Am below the outmost surface of Fig. 1a. It can be seen that TiN dendrites exist in the melted zone which is about 400 Am thick. The dendritic TiN phase distributes evenly in the NiTi matrix phase.

5600

H.C. Man, N.Q. Zhao / Surface & Coatings Technology 200 (2006) 5598 – 5605

Fig. 1. Cross section microstructure of LGN NiTi. (a) The cross section microstructure of the LGN specimen. (b) The TiN surface with the dendritic structure at the surface. (c) The microstructure at a location 40 Am below the outmost surface of (a).

Fig. 2 shows the microstructure of the overlap tracks at the top surface when 40 Am of the surface was ground off. It can be seen that the amount of TiN dendrites at the overlap

Fig. 2. The microstructure of the top surface of the overlap tracks after the TiN continuous layer was ground off.

region is slightly less than those in the central part of each individual track. Because of the near Gaussian nature of the laser beam, the centre part of the melt track is the last region to solidify. This provides a longer solidification time at the central region of the track which allows more nitriding reaction to occur and results in more TiN dendrites at the centre of the track. The continuous surface at the outmost surface of Fig. 1a was analysed by XRD and the XRD spectrum, as shown in Fig. 3, indicates that it is mainly composed of TiN compounds and no diffraction peaks of the phase of NiTi were detected. The analysis of the microstructure in the LGN surface of NiTi has been reported in our previous paper [11]. Fig. 4 exhibits the results of EDX spectra of the as received NiTi alloy (Fig. 4a) and the matrix phase of the LGN layer (Fig. 4b). Although EDX composition analysis has a 3% to 5% error, the spectra in Fig. 4 nevertheless provide an indication that the compositions of both specimens are very much the same. This is crucial as the composition affects the martensitic-austenitic phase transition temperatures.

H.C. Man, N.Q. Zhao / Surface & Coatings Technology 200 (2006) 5598 – 5605

Lin(Counts)

NiTi TiN

a b 40

50

60

70

80

5601

Fig. 5c shows the XPS spectrum of the LGN layer after the surface, including the outmost TiN layer, was ground away by about 50 Am. It exhibits the composition of the TiN(dendrites)/NiTi(matrix) composite layer which consists of 16 at.% Ni, 28.4 at.% Ti, 10.8 at.% N and some C,O as well. A narrow XPS spectrum analysis shows that the Ti element is associated with TiN, NiTi and a small amount of TiO2 whereas the Ni element is only associated with the NiTi phase only. This indicates that the amount of Ni element inside the LGN layer remains the same as the nominal NiTi alloy. 3.3. DSC analysis

90

2θ Fig. 3. The XRD spectra of the (a) as-received NiTi alloy and (b) LGN NiTi surface.

The laser processing parameters used in this work was very different from those used by Kloosterman and De Hosson. [18,19] in their work of laser nitrided titanium. A carbon dioxide laser focus beam of 0.55 mm diameter and 1200 W was used [18,19] by them and this meant that a power density of 5053 W/mm2 was irradiated onto the titanium surface. To avoid the formation of plasma at slow speed of processing, a high speed (25 – 200 mm/s) was chosen by those authors and this would result in high solidification rate and subsequently strong dendritic orientation. However, in this work, a defocus beam of 2 mm diameter was used and only 159 W/mm2 power intensity was applied. With this low power intensity, a slow processing speed was required to produce melting. The slow speed used in this work (5 mm/s) resulted in a slow solidification rate and less anisotropy in the TiN dendrites. 3.2. XPS analysis Immediately after the LGN process, it was found that the specimen surfaces have thin deposit film of brown-yellow color. This film could be cleaned away easily with a HF + HNO3 water solution reagent and a TiN layer with bright golden color would then appear. The chemical composition of the brown-yellow film was analysed by XPS. Fig. 5a, b and c show the XPS results of the LGN specimen surface just after the LGN process, the LGN specimen surface after the acid cleaning process, and the dendritic layer after 40 Am from the outmost TiN surface was ground off, respectively. It is obvious that in Fig. 5a, the Ni concentration (39.6 at.%) is much higher than those in Fig. 5b and c and according to the subsequent analysis of a narrow peak scanning, the Ni element exists as pure Ni or NiO. No Ni element was detected on the cleaned surface of the LGN specimen (Fig. 5b). This indicated that the TiN layer at the outmost surface of the LGN layer was free of Ni.

The DSC curves for the as-received hot rolled NiTi (Specimen A) are shown in Fig. 6a. Two peaks were observed on the first heating curve, and these endothermic peaks have been labeled as H1 and H2. The exothermic peak (C) observed on the cooling curve is interpreted as the reverse transformations of austenite (B2) to martensite (B19V) (A Y M). It is noticed that on the second heating curve, only one endothermic peak (M Y A), labeled as H, was exhibited instead of two peaks as appeared on the first heating curve. Fig. 6b shows the DSC curves of the NiTi alloy after the heat treatment (HT) process (Specimen B). The two endothermic peaks on the first and second heating curve coincided perfectly. The feature of the cooling curve was similar to that of Specimen A, but the transformation start temperature (Ms) shifted towards a higher temperature. Fig. 7a shows the DSC curves of the LGN NiTi specimen (Specimen C, non-heat treated). It can be seen that the heating and cooling curves have no significant difference as compared with that of Specimen A in Fig. 6a. The phase transformation temperatures for both specimens are almost the same. Fig. 7b shows the DSC curve of the LGN NiTi specimen after heat treatment process, Specimen D. It can be seen that upon heating, the H1 peak is weaker and the H2 peak is stronger than those of Specimen B. The transformation temperatures of the four specimens were summarized in Table 1. On cooling from the B2regime, the following features are noted: (1) the Ms, Mp and Mf temperatures shift towards higher temperatures after the heat treatment, (2) the martensitic transformation temperatures of non-LGN and LGN specimens (Specimen A and C, Specimen B and D) begin and end at very similar temperatures. For the heating process in which the B19V phase transforms to the B2 phase, the following points are noted: (1) for all four specimens, their As temperatures are close to each other, but the Af temperatures are different among the HT and non-HT specimen, for example A and B, C and D, respectively; (2) the phase transformation temperatures have no obvious difference after the LGN process, i.e. specimens A and C, specimens B and D, respectively.

5602

H.C. Man, N.Q. Zhao / Surface & Coatings Technology 200 (2006) 5598 – 5605

Fig. 4. EDX spectra of the as-received NiTi and the matrix phase of the LGN layer.

3.4. Discussions The shape memory behavior of NiTi alloys is affected by the alloy composition and the heat treatment process. A difference of 1 at.% in composition can cause a dramatic change in the behaviour of the SMA [20]. During laser gas nitriding, the reaction product, TiN, consumes a certain amount of Ti from the melt pool and appears in the form of dendrites upon solidification. At the same time, some Ni will be vaporized since the melting point of Ni (1453 -C) is lower than that of Ti (1720 -C). The condensed vapor of Ni in the form of NiO and Ni as appeared as the brown-yellow deposit was confirmed by XPS in Fig. 5a. The remaining Ni in the melt pool may form NiTi or a eutectic of NiTi and TiNi3. The remaining Ti in the melt pool would combine with Ni to from the NiTi phase only if the melt composition satisfies the condition of

forming the NiTi phase [21]. The results of the present work indicated that no eutectic in the LGN layer was observed. The results also indicated that it is possible to retain the NiTi phase in the LGN layer after the LGN treatment. This is a crucial condition for keeping the shape memory effect of NiTi alloy after the LGN process. It is possible to investigate the change of Ni content by means of DSC analysis because the phase transition temperatures scale with the Ni content. In general, solution annealed and water quenched Ni-rich NiTi specimens show one step martensitic transformations on cooling from B2 to B19V and one step transformation on heating from B19Vto B2 [14]. Aging of this type of materials results in the occurrence of metastable particles of type Ni4Ti3 [22,23]. The transition favors an intermediate phase, Rphase, because of its smaller transformation strain that resulted as compared to B19V[24]. A typical DSC chart of

H.C. Man, N.Q. Zhao / Surface & Coatings Technology 200 (2006) 5598 – 5605

5603

Fig. 5. XPS surface analysis of LGN NiTi specimen; (a) brown-yellow deposit film on surface immediately after the LGN process; (b) LGN specimen surface after acid cleaning; (c) the TiN/NiTi composite layer after 50 Am of the surface was removed.

aged Ni-rich NiTi SMAs would have two distinct peaks on cooling: B2 to R-phase (first distinct DSC peak), followed by R-phase to B19V(second distinct peak). Recently it was shown that there are cases where both R-phase and B19V coexist [25]. In our present work, only one peak was found on cooling, as shown in Figs. 6 and 7. The phase transformation temperature of HT specimen, B, however, increased and the peak was obviously narrower than that of the non-HT specimen, A. It can be assumed that the one peak in specimen A undoubtedly indicates a one step transformation from B2 to B19V. However, the narrow single exothermal peak of specimen B represents the transformation involving the intermediate R-phase. The transformation from B2 to R-phase requires less undercooling, so the Ms start temperature of the transformation

increases as compared with specimen A, the non-HT NiTi alloy. The DSC charts of specimen A and specimen B show two peaks on first time heating (Fig. 6a and b). But we have reasons to believe that the two peaks in both specimens were the results of different transformation mechanisms since their second heating curves have different features. As shown in Fig. 6a, Specimen A only has one peak in the reheating curve instead of two peaks in the first heating curve. On the other hand, as shown in Fig. 6b, the second heating curve of Specimen B shows exactly the same feature as that of the first time heating curve. The two peaks of the first time heating curve of Specimen A are believed to be the result of the residual stress during heat rolling. Since the stress levels are different between the grain boundary and

5604

H.C. Man, N.Q. Zhao / Surface & Coatings Technology 200 (2006) 5598 – 5605

21.5

28

H

H1 20.5 20.0

Cooling

19.5 -20

0

20

40

Temperature,

60

80

Heating Reheating

24 H2 22 20 18

C -40

H1

26

Heat Flow, mW

2nd heating

21.0

Heat Flow, mW

1st Heating H2

100

C -40

-20

0C

0

60

40

20

Temperature,

80

100

0C

b) Specimen B

a) Specimen A

Fig. 6. The DSC curves of NiTi as-received and HT specimen.

the grain after hot rolling, the free energy of transformation from B19Vto B2 will be different too, which results in two peaks in the heating curve. But upon reheating, the residual stress has already been relieved during the first heating and so only one peak has resulted. In Fig. 6b, the DSC heating curves of Specimen B showed that the two peaks of the first heating curve coincided with the two peaks of the second heating cycles. It is believed that the first peak has resulted from the transformation of B19Vto R-phase and the second peak from the transformation of R-phase to B2. According to KhalilAllafi et al. [16], they suggested that the two-step back transformations on heating from B19V are due to a first transformation peak associated with the back transformation in the Ni4Ti3 precipitate free regions and a second transformation peak in the remainder of the microstructure. The results in our experiment seem to correspond with KhalilAllafi’s suggestion. After the LGN treatment, the DSC charts of the specimen showed no considerable difference as compared with the non-LGN specimens (Figs. 6a and 7a, Figs. 6b and 7b). It is clear that DSC charts of the non-heat treated specimens have similar features before and after LGN treating. This

indicates that the LGN process has no effect on the matrix composition. The DSC charts of both non-LGN and LGN aged NiTi (Figs. 6b and 7b) showed two peaks. It can be noted that the width of the peaks after LGN are larger than that of the non-LGN specimen. The first peak in Fig. 7b is lower than the first peak in Fig. 6b whereas it was the opposite for their second peaks. As pointed out by Khalil-Allafi et al. [16], the peaks of the DSC charts is influenced by the Ni4Ti3 precipitates in Ni-rich NiTi alloys. The precipitation processes in solution annealed NiTi alloys during aging are strongly affected by the presence of external and internal stresses in the nucleation stage. Stress free aging results in heterogeneous microstructures where precipitates are mainly found around grain boundaries and defects (such as oxide particles) while grain interiors exhibit precipitate-free zones. But stressassisted aging resulted in a homogeneous distribution of precipitates throughout the microstructure resulting from the dislocations induced by the stress [16]. Grain boundaries are well known as representing locations for heterogeneous nucleation in solid state precipitation processed for a number of reasons [26]. Moreover, grain boundaries energetically favor nucleation by decreasing the interfacial

21.0 H2

20.6

Heat Flow , mV

Heat Flow, mW

21.5

H1

20.8 20.4 20.2 20.0 19.8

H1

21.0

H2

Heating

20.5. 20.0 Cooling 19.5

19.6 19.4

C

C -40

-20

0

19.0 20

40

Temperature,

60

80

100

-40

-20

0C

a) Specimen C Fig. 7. The DSC curves of LGN NiTi.

0

20

40

Temperature,

60 0C

b) Specimen D

80

100

H.C. Man, N.Q. Zhao / Surface & Coatings Technology 200 (2006) 5598 – 5605

5605

Table 1 Transformation temperatures of NiTi specimens Specimen no.

Treatment processes

Heating (-C)a H1s

H1p

H2f

H2s

H2p

H2f

Ms

Mp

A B C D

As-received NiTi HT NiTi LGN NiTi HT LGN NiTi

19.6 19.7 21.2 18.6

31.3 22.7 31.5 22.5

38.0 25.5 38.9 20.7

48.5 28.0 47.6 23.1

65.0 30.2 60.0 30.5

87.4 33.7 73.2 40.3

18.1 33.6 15.9 38.4

5.3 24.5 5.3 24.6

a b

Cooling (-C)b Mf 18.4 18.3 12.5 10.2

For the heating curves, the H1s and H1f temperatures correspond to Peak H1. The H2s and H2f temperatures correspond to Peak H2 transformation. For the cooling curves, the Ms and Mf temperatures correspond to the A – M transformation.

energy between the precipitate and the parent phase [26]. This implies that the interface may therefore be important in precipitation. After the LGN process, the distribution of the TiN dendrites in the matrix of NiTi produces a great deal of interfaces between the dendritic TiN and NiTi matrix. The interfaces thus become the nuclear position of Ni4Ti3. It is suggested here that the change in amplitude of the two peaks in Fig. 7b is due to the uniform distribution of TiN dendrites in the NiTi alloy matrix that makes the Ni4Ti3 precipitate not only near the grain boundaries and defects region but also inside the grains as well. Furthermore, there may well be a higher Ni-concentration at the interface of TiN dendrites/ NiTi that causes the precipitation process to occur throughout the nitrided layer.

4. Conclusions A TiN/NiTi two-phase surface alloyed layer on NiTi alloys was obtained by laser gas nitriding process. The alloyed layer, about 400 Am thick, consists of a continuous surface coating of about 2 Am thick TiN at the outmost surface and a layer composed of TiN dendrites distributed in the re-solidified NiTi matrix. XPS and EDX analysis results showed that the matrix of LGN NiTi layer retains the same composition as the asreceived alloy. DSC measurements show that laser gas nitriding process does not significantly affect the transformation temperatures of the shape memory alloy. The TiN dendritic structure in the LGN layer affects the precipitation of Ni4Ti3 coherent phase which in turn affects the phase transformation characteristics of solution annealed, quenched and aged NiTi alloys.

Acknowledgement The work described in this paper has been supported by the Research Grant Council of Hong Kong, China (Project No. 5270/03E). Support from the Hong Kong Polytechnic University is also acknowledged.

References [1] T. Hanawa, in: J.E. Davies (Ed.), The Bone – Biomaterial Interface, University of Toronto Press, Toronto, 1991, p. 49. [2] D. Starosvetsdy, I. Gotman, Surf. Coat. Technol. 148 (2001) 268. [3] Y. Liu, X.J. Yang, M. Chen, S. Zhu, Z. Cui, J. Mater. Sci. Technol. 18 (2002) 534. [4] C. Trepanier, M. Tabrizian, L. Yahia, L. Bilodeau, D. Piron, J. Biomed. Mater. Res. 43 (1998) 433. [5] H.C. Man, Z.D. Cui, T.M. Yue, Scr. Mater. 45 (2001) 1447. [6] Y. Fu, X.F. Wu, Y. Wang, B. Li, S.Z. Yang, Appl. Surf. Sci. 157 (2000) 167. [7] K. Endo, R. Sachdeva, Y. Araki, H. Ohno, Dent. Mater. J. 13 (1994) 228. [8] S.K. Wu, H.C. Lin, C.L. Chu, Surf. Coat. Technol. 92 (1997) 206. [9] D. Starosvetsdy, I. Gotman, Biomaterials 22 (2001) 1853. [10] H.C. Man, Z.D. Cui, T.M. Yue, J. Laser Appl. 14 (4) (2002) 242. [11] Z.D. Cui, H.C. Man, X.J. Yang, Appl. Surf. Sci. 208 (2003) 388. [12] Z.D. Cui, H.C. Man, F.T. Cheng, Surf. Coat. Technol. 162 (2003) 147. [13] C.M. Wayman, T.W. Duerig, in: T.W. Duerig, K.N. Melton, D. Stockel, C.M. Wayman (Eds.), Engineering Aspects of Shape Memory Alloys, Butterworth-Heinemann, London, 1990 (3 pp.). [14] T. Saburi, in: K. Otsuka, C.M. Wayman (Eds.), Shape Memory Materials, Cambridge University Press, Cambridge, 1998, p. 49. [15] W. Tang, B. Sundmann, R. Sandstro¨m, C. Quiu, Acta Mater. 47 (1999) 3457. [16] Jafar Khalil-Allafi, Antonin Dlouhy, Gunther Eggeler, Acta Mater. 50 (2002) 4255. [17] K. Otsuka, in: T.W. Duerig, K.N. Melton, D. Stockel, C.M. Wayman (Eds.), Engineering Aspects of Shape Memory Alloys, ButterworthHeinemann, London, 1990, p. 36. [18] A.B. Kloosterman, J.Th.M. De Hosson, J. Mater. Sci. 32 (1997) 6201. [19] A.B. Kloosterman, J.Th.M. De Hosson, Scr. Metall. Mater. 33 (4) (1995) 567. [20] W. Tang, B. Sundman, R. Sandstro¨m, C. Qiu, Acta Mater. 47 (1999) 3457. [21] T.B. Massalski, J.L. Murray, L.H. Bennett, H. Baker, Binary Alloy Phase Diagrams, American Society for Metals, Metals Park, Ohio, 1986 (1768 pp.). [22] D.Y. Li, L.Q. Chen, Acta Mater. 46 (1998) 639. [23] T. Tadaki, Y. Nakata, K. Shimizu, K. Otsuka, Trans. JIM 27 (1986) 731. [24] H. Sitepu, W.W. Schmahl, J. Khalil-Allafi, G. Eggeler, A. Dlouhy, D.M. To¨bbens, M. Tovar, Scr. Mater. 46 (2002) 543. [25] X. Ren, N. Miura, J. Zhang, K. Otsuka, K. Tanaka, M. Koiwa, S. Suzuki, Y.I. Chumlyakov, Mater. Sci. Eng., A Struct. Mater.: Prop. Microstruct. Process. 312 (2001) 196. [26] D.A. Porter, K.E. Easterling, 2nd edR, Chapman and Hall, London, 1997 (reprinted).