Improvement of High-Temperature Shape-Memory Effect by Multi-ComponentAlloying for TiPd Alloys

Hiromichi Matsuda1,2, Hirotaka Sato1,2, Masayuki Shimojo1 and Yoko Yamabe-Mitarai2,+

1Shibaura Institute of Technology, Tokyo 135-8548, Japan2National Institute for Materials Science, Tsukuba 305-0047, Japan

The influence of multi-component alloying on the phase transformation and shape-memory effect was investigated to develop new high-temperature shape memory alloys (HT-SMAs). Four alloys®35Ti­20Pd­15Ni­15Pt­15Zr, 40Ti­20Pd­15Ni­15Pt­10Zr (high-entropy alloys,HEAs), 45Ti­20Pd­5Ni­25Pt­5Zr, and 45Ti­20Pd­10Ni­20Pt­5Zr (medium-entropy alloys, MEAs, at%)®were prepared. At roomtemperature, the B2 structure was stable in the HEAs, and no martensitic transformation (MT) was observed. However, in the MEAs, anMT from the B2 structure to a B19 structure was clearly observed. The MT temperature of the MEAs was comparable to or higher than those ofbinary and ternary TiPd alloys. The strengths of both the martensite and austenite phases in 45Ti­20Pd­5Ni­25Pt­5Zr were higher than those in45Ti­20Pd­10Ni­20Pt­5Zr and ternary TiPd alloys. We attempted to explain the high strength using the ¤ parameter, which indicates the latticedistortion for various atomic sizes, but a clear correlation was not observed, as there were no significant differences in the ¤ parameter among thetested alloys. The shape recovery was investigated via a thermal cyclic test under an applied stress in the range of 15­200MPa. Although a smallplastic strain was introduced during the thermal cyclic test, a shape recovery over 80% was obtained for both MEAs. Training, that is, thethermal cyclic test under the same applied stress, was conducted to investigate the change of the irrecoverable strain and the work output. For45Ti­20Pd­5Ni­25Pt­5Zr, the irrecoverable strain was deleted after 50 cycles, and perfect recovery was obtained. The largest work output(3.5 J/cm3) was obtained under 200MPa. In 45Ti­20Pd­10Ni­20Pt­5Zr, perfect recovery was obtained from the first cycle. However, therecoverable strain was small, and the largest work output was 1.5 J/cm3 under 200MPa. The shape recovery of 45Ti­20Pd­5Ni­25Pt­5Zr ispromising for new HT-SMAs compared with the ternary Ti­Pd­Zr alloys and other HEA-SMAs. [doi:10.2320/matertrans.MT-MAW2019012]

(Received May 24, 2019; Accepted July 26, 2019; Published October 25, 2019)

Keywords: TiPd, TiPt, multi-component alloys, thermal cyclic test, shape recovery, high-temperature shape-memory alloys, B2, B19

1. Introduction

Shape-memory alloys (SMAs) undergo a martensitictransformation (MT), in which their crystal structure changesfrom the austenite phase to the martensite phase withoutcomposition change via diffusion during cooling. SMAs canrecover the introduced strain in the martensite phase via thereverse transformation. The MT temperature (MTT) controlsthe operation temperature of SMAs. Presently, TiNi alloyswith an MT from a B2 structure to a B19A structure areavailable SMAs. For example, they are used in a wide rangeof fields as hot­cold water mixing faucets, shower valves,and oil flow control devices. However, because the MTT ofthe TiNi alloy is <100°C, the operation temperature islimited. The development of high-temperature SMAs (HT-SMAs) will allow the use of SMAs in aerospace andautomotive applications.1) Therefore, HT-SMAs with theMTT above 100°C have been developed. In an early stage,increasing the MTT of TiNi via the addition of alloyingelements such as Hf, Zr, Pd, Pt, and Au was attempted.1­10)

The MT from the B2 structure to the B19 structure occurs inTiPd and TiPt, and their MTTs are >500°C. TiNi, TiPd, andTiPt are fully solid solutions, and addition of Pd or Pt to TiNiincreases the MTT above 500°C.11,12) However, perfect shaperecovery is difficult to obtain. Recently, the addition of Hfand Zr has attracted attention, because nanosized precipitatesof (Ti, Hf )3Ni4 or (Ti, Zr)3Ni4, which are referred to as the Hphase, enhanced the strength of TiNi alloys, and as a result,perfect shape recovery was obtained under a high appliedstress.13­19) However, the MTTs of TiNi­Hf and TiNi­Zr withthe H phase are <300°C.

We focused on TiPt,20­24) TiAu,25,26) and TiPd27­33) alloyswith the MT from the B2 structure to the B19 structure forHT-SMA because their MTTs are approximately 1000, 600,and 570°C, respectively. Because the shape-recovery ratioof the binary Ti­50Pd (at%) determined via a simplecompression test was 39.5%,27) improvement of the strengthof alloys has been attempted via the addition of alloyingelements such as Cr, Nb, Zr, Hf, Mo, W, Ta, Ir, Ru, andCo.28,31) It was found that Zr is the most promising alloyingelement, and the shape recovery was improved to 84.3% viathe addition of 5 at% Zr.27) Then, the composition de-pendence of Zr was investigated, and the martensite finishtemperatures (Mf ) of Ti­50Pd­5Zr, ­7Zr, and 10Zr were 445,400, and 302°C, respectively (reduced from 480°C for Ti­50Pd).32) Perfect shape recovery was obtained for Ti­50Pd­7Zr and 10Zr alloys via cyclic thermal training, althoughit was not obtained for Ti­50Pd­5Zr via cyclic thermaltraining.32) The simultaneous addition of Zr and V to TiPdwas also attempted, because V has the largest strengtheningeffect on TiPd among alloying elements.29­31) Perfectrecovery was obtained for Ti­50Pd­2.5Zr­2.5V and Ti­50Pd­1Zr­4V via thermal cyclic training, but the Mf

decreased to 368 and 424°C, respectively.32) To maintain ahigh MTT and strength, the addition of Pd to TiPt was alsoattempted. The Mf of the Ti­15Pd­35Pt­5Zr was 590°C,which is very high, but perfect shape recovery was notobtained via thermal cyclic training.33)

Recently, high-entropy alloys (HEAs; multi-componentequiatomic or near-equiatomic alloys) have attractedattention as new-concept alloys because their high entropyeffect, severe lattice distortion, sluggish diffusion, andcocktail effect lead to unexpected properties, for example,compatibility of high strength and high ductility.34,35) The+Corresponding author, E-mail: MITARAI.Yoko@nims.go.jp

Materials Transactions, Vol. 60, No. 11 (2019) pp. 2282 to 2291©2019 The Japan Institute of Metals and Materials

severe lattice distortion and sluggish diffusion are expected tocontribute to strong solid-solution hardening of HT-SMAs.The first attempt of an HEA to an SMA was performed forTiZrHfNiCu. The MT from the B2 structure to the B19Astructure with Af of 337°C and Mf of 127°C (higher thanthat of TiNi) was observed.36) The hardness of TiZrHfNiCuwas twice that of TiNi, and perfect shape recovery wasobtained. However, the recovery strain was 10 times smallerthan that of TiNi. Another example was TiZrHfNiCoCu.37)

The MT from the B2 structure to the B19A structure with Af of72°C and Mf of ¹80°C occurred. Large strain recovery of4.9% was obtained under 650MPa in TiZrHfNiCoCu, whilea similar recovery strain of 4.5% in TiNi was obtained under200MPa. The resistance to high stress in TiZrHfNiCoCuwas attributed to the significant solid-solution hardeningeffect. Application of an HEA to a HT-SMA was attemptedfor NiPdTiHf and NiPdTiHfZr.38) The MT from the B2structure to the B19 structure, as well as TiPd, occurred. Thehighest MT, Af of approximately 780°C, and Mf ofapproximately 600°C were obtained for NiPdTiHfZr. Thisindicates that HEAs are promising materials for HT-SMAs.

In this study, the phase transformation and shape recoveryof TiPd-based multi-component alloys, i.e., TiPdNiPtZr, wereinvestigated. The potential of TiPdNiPtZr for HT-SMAs wasdiscussed.

2. Experimental Procedure

15-g ingots of TiPdNiPtZr alloys were melted via the arc-melting method. The nominal compositions of the alloysare presented in Table 1. The multi-component alloys areclassified according to the mixing entropy using thefollowing equations.39)

HEA: �Smix � 1:5R ð1ÞMedium-entropy alloy ðMEAÞ: 1:0R � �Smix � 1:5R ð2Þ

Low-entropy alloy ðLEA, conventionalsolid-solution alloyÞ: �Smix � 1:0R ð3Þ

Here, ¦Smix represents the mixing entropy, and R is a gasconstant (8.314 J/Kmol).

¦Smix is defined by the following equation.

�Smix ¼ �RXn

i¼1xilnxi ð4Þ

Here, xi is the mole fraction of the component i, and n is thenumber of constituent elements. The calculated mixing

entropies are presented in Table 1. For reference, the mixingentropies of Ti­50Pd­5Zr, Ti­50Pd­7Zr, Ti­50Pd­10Zr, andTi­50Pt­5Zr are also shown.

The ingots were encapsulated in a quartz tube with Ar gas,heated at 1000°C for 3 h, and then quenched in water. Themicrostructures of the heat-treated samples were observedusing field-emission scanning electron microscopy (JSM7200 F) with an energy dispersive X-ray spectromentry(EDS) and an electron backscatter diffraction (EBSD) patternanalyzer at 20 kV.

A disc-shaped sample having a diameter of 4mm and athickness of 1mm was cut from the heat-treated ingotsusing a wire electric discharge machine to measure theMT, the austenite start and finish temperatures (As and Af,respectively), and the martensite start and finish temperatures(Ms and Mf, respectively) via differential scanning calorim-etry (DSC) (Material Analysis and Characterization,DSC3200s). The weight of each DSC sample wasapproximately 100mg. The measurements were performedwith a scanning rate of 10°C/min from room temperatureto 700°C. A plate 10mm long and 1.0mm thick was cut fromthe as-cast ingot for X-ray diffraction (XRD) measurements(Rigaku Co., Ltd., RINT TTR-III), which were performedto identify the crystal structure using Cu K¡ radiation at50 kV and 300mA.

To investigate the strength, a compression test wasperformed on heat-treated samples with dimensions of2.0 © 3.0 © 4.0mm3, with an initial strain rate of 3.0 ©10¹4/s at 30°C below Mf and at 30°C above Af (ShimazuCorp., AG-X). A single compression test for obtaining theplastic deformation was conducted for the martensite phase.For the austenite phase, a loading­unloading compressiontest was performed to remove the effect of the stress-inducedmartensite transformation. The sample was deformed to0.25% initially and was unloaded, and the applied strainincreased for each loading­unloading cycle until plasticdeformation occurred. The 0.2% proof stress of the austeniticphase was obtained from the stress­strain curve by plottingthe plastic strain for the maximum stress for each loading­unloading cycle.

Thermal cyclic tests were performed in the temperaturerange of Mf ¹ 30°C to Af + 30°C under compressive stressesof 15, 50, 100, 150, and 200MPa to investigate the shape-memory properties. In the thermal cycling test, the samplewas first heated to Af + 30°C, then cooled to Af ¹ 30°C,and again heated to Af + 30°C. The sample length wasdirectly measured via image tracking using a charge-coupleddevice camera. The shape recovery and the work output werecalculated according to the strain­temperature (ST) curveusing the following equations.

Shape - recovery ratio; r ð%Þ ¼ ¾r=¾t ð5ÞWork output, w ðJ=cm3Þ ¼ ¾r � · ðMPaÞ ð6Þ

Here, ¾r represents the recoverable strain, ¾t represents thetransformation strain, and · represents the applied stress.

The thermal cyclic test was repeated under the same stress,and the changes in the irrecoverable strain and the workoutput were investigated. The repeated thermal cyclic testwas called “training.”

Table 1 Nominal composition of alloys (at%) and mixing entropy "Smix

(J/molK).

Improvement of High-Temperature Shape-Memory Effect by Multi-Component Alloying for TiPd Alloys 2283

3. Results and Discussion

3.1 MicrostructuresBack-scattered electron images of the samples heat-treated

at 1000°C for 3 h followed by water quenching are shown inFig. 1. The dendrite structure remained in all the samples,even after solution treatment at 1000°C. Because thehomogenized microstructure was obtained in Ti­Pd­Zrternary alloys via the same heat treatment at 1000°C for3 h, the dendrite structure indicates slow diffusion in theHEAs and MEAs. The composition of the bright phase alongdendrite boundaries investigated using EDS was 8³12 at%for Pd, 2³7 at% for Ni, 26³30 at% for Pt, and 15³23 at% forZr for all tested alloys as shown in Table 2. Although thecomposition deviation was found among the tested alloys,the ratio between (Ti, Zr) and (Pd, Pt, Ni) was approximately57:43. Since the compositions of Zr and Pt were higher in thebright phase along dendrite boundaries than those of thematrix, it is considered that the bright phase is segregationof Zr and Pt. The precipitates with dark contrast were alsoobserved as shown by arrows in Figs. 1(b), (c), and (d). Theratio between (Ti, Zr) and (Pd, Pt, Ni) was approximately 2:1as shown in Table 2. Then, it is considered these precipitateswith dark contrast are Ti2Pd type precipitates.

In the HEAs, i.e., 35Ti­20Pd­15Ni­15Pt­15Zr and 40Ti­20Pd­15Ni­15Pt­10Zr, the twin structure formed by the MTwas not observed, as shown in Figs. 1(a) and (b). While, thetwin structure was clearly observed in the images with highmagnification in the MEAs, i.e., 45Ti­20Pd­5Ni­25Pt­5Zrand 45Ti­20Pd­10Ni­20Pt­5Zr, as shown in Figs. 1(c) and(d). The crystal structure of the HEAs was investigated viaEBSD, and the B2 structure was observed in the whole area,indicating no phase transformation.

3.2 Martensitic transformationThe DSC curves are shown in Fig. 2. Both endothermic

and exothermic peaks were observed for the MEAs,indicating that a phase transformation occurred, as shownin Figs. 2(c) and (d). However, the endothermic andexothermic peaks were not clearly observed for 35Ti­20Pd­15Ni­15Pt­15Zr as shown in Fig. 2(a). On the otherhand, the small broad peaks were observed in 40Ti­20Pd­15Ni­15Pt­10Zr as shown in Fig. 2(b). However, the EBSDanalysis indicated the crystal structure of 40Ti­20Pd­15Ni­15Pt­10Zr was the B2 structure at room temperature. Theseresults indicate that phase transformation from the B2 toB19 structures was not occurred above room temperature inthe HEAs. The microstructure observations supported theDSC measurement; that is, the twin structure formed duringthe MT was observed in the MEAs (Figs. 1(c) and (d)).However, the twin structure was not observed in the HEAs(Figs. 1(a) and (b)). The stable B2 structure in HEAssuggests that the MTT became lower than room temperature.The reason for this is unclear. Since no clear evidence ofmartensitic phase transformation was found in 40Ti­20Pd­15Ni­15Pt­10Zr, further investigation using high temper-ature X-ray diffractiometry is necessary to understand thebroad peaks shown in Fig. 2(b). In a previous study, theMTT of the equiatomic alloy Ni25Pd25Ti25Hf25 was higher(for example, Af was 714°C) than that of the near-equiatomicalloy Ni35Pd15Ti20Hf30 (Af was 686°C).36) In our case, thealloy composition of the HEAs and MEAs was notequiatomic. Even so, the difference of phase transforma-tion appeared. To understand the phase-transformationbehavior, the MTT change for various combinations ofalloying elements with different compositions must beinvestigated.

The MTTs determined via DSC are presented in Table 3together with the MTTs of the binary alloy Ti­50Pd,27,31) the

Fig. 1 Back-scattered electron of (a) 35Ti­20Pd­15Ni­15Pt­15Zr, (b)40Ti­20Pd­15Ni­15Pt­10Zr, (c) 45Ti­20Pd­5Ni­25Pt­5Zr, and (d)45Ti­20Pd­10Ni­20Pt­5Zr.

Table 2 Phase composition (at%) measured by EDS.

H. Matsuda, H. Sato, M. Shimojo and Y. Yamabe-Mitarai2284

ternary alloy Ti­50Pd­5Zr,27,31) Ti­50Pd­7Zr,32) Ti­50Pd­10Zr,32) and Ti­50Pt­5Zr20,23) for reference. The MTTdecreased with an increase in the Zr content in Ti­50Pdalloys. For example, the Mf and Af decreased from 480 and550°C for Ti­50Pd to 302 and 416°C for Ti­50Pd­10Zr,respectively. The decrease in Mf and Af for 1 at% Zr was 17.8and 13.4°C, respectively. In the MEA with 5 at% Zr, Pd wasreplaced with Ni and Pt. Compared with the MTT of Ti­50Pd­5Zr, the As, Af, and Ms of 45Ti­20Pd­5Ni­25Pt­5Zrwere higher than those of Ti­50Pd­5Zr, although the Mf wasslightly lower than that of Ti­50Pd­5Zr. However, in 45Ti­20Pd­10Ni­20Pt­5Zr, where the amount of Ni increasedfrom 5 to 10 at% and the amount of Pt decreased from 25 to20 at%, the MTT was drastically reduced compared with45Ti­20Pd­5Ni­25Pt­5Zr, and they were almost the same asthose of Ti­50Pd­10Zr. In comparison with MTT in MEAsand Ti­50Pt­5Zr because MEAs included Pt to increase

MTT, MTTs of MEAs were drastically lowered by additionof Pd and Ni. From a different point of view, MTTs of 45Ti­20Pd­5Ni­25Pt­5Zr was held close to those of Ti­50Pd byaddition of 25 at% Pt.

The temperature hysteresis (Af ¹ Ms) of the binary andternary Ti­Pd­Zr was approximately 50°C up to 7 at% Zr.It increased to 94°C for Ti­50Pd­10Zr. The temperaturehysteresis of the MEA was approximately 100°C for 45Ti­20Pd­5Ni­25Pt­5Zr and 45Ti­20Pd­10Ni­20Pt­5Zr andwas closed to that of Ti­50Pt­5Zr. The increase in thetemperature hysteresis indicates that a large temperaturechange is necessary for shape recovery, and it is disadvanta-geous to use SMAs.

XRD patterns measured at room temperature and 750°Cfor the as-cast MEAs are shown in Fig. 3. The B19 and theB2 structures were observed at room temperature and 750°C,respectively, in both alloys, although the peaks of TiO2 were

Fig. 2 DSC curves of (a) 35Ti­20Pd­15Ni­15Pt­15Zr, (b) 40Ti­20Pd­15Ni­15Pt­10Zr, (c) 45Ti­20Pd­5Ni­25Pt­5Zr, and (d) 45Ti­20Pd­10Ni­20Pt­5Zr.

Table 3 MTT (°C) and temperature hysteresis (Af ¹Ms) of Ti­50Pd and Ti­50Pt alloys and MEAs.

Improvement of High-Temperature Shape-Memory Effect by Multi-Component Alloying for TiPd Alloys 2285

also observed owing to oxidation during the measurement at750°C. This indicates that the MT occurred during coolingafter melting and solidification, and the crystal structure ofTiPd was not changed by multi-component alloying.

The temperature hysteresis of HE-SMAs was reported inprevious studies. For example, the Af and Ms ofTi16.667Hf16.667Zr16.667Ni25Cu25 were 611 and 499K, respec-tively and the temperature hysteresis was 112K.37) Foran alloy obtained by adding Co to TiHfZrNiCu, i.e.,Ti16.667Hf16.667Zr16.667Ni25Co10Cu15, the Af and Ms were

71.1 and 36°C, respectively, and the temperature hysteresiswas 35.1°C, which was smaller than that of TiHfZrNiCu.37)

The Af and Ms of Ni35Pd15Ti20Hf30 were 686 and 525°C,respectively, and the temperature hysteresis was 161°C,which is large.38) In the case of the equiatomic compositionfor NiPdTiHf, the DSC peaks were sharp compared withthose of Ni35Pd15Ti20Hf30, and the Af and Ms ofNi25Pd25Ti25Hf25 were 714 and approximately 620°C,respectively.38) The temperature hysteresis was 94°C, whichwas smaller than that of Ni35Pd15Ti20Hf30. In most cases,the temperature hysteresis of multi-component alloys above90°C was larger than that of binary Ti­Pd alloys (40°C).The large temperature hysteresis was attributed to theinhomogeneous microstructure of the as-cast samples.38)

Another explanation was that the severe lattice distortionrestricted the collective growth of martensite plates, causinglarge temperature hysteresis.37) In our case, the dendritestructure remained after heat treatment, and the micro-structure inhomogeneity and severe lattice distortion mayhave been the reasons for the large temperature hysteresis.In both TiHfZrNiCu and NiPdTiHf, the mixing entropy ofTiHfZrNiCu with Co (1.75R) and the equiatomic NiPdTiHf(1.4R) were slightly higher than those in TiHfZrNiCu (1.5R)and Ni35Pd15Ti20Hf30 (1.3R), respectively. The temperaturehysteresis of the alloys with a higher mixing entropy wassmaller. However, in our case as shown in Tables 1 and 3,although the mixing entropy of HEAs was higher than thoseof MEAs, the temperature hysteresis was not observed due tolower MTT below room temperature or no MT in HEAs.Compared with the temperature hysteresis of MEAs andbinary and ternary alloys, the temperature hysteresis ofMEAs was larger than for binary and ternary alloys up to7 at% Zr although the mixing entropy of MEAs was higherthan those of binary and ternary alloys. To understand theeffect of the mixing entropy on the phase transformation,more data are necessary.

3.3 Mechanical propertyThe 0.2% proof stress of the MEAs and ternary Ti­Pd32)

and Ti­Pt23) alloys is presented in Table 4. The 0.2% proofstress is plotted with respect to the test temperature in Fig. 4.Although the strength decreased with an increase in thetesting temperature, the strength of the martensite and

Fig. 3 XRD patterns at room temperature and 750°C for (a) 45Ti­20Pd­5Ni­25Pt­5Zr and (b) 45Ti­20Pd­10Ni­20Pt­5Zr.

Table 4 0.2% proof stress (·0.2, MPa) of ternary Ti­50Pd alloys, and Ti­50Pt and the MEAs. Testing temperatures were at Mf ¹ 30°C andAf + 30°C for Ti­50Pd alloys and MEAs, and at Mf ¹ 50°C and Af + 50°C for Ti­50Pt­5Zr.

H. Matsuda, H. Sato, M. Shimojo and Y. Yamabe-Mitarai2286

austenite phases of 45Ti­20Pd­5Ni­25Pt­5Zr were higherthan those of the ternary TiPd alloys. However, the strengthof 45Ti­20Pd­10Ni­20Pt­5Zr was approximately equal tothat of the ternary TiPd alloys. The 0.2% flow stress of themartensite phase of Ti­50Pt­5Zr was high for the testingtemperature, but the 0.2% flow stress of the austenite phasewas smaller than those of other alloys. The behavior ofmechanical properties of MEAs was closer to that of Ti­Pdalloys rather than that of Ti­Pt alloys.

The solid-solution hardening is often governed by themisfit of the atomic size between solvent and solute atoms.However, in multi-component alloys, a solvent element doesnot exist. Instead of the misfit of the atomic size betweensolvent and solute atoms, a new parameter (¤, representingthe misfit for the average atomic size of the constituentelements) is defined as follows.40)

¤ ¼ 100�ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiXn

i¼1xi 1� ri

�r

� �2rð7Þ

�r ¼Xn

i¼1xiri ð8Þ

Here, xi represents the concentration of element i, rirepresents the atomic radius of element i, and r representsthe average atomic radius of the constituent alloyingelements. The atomic radii of the constituent elements arepresented in Table 5. The average radius and the ¤ parameterscalculated using the atomic radii in Table 5 are presented inTable 6. The average radius was almost the same for all thetested alloys. The ¤ parameters of the MEAs, Ti­50Pd­5Zrand Ti­50Pt­5Zr were compared. The ¤ parameters of bothMEAs were larger than those of Ti­50Pd­5Zr and Ti­50Pt­5Zr. This indicates that severe lattice distortion occurred inthe MEAs. While the ¤ parameters of the Ti­Pd­Zr ternaryalloys increased with the Zr content, owing to the largeatomic size of Zr compared with other elements. The ¤

parameters of Ti­50Pd­10Zr and the MEAs were almostthe same. The large ¤ parameters of the MEAs and Ti­50Pd­10Zr indicate a high degree of solid-solution hardening inboth austenite and martensite. However, a large solid-solutionhardening effect was observed only for 45Ti­20Pd­5Ni­25Pd­5Zr. The strength of the other alloys exhibited almostthe same dependence on the test temperature.

3.4 Shape-memory effectTo investigate the shape-memory effect, thermal cyclic

tests were conducted at 15, 50, 100, 150, and 200MPa for theMEAs. The ST curves for each applied stress are shown inFig. 5. The transformation strain increased with increase ofthe applied stress for both alloys. An almost closed loop, thatis, perfect recovery, was obtained up to 100MPa, but theplastic strain appeared at ²150MPa for 45Ti­20Pd­5Ni­25Pt­5Zr. In the case of 45Ti­20Pd­10Ni­20Pt­5Zr, perfectrecovery was obtained up to 200MPa. This is because theMTT of 45Ti­20Pd­10Ni­20Pt­5Zr was lower than that of45Ti­20Pd­5Ni­25Pt­5Zr, and consequently, the strength ofthe martensite and austenite phases around the MTT washigher for 45Ti­20Pd­10Ni­20Pt­5Zr than for 45Ti­20Pd­5Ni­25Pt­5Zr. The high strength around the MTT prohibitedplastic deformation during the thermal cyclic test. Anotherdifference between 45Ti­20Pd­5Ni­25Pt­5Zr and 45Ti­20Pd­10Ni­20Pt­5Zr was the sharpness of the ST curves.

Fig. 4 0.2% proof stress of ternary Ti­50Pd and Ti­50Pt alloys and theMEAs with respect to the testing temperature.

Table 5 Atomic radii of constituent elements.41)

Table 6 Average atomic radius, �r, and the ¤ parameter.

Improvement of High-Temperature Shape-Memory Effect by Multi-Component Alloying for TiPd Alloys 2287

The ST curves were sharp for 45Ti­20Pd­10Ni­20Pt­5Zr,indicating that the temperature hysteresis during the thermalcyclic test was smaller for 45Ti­20Pd­10Ni­20Pt­5Zr thanfor 45Ti­20Pd­5Ni­25Pt­5Zr.

The shape recovery ratio and the work output werecalculated using eqs. (5) and (6), as described in theexperimental procedure, using the recoverable strain (¾r),the transformation strain (¾t), and the applied stress obtainedfrom Fig. 5. The shape recovery ratio and the work output areshown in Fig. 6. It appeared that there was no plastic strainup to 150 and 200MPa in 45Ti­20Pd­5Ni­25Pt­5Zr and45Ti­20Pd­10Ni­20Pt­5Zr, respectively. However, the en-larged ST curve indicated small plastic strain for all the testedstresses. For both MEAs, the recovery ratio was not 100%.As shown in Fig. 5(a), the plastic strain became large withan increase in the applied stress for 45Ti­20Pd­5Ni­25Pt­5Zr; then, the recovery ratio decreased significantly above100MPa, as shown in Fig. 6(a). The recovery ratio of 45Ti­20Pd­10Ni­20Pt­5Zr was close to 100%, as shown inFig. 6(b). The work output increased with the applied stress.At 200MPa, the work outputs of 45Ti­20Pd­5Ni­25Pt­5Zrand 45Ti­20Pd­10Ni­20Pt­5Zr were 3.5 and 4.5 J/cm3,respectively. Because the Af values of 45Ti­20Pd­5Ni­25Pt­5Zr and 45Ti­20Pd­10Ni­20Pt­5Zr were 600 and 400°C,respectively, during the thermal cyclic test, a large workoutput (3.5 J/cm3) was obtained in 45Ti­20Pd­5Ni­25Pt­

5Zr at 600°C, which is higher than the correspondingtemperatures for ternary Ti­50Pd­7Zr (7 J/cm3 at 430°C) andTi­50Pd­10Zr (5.8 J/cm3 at 400°C). In Ti­50Pt­5Zr, STcurves became trumpet-like shape indicating progress ofdeformation during thermal cyclic test.33) It can be said thatshape memory effect of MEAs is closed to Ti­Pd alloys.

3.5 Training effectTo understand the shape-recovery change during the

thermal cyclic test, repeated thermal cyclic tests, i.e., training,was performed. The training conditions are presented inTable 7. The training was performed for 45Ti­20Pd­5Ni­25Pt­5Zr at 200, 300, and 400MPa for 80, 103, and 71cycles, respectively. For 45Ti­20Pd­10Ni­20Pt­5Zr, thetraining was performed at 200, 300, and 550MPa for 91,93, and 39 cycles, respectively. Before the training, a stress50MPa larger than the testing stress was applied for the firstcycle, and then the applied stress decreased to the testingstress for training. The ST curves of the first and last cyclesare presented in Fig. 7. For 45Ti­20Pd­5Ni­25Pt­5Zr,plastic strain was observed for the first cycle at the appliedstresses of 200, 300, and 400MPa. However, the plasticstrain decreased during the thermal cycles and finally reached0, and perfect recovery was obtained, as indicated by the lastcycle in Figs. 7(a)­(c). Simultaneously, the transformationstrain decreased from the first cycle to the last cycle duringthe training. The transformation strain also decreased with anincrease in the applied stress.

Fig. 5 ST curves of (a) 45Ti­20Pd­5Ni­25Pt­5Zr and (b) 45Ti­20Pd­10Ni­20Pt­5Zr. From bottom to top, the applied stress is 15, 50, 100,150, and 200MPa, where ¾t represents the transformation strain (%) and ¾rrepresents the recoverable strain (%).

Fig. 6 Recovery ratio and work output: (a) 45Ti­20Pd­5Ni­25Pt­5Zr and(b) 45Ti­20Pd­10Ni­20Pt­5Zr.

H. Matsuda, H. Sato, M. Shimojo and Y. Yamabe-Mitarai2288

Similar behavior was observed for 45Ti­20Pd­10Ni­25Pt­5Zr, as shown in Figs. 7(d)­(f ). The plastic strainwas unclear, and almost perfect recovery was obtained for thefirst cycle up to 550MPa. The transformation straindecreased during the training and with an increase in theapplied stress, similar to the results for 45Ti­20Pd­5Ni­25Pt­5Zr. However, the decrease in the transformation strainwas more drastic for 45Ti­20Pd­10Ni­25Pt­5Zr. Forexample, the transformation strain at 300MPa (Figs. 7(b)and (e)) was approximately 0.5% for 45Ti­20Pd­10Ni­25Pt­5Zr, while it was approximately 1% for 45Ti­20Pd­5Ni­25Pt­5Zr.

The changes in the plastic strain, recoverable strain, andwork output during the thermal cyclic test are shown in Fig. 8with respect to the cycle number. The plastic strain drasticallydecreased during the cyclic test and reached 0 after 50 cyclesfor 45Ti­20Pd­5Ni­25Pt­5Zr, as shown in Fig. 8(a). Perfectrecovery was obtained after 50 cycles. The recoverable strainwas the largest under 200MPa and decreased with anincrease in the applied stress (Fig. 8(b)). The recoverablestrain decreased during the thermal cyclic test, but it wassaturated after 50 cycles. Consequently, although the workoutput decreased during the thermal cyclic test, it was also

saturated (Fig. 8(c)). The final work output was approx-imately 3.5, 3.0, and 2.0 J/cm3 under 200, 300, and 400MPa,respectively. The largest work output was obtained under200MPa. For 45Ti­20Pd­10Ni­25Pt­5Zr, the plastic strainwas not observed for the entire thermal cyclic test. Therecoverable strain was almost constant under 300MPa andslightly decreased under 200MPa during the thermal cyclictest, as shown in Fig. 8(d). The recoverable strain was largerunder 200MPa than under 300MPa. The work outputreflected the recoverable strain; that is, a larger work outputwas obtained under 200MPa than under 300MPa, and thework output was almost constant (Fig. 8(e)). The work outputwas not estimated for 550MPa, because the phase trans-formation was unclear under this condition.

Although the behavior observed during the thermal cyclictest was not clearly understood, when the applied stress wasincreased, it is considered that dislocations were introducedby plastic deformation under 400MPa, and these dislocationsprevented the phase transformation. Consequently, therecoverable strain decreased under a high applied stress.Similar behavior was observed for Ti­35Pt­15Pd­5Zr.33)

Dislocations were introduced in the initial cycles, and thenthe irrecoverable strain drastically decreased under 200MPa

Table 7 Training conditions.

Fig. 7 ST curves of (a­c) 45Ti­20Pd­5Ni­25Pt­5Zr ((a) for 1 cycle and 80 cycles at 200MPa, (b) for 1 cycle and 103 cycles at 300MPa,(c) for 1 cycle and 71 cycles at 400MPa) and (d­f ) 45Ti­20Pd­10Ni­20Pt­5Zr ((d) for 1 cycle and 91 cycles at 200MPa, (e) for 1 cycleand 93 cycles at 300MPa, (f ) for 1 cycle and 39 cycles at 550MPa).

Improvement of High-Temperature Shape-Memory Effect by Multi-Component Alloying for TiPd Alloys 2289

compared with 50 and 100MPa. However, the total numberof dislocations introduced was higher under 200MPa, and thedecrease in the recoverable strain was large. Thus, asignificant reduction in the work output was observed. Inthe present study, similar behavior may have occurred under400MPa.

Perfect recovery was obtained for the MEAs in the thermalcyclic test. Compared with the ternary Ti­Pd­Zr alloys orquaternary Ti­Pd­Zr­V alloys,32) the MEAs achieved perfectrecovery under a larger applied stress (between 200­400MPa). The largest applied stresses for obtaining perfectrecovery were 150 and 200MPa for Ti­50Pd­1Zr­4V andTi­50Pd­10Zr, respectively, but they were <65MPa forother alloys. This is attributed to the effect of the latticedistortion or the cocktail effect with an increase in the mixingentropy, although it is not clear for the solid-solutionhardening effect, as shown in Fig. 4. Compared with otherHE-SMAs, the transformation strain of Ni35Pd15Ti30Hf20 wasapproximately 0.8% under 75MPa. Although the recoverablestrain was unclear owing to the limitations of the testmachine, if it is considered that the recoverable strain was

equivalent to the transformation strain, the work outputwas approximately 0.6 J/cm3, which is very small.38) Ouralloys were MEAs, not HEAs, but the results indicate that themulti-component alloys are promising as HT-SMAs.

4. Conclusions

The phase transformation and shape-memory effect ofmulti-component alloys were investigated.(1) Four alloys were prepared: 35Ti­20Pd­15Ni­15Pt­

15Zr, 40Ti­20Pd­15Ni­15Pt­10Zr, 45Ti­20Pd­5Ni­25Pt­5Zr, and 45Ti­20Pd­10Ni­20Pt­5Zr (at%). Ac-cording to the mixing entropy, 35Ti­20Pd­15Ni­15Pt­15Zr and 40Ti­20Pd­15Ni­15Pt­10Zr were identifiedas HEAs, and 45Ti­20Pd­5Ni­25Pt­5Zr and 45Ti­20Pd­10Ni­20Pt­5Zr were identified as MEAs.

(2) In the HEAs, the martensitic phase transformation wasnot observed, and the B2 structure was stable at roomtemperature. However, in the MEAs, the martensiticphase transformation from the B2 structure to the B19structure was clearly observed.

Fig. 8 (a) Plastic strain, (b, d) recoverable strain, and (c, e) work output of (a­c) 45Ti­20Pd­5Ni­25Pt­5Zr and (d­e) 45Ti­20Pd­10Ni­20Pt­5Zr.

H. Matsuda, H. Sato, M. Shimojo and Y. Yamabe-Mitarai2290

(3) The Af and Mf of 45Ti­20Pd­5Ni­25Pt­5Zr were 598and 432°C, respectively, which were comparable tothose of binary Ti­50Pd (480 and 550°C, respectively).The Af and Mf of 45Ti­20Pd­10Ni­20Pt­5Zr wereslightly lower (442 and 256°C, respectively) andcomparable to those of Ti­50Pd­10Zr (416 and302°C, respectively). The temperature hysteresis (Af ¹Ms) of the MEAs was two times larger than that of thebinary and ternary TiPd alloys. The large temperaturehysteresis was attributed to the restriction of the growthof martensite plates during the MT due to the largelattice distortion.

(4) The strengths of both the martensite and austenitephases in 45Ti­20Pd­5Ni­25Pt­5Zr were higher thanthose for the other three alloys. The ¤ parameter, whichindicates the lattice distortion with various atomic sizes,was used to consider the solid-solution hardening effect.There was no large difference in the ¤ parameter amongthe tested alloys, although the ¤ parameter increasedwith the Zr content (large atomic size) and for multi-component alloying. Thus, the significant solid-solutionhardening effect of 45Ti­20Pd­5Ni­25Pt­5Zr was notexplained well using the ¤ parameter.

(5) The shape recovery of the MEAs was examined usinga thermal cyclic test. Although a small plastic strainwas introduced during the thermal cyclic test, shaperecovery of >80% was obtained for both MEAs.

(6) Training was conducted to investigate the changes inthe irrecoverable strain and the work output. For45Ti­20Pd­5Ni­25Pt­5Zr, the irrecoverable strainwas deleted after 50 cycles, and perfect recovery wasobtained. The largest work output (3.5 J/cm3) wasobtained under 200MPa. For 45Ti­20Pd­10Ni­20Pt­5Zr, perfect recovery was obtained from the first cycle.However, the recoverable strain was small, and thelargest work output was 1.5 J/cm3 under 200MPa.

Acknowledgments

The study was partly supported by “Precious MetalsResearch Grant of TANAKA Memorial Foundation”, forwhich the authors express thanks.

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