Shape memory characteristics of as-spun and annealed Ti51Ni49 crystalline ribbons

Intermetallics 18 (2010) 965–971

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Shape memory characteristics of as-spun and annealed Ti51Ni49 crystalline ribbons S.H. Chang a, S.K. Wu b, *, L.M. Wu b a b

Department of Chemical and Materials Engineering, National I-Lan University, I-Lan 260, Taiwan Department of Materials Science and Engineering, National Taiwan University, Taipei 106, Taiwan

a r t i c l e i n f o

a b s t r a c t

Article history: Received 24 November 2009 Received in revised form 22 December 2009 Accepted 14 January 2010 Available online 16 February 2010

The as-spun Ti51Ni49 ribbon is essentially crystalline and shows a two-stage B2 0 R and R 0 B190 martensitic transformation with a significant transformation enthalpy. The as-spun ribbon already has plate-like Guinier-Preston (GP) zones which are proposed to induce the R-phase in martensitic transformation. After being annealed at 200  C for 1 h, the size of the GP zones increases but their number decreases simultaneously. The maximal recoverable strain of the as-spun and 200  C annealed Ti51Ni49 ribbons reaches 6.2% and 6.4%, respectively. With the increase of the annealing temperature up to 500  C and 600  C, only Ti2Ni precipitates can be observed in the ribbons and the two-stage martensitic transformation tends to coalesce. Experimental results reveal that the ribbons with GP zones show a higher recoverable strain, a lower permanent strain and a lower Ms temperature than those with Ti2Ni precipitates. The reason why the GP zones are formed in as-spun Ti51Ni49 ribbon during melt-spinning procedure is also discussed. Ó 2010 Elsevier Ltd. All rights reserved.

Keywords: B. Martensitic transformations B. Shape-memory effects B. Precipitates C. Rapid solidification processing

1. Introduction TiNi-based alloys are known as the most important shape memory alloys (SMAs) because they exhibit thermolelastic martensitic transformation and show excellent shape memory effect (SME), superelasticity and damping capacity [1]. Recently, melt-spinning technique has been utilized to fabricate amorphous or crystalline Ti–Ni binary and Ti–Ni–Cu ternary SMA ribbons for various applications. Among them, Ti50Ni25Cu25 ribbon was widely studied because of its unique properties, such as small transformation hysteresis, low flow stress level in the martensite state, low sensitivity of the martensitic transformation start (Ms) temperature and easy fabrication for fully amorphous melt-spun ribbons [2–18]. These advantages make Ti50Ni25Cu25 ribbon a good candidate for applications, such as actuator and sensor that require short response time at thermal cycle. Unfortunately, amorphous Ti50Ni25Cu25 ribbon does not possess any martensitic transformation behavior and cannot show SME. Therefore, a proper thermal annealing procedure is required for practical applications.

* Corresponding author. Department of Materials Science and Engineering, National Taiwan University, 1, Roosevelt Rd., Sec. 4, Taipei 106, Taiwan. Tel.: þ886 2 2363 7846; fax: þ886 2 2363 4562. . E-mail address: [email protected] (S.K. Wu). 0966-9795/$ – see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2010.01.010

Several literatures have reported the annealing effect on the microstructures, textures, martensitic transformations, shape memory properties and crystallization kinetics of melt-spun Ti50Ni25Cu25 ribbon [13–16]. Nevertheless, it is very difficult to control the appropriate annealing conditions for Ti50Ni25Cu25 ribbon [17,18]. This is because no obvious martensitic transformation can be observed before Ti50Ni25Cu25 ribbon is sufficiently crystallized. On the other hand, the crystallized ribbon becomes very fragile and its SME is deteriorated significantly after prolonged annealing interval. On the contrary, as-spun Ti-Ni binary SMA ribbons manufactured by similar melt-spinning technique are usually crystalline and can be expected to behave better SME. Some literatures have investigated the textures, shape memory properties and the microstructures of as-spun Ti51Ni49, Ti50Ni50 and Ti49Ni51 ribbons [19–21]. However, the annealing effect on the Ti– Ni binary melt-spun ribbons has not been studied in detail yet. In this study, as-spun Ti51Ni49 ribbon was selected to investigate its microstructure and SME. The annealing effect on the precipitation formation and the shape memory behavior of Ti51Ni49 melt-spun ribbons were also systematically studied and discussed.

2. Experimental procedures Ti51Ni49 melt-spun ribbons were prepared by the Institute for Materials Research, Tohoku University, JAPAN, using a single-roller

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melt-spinning technique. Ti51Ni49 ingot was first prepared by conventional vacuum arc remelter to form a homogenized ingot. Then the ingot was cut into appropriate size and induction-melted in a quartz crucible filled with argon. The melted Ti51Ni49 alloy was subsequently ejected with pressurized argon gas onto a watercooled copper roller with a surface velocity of 42 m/s. The as-spun ribbons was 14 mm in thickness and 1.05 mm in width. The martensitic transformation temperatures and the transformation enthalpies of the as-spun and annealed Ti51Ni49 ribbons were determined by TA Q10 differential scanning calorimetry (DSC). The specimen weight used in DSC test was about 2–4 mg and the heat ing/cooling rate (T ) was set as 10  C/min. The shape memory behavior of the ribbons was investigated by thermal cycling tests under various constant stresses using TA 2980 dynamic mechanical analysis (DMA) with a film tension clamp and a liquid nitrogen cooling apparatus. The T used in the thermal cycling tests was set as 3  C/min and the dimension of DMA specimens was 10 mm in gauge length, 1.05 mm in width and 14 mm in thickness. Microstructural observation for Ti51Ni49 ribbon was performed with a Tescan 5136MM scanning electron microscope (SEM). The specimens for the transmission electron microscopy (TEM) were prepared by an argon ion miller and observed using a Philips-TECNAI G2 F20 FEGTEM microscope operated at 200 kV. 3. Results 3.1. Characteristics of the as-spun Ti51Ni49 ribbon 3.1.1. DSC and SEM results Fig. 1(a) shows the DSC curve of the as-spun Ti51Ni49 ribbon. As shown in Fig. 1(a), there are B2 2 R and R 2 B190 martensitic transformation peaks in cooling, but only a B190 / B2 peak in heating. Besides, both transformation enthalpies (DHs) of the transformations peaks in cooling and heating curves are close to 24 J/g, which is almost identical to those of Ti50Ni50 bulk SMA, typically about 25 J/g. This implies that the as-spun Ti51Ni49 ribbon used in this study is crystalline. Fig. 1(b) shows a top-view SEM micrograph of the free surface of the as-spun Ti51Ni49 ribbon. As shown in Fig. 1(b), the ribbon consists of spherical crystalline grains of a uniform grain size with a diameter of about 3 mm. Therefore, the multi-stage martensitic transformation induced by the inhomogeneous grain size distribution in the ribbon will certainly not appear in this study [22]. According to the DSC and SEM results, the as-spun Ti51Ni49 ribbon used in this study is essentially crystalline and adequate to conduct the following DMA experiments.

3.1.2. DMA and TEM results Fig. 2(a) shows the DMA strain-temperature curve of the asspun Ti51Ni49 ribbon measured under a constant stress of 238 MPa. As illustrated in Fig. 2(a), the martensitic transformation temperatures such as Ms (martensite start), Mf (martensite finish), As (austenite start), Af (austenite finish) temperatures as well as the recoverable strain (3a) and the permanent strain (3p) of the ribbon can be determined from the strain–temperature curve. Fig. 2(b) shows the DMA strain–temperature curves of the as-spun Ti51Ni49 ribbon measured under different constant stresses from 68 MPa to 238 MPa. In Fig. 2(b), it shows the determined 3a increases with the increase of the applied stress and reaches a maximum value of 6.2%. When the applied stress is lower than 136 MPa, the specimen can totally recover throughout the thermal cycle and no permanent strain 3p is retained. Besides, as shown in Fig. 2(b), the B2 2 R and R 2 B190 two-stage martensitic transformation only appears when the applied stress is 68 MPa. The Rs (R-phase start) and Rf (R-phase finish) temperatures can also be determined, as illustrated in Fig. 2(b). With the increase of the applied stress, however, only a single B2 2 B190 martensitic transformation can be found in the DMA strain–temperature cooling curve. This feature comes from the fact that the stress dependence of the Ms temperature is much larger than that of the Rs temperature, which leads to the coalescence of the two-stage martensitic transformation when the applied stress is increased [23,24]. Many literatures reported that the formation of R-phase in martensitic transformation of TiNi-based bulk SMAs corresponds to the aging of Ni-rich TiNi SMAs, the substitution of Fe or Al for Ti or Ni in TiNi SMAs, and the introduction of structural defects or dislocations into TiNi-based SMAs by means of heavy cold-working, thermal cycling or exploring under high energy radiation [25–27]. The multiple martensitic transformation appearing in Ti-rich TiNi SMAs is uncommon but has been observed in the annealed specimens of sputtering-deposited TiNi/TiNiCu thin films [28–31] and TiNi gas-atomized powders [32]. According to the reported studies, the R-phase appearing in the martensitic transformation of annealed Ti-rich TiNi thin films is induced either by the coherent stress field around the plate-like Guinier-Preston (GP) zones or by the semicoherent stress field around the spherical Ti2Ni precipitates [28–32]. This feature is similar to the formation of R-phase in the martensitic transformation of aged Ni-rich TiNi bulk SMAs in which the stress field is developed by Ti3Ni4 precipitates [33]. Recently, Khantachawana et al. [21] and Nam et al. [34] reported that the Rphase can also be observed in the martensitic transformation of Tirich Ti–Ni binary and Ti–Ni–Cu ternary melt-spun ribbons, however,

Fig. 1. (a) DSC curve and (b) SEM image for the as-spun Ti51Ni49 ribbon.

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Fig. 2. DMA strain–temperature curve for the as-spun Ti51Ni49 ribbon measured (a) at a constant applied stress of 238 MPa and (b) at different applied stresses from 68 MPa to 238 MPa.

only Ti2Ni precipitates rather than GP zones are observed in these ribbons. Fig. 3(a) and (b) show the TEM bright field (BF) image and the selected area diffraction pattern (SADP), respectively, of the asspun Ti51Ni49 ribbon used in this study. As shown in Fig. 3(a), the plate-like GP zones with the length of about 10 nm can be observed, as the single arrows indicated. This manifests clearly that the GP zones form in as-spun Ti51Ni49 ribbons and they can induce the R-phase, as 1/3 <011>* diffraction spots shown in Fig. 3(b). We suggest that either GP zones or Ti2Ni precipitates are formed in the as-spun Ti–Ni/Ti–Ni–Cu ribbons may correspond to alloy’s composition and different cooling rate during the melt-spinning process. However, only GP zones rather than Ti2Ni precipitates are observed in the as-spun Ti51Ni49 ribbons of this study. 3.2. Characteristics of the Ti51Ni49 ribbon annealed at different temperatures 3.2.1. DSC and DMA results In order to investigate the annealing effect on the shape memory behavior of Ti51Ni49 ribbons, the as-spun ribbons were evacuated annealing at different temperatures for 1 h. Fig. 4(a)–(c) show the DSC cooling and heating curves of the Ti51Ni49 ribbons annealed at 200  C, 500  C and 600  C, respectively, for 1 h. The DSC curve of the ribbon annealed at 200  C, as shown in Fig. 4(a), is similar to that of the as-spun ribbon shown in Fig. 1(a). However, with the increase of the annealing temperature up to 500  C and

600  C, as shown in Fig. 4(b) and (c), respectively, the B2 2 R and R 2 B190 two-stage martensitic transformation peaks start to merge. Meanwhile, with the increase of the annealing temperature, the DH value only increases gradually from z24 J/g to z25 J/g, but the B2 2 R transformation peak temperature raises significantly from z47  C to z 55  C. The annealing temperature effect is inconspicuous on the DH value because the as-spun Ti51Ni49 ribbon is already crystalline. Besides, the increase of DH value can be attributed to the increase of the martenstic transformation peak temperature since the entropy change is usually maintaining constant for martensitic transformation [35]. Fig. 5(a)–(c) show the DMA strain–temperature curves of Ti51Ni49 ribbons annealed at 200  C, 500  C and 600  C, respectively, for 1 h. Each annealed ribbon in Fig. 5 exhibits a well-defined strain–temperature curve which reveals that the ribbons still possess good SME after annealing. In addition, as shown in Fig. 5, the Rs and Rf temperatures can only be obtained in the strain–temperature curve (under 68 MPa stress) of the ribbon annealed at 200  C. This corresponds to the coalescence of the B2 2 R and R 2 B190 martensitic transformation peaks in the DSC curves for the ribbons annealed at 500  C and 600  C, as demonstrated in Fig. 4. Fig. 6(a) and (b) plot the evolution of the 3a and 3p, respectively, for the Ti51Ni49 ribbons measured from the DMA results of Fig. 2(b) and Fig. 5. As shown in Fig. 6(a), the 3a values of both as-spun and annealed Ti51Ni49 ribbons increase with the increase of the applied stress. Under a constant applied stress of

Fig. 3. TEM (a) BF image and (b) corresponding SADP for the as-spun Ti51Ni49 ribbon.

1.0

0.5

cooling

B2 R o B19' R 47.1 C o 21.5 C 24.5J/g ΔH

0.0 24.7J/g

heating

-0.5

-1.0 -200 -150 -100

exo up

0

50 o

Temperature( C)

100

1.0

0.5

R

B19'

150

200

c

B2 o 55.6 C

24.6J/g Δ H 24.7J/g

0.0 heating

o

-1.0 -200 -150 -100

exo up

83.5 C B2 B19'

-50

0

50

1.0 B2

R

B19'

o

55.7 C

cooling

-0.5

o

58.5 C B2 B19'

-50

b

Heat Flow(W/g)

a

Heat Flow(W/g)

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Heat Flow(W/g)

968

100

0.5

o

39.3 C

cooling 25.3J/g Δ H 25.3J/g

0.0 heating

-0.5

150

200

-1.0 -200 -150 -100

exo up

o

Temperature( C)

o

80.0 C B2 B19'

-50

0

50

100

150

200

o

Temperature( C)

Fig. 4. DSC curves for the Ti51Ni49 ribbons annealed at (a) 200  C, (b) 500  C and (c) 600  C for 1 h.

238 MPa, the maximum 3a value of the as-spun Ti51Ni49 ribbon and the ribbon annealed at 200  C can reach 6.2% and 6.4%, respectively. However, with further increasing annealing temperature to above 500  C, the maximum 3a value measured under the same applied stress decreases significantly to below 4.5%. In Fig. 6(b), the 3p value for each specimen also increases with the increase of the applied stress, in which the ribbons annealed at 500  C and 600  C both show a maximum 3p value of about 0.3% when the applied stress is 238 MPa. According to the 3a and 3p values shown in Fig. 6, the Ti51Ni49 ribbons can be divided into two groups. One is the as-spun Ti51Ni49 ribbon and the ribbon annealed at 200  C, and the other is the ribbons annealed at 500  C and 600  C, as presented by the empty and solid symbols in Fig. 6, respectively. Under the same applied

stress, the former group always exhibits a larger 3a value and a lower 3p value than the latter one. 3.2.2. TEM results Fig. 7(a) and (b) show the TEM BF image and the corresponding SADP, respectively, of Ti51Ni49 ribbon annealed at 200  C for 1 h. As shown in Fig. 7(a), the GP zones which are existent in the as-spun ribbon still can be observed. The GP zones are aligned along <010>B2 directions of the B2 parent phase and distributed about 5–20 nm apart. Compared the GP zones of Fig. 7(a) with those of Fig. 3(a), the maximum length of the GP zones increases from z10 nm to z20 nm after 200  C annealing but the number of the GP zones decreases. Fig. 7(c) and (d) show the BF image and the corresponding SADP of Ti51Ni49 ribbon annealed at 600  C for 1 h. As shown in Fig. 7(c), the

Fig. 5. DMA strain-temperature curve measured at different applied stresses from 68 MPa to 238 MPa for Ti51Ni49 ribbons annealed at (a) 200  C, (b) 500  C and (c) 600  C for 1 h.

8

6

b Permanent Strain, p (%)

a Recoverable Strain, a (%)

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4

As recieved o Annealed at 200 C for 1h o Annealed at 500 C for 1h o Annealed at 600 C for 1h

2

0

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0.0 0

50

100

150

200

250

300

Applied Stress (MPa)

0

50

100

150

200

250

300

Applied Stress (MPa)

Fig. 6. Evolution of the (a) recoverable strain (3a) and (b) permanent strain (3p) as a function of the applied stress for the Ti51Ni49 ribbons.

GP zones no longer exist, instead, only precipitates with Moire´ fringes can be observed. The sputtering-deposited Ti52.4Ni38.3Cu9.3 thin films annealed at 600  C, 650  C and 750  C for 600 s are also found the precipitates with Moire´ fringes [31]. As can be seen in Fig. 7(d), the Moire´ fringes are parallel to the {011}B2 planes of the B2 parent phase and the extra spots of the SADP can be identified as

Ti2Ni precipitate [36]. The BF images and the SADP of the ribbon annealed at 500  C are not presented here since they are similar to those of the ribbon annealed at 600  C but having a smaller size of Ti2Ni precipitates. Only Ti2Ni precipitates with Moire´ fringes rather than GP zones can be observed in 500  C and 600  C annealed ribbons. Compared with the DMA strain-temperature curves shown

Fig. 7. TEM (a) BF image and (b) corresponding SADP for the Ti51Ni49 ribbon annealed at 200  C for 1 h. TEM (c) BF image and (d) corresponding SAPD for the Ti51Ni49 ribbon annealed at 600  C for 1 h.

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Table 1 Microstructure of precipitates, recoverable strain (3a), permanent strain (3p) and Ms temperature for the Ti51Ni49 ribbons and the Ti52.1Ni47.9 thin films [30,37]. Composition (at. %)

Heat treatment

Precipitates

Recoverable strain, 3a (%)

Permanent strain, 3p (%)

Ms temperature ( C)

Ti51Ni49 ribbon (238 MPa)

As received 200  C  1 h 500  C  1 h 600  C  1 h

GP zones GP zones Ti2Ni Ti2Ni

6.2 6.4 4.5 4.1

0.1 0.2 0.3 0.3

55.0 52.8 61.1 61.0

Ti52.1Ni47.9 thin film [30,37] (400 MPa)

450  C  1 h 500  C  1 h 550  C  1 h

GP zones GP zones þ Ti2Ni Ti2Ni

5.5 5.1 3.9

0.0 0.1 0.4

– – –

in Fig. 6, we conclude that the Ti51Ni49 ribbons with GP zones possess a larger 3a value and a lower 3p value than those with Ti2Ni precipitates. 4. Discussion 4.1. The formation of the GP zones in as-spun and annealed Ti51Ni49 ribbons From Section 3.1, the as-spun Ti51Ni49 ribbon used in this study is regarded as a fully crystalline one with transformation enthalpy being almost identical to that of Ti50Ni50 bulk SMA. The TEM observation shown in Fig. 3 demonstrates that the GP zones of the as-spun Ti51Ni49 ribbons are aligned along <001>B2* directions and the reported studies indicate that there is coherent stress around them which can induce the R-phase during the martensitic transformation [28,30]. The results of Section 3.1 clearly indicate that, for Ti51Ni49 ribbons melt-spun by the method illustrated in Section 2, the GP zones are already formed during the rapid solidification process (RSP). This characteristic is different from that observed in Ti-rich Ti–Ni SMAs thin films in which the as-sputtered thin films are always amorphous [28,29,31,37,38]. To form the GP zones in Ti-rich Ti–Ni thin films, annealing at temperature above 450  C for suitable time is needed. This means that, in the as-sputtered condition, supersaturated Ti atoms are dissolved in the amorphous thin films due to the cooling rate of the sputtering thin films can reach about 108  C/sec [39]. Therefore, the GP zones can only be formed through the suitable annealing conditions, a behavior similar to those GP zones precipitated in age-hardening alloys, such as Al–Cu aluminum alloys, Cu–Be copper alloys, etc. [40,41]. In contrast, the cooling rate of the ribbons fabricated by RSP is in the range of 104 w 106  C/sec which is much lower than that of the sputtering-deposited thin films [4]. Besides, the as-spun ribbons are fabricated by the solidification of the liquid phase, but the as-sputtered thin films are deposited from the sputtered atoms which are bombarded by argon plasma [42]. Because the rapid solidification, the as-spun ribbons can form a supersaturated solid solution first when its temperature decreases to below the liquidus. For the aging of the supersaturated Ti-rich Ti– Ni solid solution, the GP zones can be regarded as the transition phase and the Ti2Ni precipitates are the equilibrium phase. The activation energy barrier to the formation of the GP zones is smaller than that of the Ti2Ni precipitates. Therefore, it is possible that the GP zones can be formed in the as-spun ribbons ranged from the liquidus to the room temperature. The reported study indicated that, for Ti36 wt% W alloy, the solidified ribbon contains the omega (u) phase [43]. This manifests the possibility of the formation of GP zones in the as-spun Ti51Ni49 ribbons. 4.2. Comparison shape memory characteristics of Ti51Ni49 ribbons with those of annealed Ti52.1Ni47.9 thin films Table 1 summarizes the precipitates microstructure, Ms temperature, 3a, and 3p values measured under a constant stress of

238 MPa for the as-spun Ti51Ni49 ribbon and the ribbons annealed at 200  C, 500  C and 600  C for 1 h. The results of the annealed Ti52.1Ni47.9 thin films measured under 400 MPa are also listed in Table 1 for comparison [30,37]. Table 1 shows that Ti51Ni49 ribbons and Ti52.1Ni47.9 thin films with GP zones all exhibit a large 3a value above 5.5%. When the ribbons and the thin films are annealed at temperatures above 500  C, only Ti2Ni precipitates can be observed in the specimens and show a lower 3a value below 4.5%. Similar results have also been reported by Kajiwama et al. [28] and Tomozawa et al. [31] that the annealed Ti-rich TiNi and TiNiCu thin films with GP zones have a larger 3a than those with Ti2Ni precipitates. This characteristic can be explained by the fact that the twinning in martensite phase can pass through the coherent stress field around the GP zones more easily than the semicoherenet interfaces around the Ti2Ni precipitates [30]. Table 1 also depicts that both Ti51Ni49 ribbons and Ti52.1Ni47.9 thin films with GP zones exhibit a smaller 3p value than those with Ti2Ni precipitates. This is because the slip dislocations can not cut through the coherent stress field around the GP zones and it causes a smaller 3p value after thermal cycling [28,30]. In addition, although the as-spun Ti51Ni49 ribbon and the ribbon annealed at 200  C for 1h both contain GP zones rather than Ti2Ni precipitates, the annealed ribbon always exhibits a larger 3p value than the asspun one. This is associated with the increasing size and the decreasing number of the GP zones in the annealing ribbon in which the strength of the annealing ribbon is less than that of the as-spun one. Moreover, it is also interesting to note that the 3a values of the ribbons with GP zones are larger than those of the thin films with GP zones, even though the applied stress on the ribbons (238 MPa) is smaller than that on the thin films (400 MPa). This may come from the fact that the surface energy which impedes the martensitic transformation becomes more dominating in the thin film and requires more energy to achieve a considerable 3a value since the thickness of the thin film (about 6–7 mm) is smaller than that of the melt-spun ribbon (about 14 mm). The other reason is the roughness of the free surface of the ribbon, as shown in Fig. 1(b), which is much higher than that of the thin films. This will cause the stress-concentration at the free surface of the ribbon tested by DMA tensile clamp for determining the strain–temperature curve. 5. Conclusions The as-spun Ti51Ni49 ribbon is essentially crystalline and shows a two-stage B2 2 R and R 2 B190 forward martensitic transformation and a one-stage B190 2 B2 reverse martensitic transformation in DSC curves with significant transformation enthalpies of z24 J/g. With the increase of the annealing temperature from 200  C to 600  C, the B2 2 R and R 2 B190 martensitic transformation peaks in the DSC curve gradually merge into a single peak with DH values around 25 J/g. The ribbons annealed at 500  C and 600  C exhibit higher martenstic transformation peak temperatures, which are attributed to their higher transformation enthalpies in DSC measurements. For Ti51Ni49 ribbons melt-spun by

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the method illustrated in Section 2, the GP zones are already formed in the as-spun ribbon. During the rapid solidification, the melt-spun ribbons are first situated at a supersaturated solid solution state when they are cooled to a temperature just below the liquidus. Meanwhile, with further cooling to room temperature, the melt-spun ribbon under the supersaturated solid solution state precipitates the GP zones, instead of the Ti2Ni precipitates, since the formation of the GP zones possesses a lower activation energy barrier than that of the Ti2Ni precipitates. After annealing at 200  C for 1 h, the size of the GP zones in the ribbon increases from about 10 nm to about 20 nm while the number of the GP zones decreases simultaneously. With further increasing the annealing temperature up to 500  C and 600  C, the GP zones in the ribbons diminish and only Ti2Ni precipitates with Moire´ fringes linger. The ribbons with GP zones exhibit a larger recoverable strain and a lower permanent strain than those with Ti2Ni precipitates due to the different coherency of the stress field around them. The Ti51Ni49 ribbon annealed at 200  C for 1 h exhibits a recoverable strain 3a of 6.4% which is higher than that of Ti52,1Ni47,9 thin film annealed at 450  C for 1 h, say 5.5%. Acknowledgements The authors gratefully acknowledge the financial support from National Science Council (NSC), Taiwan, Republic of China, under the grants NSC-97-2221-E002-035-MY3 and NSC-97-2221-E-197003. We also sincerely acknowledge the assistance of Professor H. Kimura, Institute for Materials Research, Tohoku University, Japan, in ribbons preparation. References [1] Wayman CM, Duerig TW. In: Duerig TW, Melton KN, Sto¨ckel D, Wayman CM, editors. Engineering aspects of shape memory alloys. London: ButterworthHeinemann Press; 1990. p. 3–20. [2] Xie ZL, Van Humbeeck J, Liu Y, Delaey L. TEM study of Ti50Ni25Cu25 melt spun ribbons. Scr Mater 1997;37:363–71. [3] Ro¨sner H, Shelyakov AV, Glezer AM, Feit K, Schlobmacher P. A study of an amorphous–crystalline structured Ti–25Ni–25Cu (at.%) shape memory alloy. Mater Sci Eng A 1999;273–275:733–7. [4] Ro¨sner H, Schlobmacher P, Shelyakov AV, Glezer AM. Formation of TiCu platelike precipitates in Ti50Ni25Cu25 shape memory alloys. Scr Mater 2000;43:871–6. [5] Satto C, Ledda A, Potapov P, Janssens JF, Schryvers D. Phase transformations and precipitation in amorphous Ti50Ni25Cu25 ribbons. Intermetallics 2001;9:395–401. [6] Ro¨sner H, Schlobmacher P, Shelyakov AV, Glezer AM. The influence of coherent TiCu plate-like precipitates on the thermoelastic martensitic transformation in melt-spun Ti50Ni25Cu25 shape memory alloys. Acta Mater 2001;49:1541–8. [7] Santamarta R, Schryvers D. Microstructure of a partially crystallised Ti50Ni25Cu25 melt-spun ribbon. Mater Trans 2003;44:1760–7. [8] Liu Y. Mechanical and thermomechanical properties of a Ti50Ni25Cu25 melt spun ribbon. Mater Sci Eng A 2003;354:286–91. [9] Santamarta R, Cesari E, Pons J, Goryczka T. Shape memory properties of Ni–Ti based melt-spun ribbons. Metall Mater Trans 2004;35A:761–70. [10] Cheng GP, Xie ZL. Some results on shape memory properties of Ti50Ni25Cu25 melt-spun ribbons. J Alloys Compd 2005;396:128–32. [11] Xie ZL, Chen GP, Liu Y. Microstructure and texture development in Ti50Ni25Cu25 melt-spun ribbon. Acta Mater 2007;55:361–9. [12] Tong Y, Liu Y, Xie Z, Zarinejad M. Effect of precipitation on the shape memory effect of Ti50Ni25Cu25 melt-spun ribbon. Acta Mater 2008;56:1721–32.

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[13] Chang SH, Wu SK, Kimura H. Crystallization kinetics of Ti50Ni25Cu25 melt-spun amorphous ribbons. Mater Trans 2006;47:2489–92. [14] Cheng GP, Xie ZL, Liu Y. Transformation characteristics of annealed Ti50Ni25Cu25 melt spun ribbon. J Alloys Compd 2006;415:182–7. [15] Cheng GP, Xie ZL, Liu Y. Texture and shape memory property of annealed Ti50Ni25Cu25 ribbons. Mater Sci Eng A 2006;425:268–71. [16] Liu Y, Xie ZL, Tong YX, Lim CW. Properties of rapidly annealed Ti50Ni25Cu25 melt-spun ribbon. J Alloys Compd 2006;416:188–93. [17] Chang SH, Wu SK, Kimura H. Annealing effects on the crystallization and shape memory effect of Ti50Ni25Cu25 melt-spun ribbons. Intermetallics 2007;15: 233–40. [18] Chang SH, Wu SK, Kimura H. Martensitic transformation of annealed Ti50Ni25Cu25 melt-spun ribbons. Mater Sci Eng A 2008;476:316–21. [19] Saravanan V, Khantachawana A, Miyazaki S. Texture analysis and properties of rapidly solidified Ti52Ni38Cu10 shape memory alloy. Mater Trans 2004;45:208–13. [20] Crone WC, Wu D, Perepezko JH. Pseudoelastic behavior of nickel–titanium melt-spun ribbon. Mater Sci Eng A 2004;375–377:1177–81. [21] Khantachawana A, Mizubayashi H, Miyazaki S. Texture and microstructure of Ti–Ni melt-spun shape memory alloy ribbons. Mater Trans 2004;45:214–8. [22] Lin KN, Wu SK. Martensitic transformation of grain-size mixed Ti51Ni49 meltspun ribbons. J Alloys Compd 2006;424:171–5. [23] Nam TH, Noh JP, Jung DW, Kim YW, Im HJ, Ahn JS, et al. The R phase transformation in Ti–49Ni (at.%) shape memory alloy ribbons fabricated by melt spinning. J Mater Sci Lett 2002;21:11–3. [24] Stachowiak GB, McCormick PG. Shape memory behaviour associated with the R and martensitic transformations in a NiTi alloy. Acta Metall 1988;36:291–7. [25] Fan G, Chen W, Yang S, Zhu J, Ren X, Otsuka K. Origin of abnormal multi-stage martensitic transformation behavior in aged Ni-rich Ti–Ni shape memory alloys. Acta Mater 2004;52:4351–62. [26] Fan G, Zhou Y, Otsuka K, Ren X. Ultrahigh damping in R-phase state of Ti–Ni– Fe alloy. Appl Phys Lett 2006;89:161902. [27] Liu Y, McCormick PG. Thermodynamic analysis of the martensitic transformation in NiTi – I. Effect of heat treatment on transformation behaviour. Acta Metall Mater 1994;42:2401–6. [28] Kajiwara S, Kikuchi T, Ogawa K, Matsunaga T, Miyazaki S. Strengthening of Ti–Ni shape-memory films by coherent subnanometric plate precipitates. Philos Mag Lett 1996;74:137–44. [29] Ishida A, Sato M, Kimura T, Sawaguchi T. Effects of composition and annealing on shape memory behavior of Ti-rich Ti–Ni thin films formed by sputtering. Mater Trans JIM 2001;42:1060–7. [30] Zhang JX, Sato M, Ishida A. On the Ti2Ni precipitates and Guinier–Preston zones in Ti-rich Ti–Ni thin films. Acta Mater 2003;51:3121–30. [31] Tomozawa M, Kim HY, Miyazaki S. Shape memory behavior and internal structure of Ti–Ni–Cu shape memory alloy thin films and their application for microactuators. Acta Mater 2009;57:441–52. [32] Yamamoto T, Kato H, Murakami Y, Kimura H, Inoue A. Martensitic transformation and microstructure of Ti-rich Ti–Ni as-atomized powders. Acta Mater 2008;56:5927–37. [33] Wu SK, Lin HC, Chou TS. A study of electrical resistivity, internal friction and shear modulus on an aged Ti49Ni51 alloy. Acta Metall Mater 1990;38:95–102. [34] Nam TH, Lee JH, Jung DW, Yu CA, Liu Y, Kim YW. Transformation behaviour of Ti–Ni and Ti–Ni–Cu alloy ribbons with nano Ti2Ni particles. Mater Sci Eng A 2007;449–451:1041–4. ˜ osa Ll, Yoshikawa M. Scr Metall Mater [35] Friend CM, Ortı´n J, Planes A, Man 1990;24:1641–5. [36] Wu LM, Wu SK. Philos Mag Lett in press, doi:10.1080/0950083100360773. [37] Ishida A, Martynov V. Sputter-deposited shape-memory alloy thin films: properties and applications. MRS Bull 2002;27:111–4. [38] Kajiwara S, Yamazaki K, Ogawa K, Kikuchi T, Miyazaki S. Main factors for obtaining good shape memory properties in sputter-deposited thin films of Ti–Ni based alloys. Trans Mater Res Soc Jpn 2001;26:183–8. [39] Luborsky FE, editor. Amorphous metallic alloys. London: Butterworth; 1983. p. 27. [40] Smith WF. Structure and properties of engineering alloys. Int’l ed.. New York: McGraw-Hill; 1993. pp. 176–232 [41] Porter DA, Easterling KE. Phase transformations in metals and alloys. UK: Van Nostrand Reinhold; 1981. p. 291. [42] Smith DL. Thin-film deposition. Int’l ed.. New York: McGraw-Hill; 1999. p. 431 [43] Anantharaman TR, Suryanarayana C. Rapid solidified metals. Brookfield, VT, USA; 1987. pp. 187–204.