Journal of Alloys and Compounds 616 (2014) 275–283

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Journal of Alloys and Compounds

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Effect of precipitation hardening and thermomechanical trainingon microstructure and shape memory properties of Ti50Ni15Pd25Cu10

high temperature shape memory alloys

http://dx.doi.org/10.1016/j.jallcom.2014.07.1160925-8388/� 2014 Elsevier B.V. All rights reserved.

⇑ Corresponding author. Tel.: +92 5190856035; fax: +92 519085002.E-mail address: saifur.rehman@smme.nust.edu.pk (S.u. Rehman).

Saif ur Rehman a,⇑, Mushtaq Khan a, A. Nusair Khan b, M. Imran Khan c, Liaqat Ali a, Syed Husain Imran Jaffery a

a School of Mechanical and Manufacturing Engineering, National University of Science and Technology, Islamabad, Pakistanb Institute of Industrial Control System, Rawalpindi, Pakistanc Faculty of Materials Science and Engineering, GIK Institute of Engineering Sciences and Technology, Topi, KPK, Pakistan

a r t i c l e i n f o

Article history:Received 14 May 2014Received in revised form 12 July 2014Accepted 14 July 2014Available online 22 July 2014

Keywords:Transformation temperatureThermal cyclingThermomechanical trainingShape memory properties

a b s t r a c t

Microstructural analysis, evolution of phase transformation temperatures and shape memory propertiesfor Ti50Ni15Pd25Cu10 high temperature shape memory alloys were investigated in solution treated andaged conditions. It is shown that aging at 600 �C for 3 h resulted in the formation of two types of fine pre-cipitates i.e. TiPdCu and Ti2Pd. It is observed that due to the formation of these precipitates, martensitestart temperature under stress free and constrained (500 MPa) conditions are decreased by 32 �C and29 �C respectively. Recovery ratio of 76% obtained for solution treated sample increased to 94% at500 MPa for the 600 �C aged sample; resultantly the irrecoverable strain is decreased by 18%. Trainingfor 5 thermal cycles was carried out under stress free condition resulted in the decrease of thermal hys-teresis by 5 �C for the solution treated sample, whereas thermal hysteresis of the aged sample remainedstable. Similarly thermomechanical training cycles at 500 MPa demonstrated an improvement in therecovery ratio by 14% and 5% for solution treated and 600 �C aged samples respectively. Aging at600 �C for 3 h and thermomechanical training greatly improved the cyclic stability and shape memoryproperties of Ti50Ni15Pd25Cu10 high temperature shape memory alloys.

� 2014 Elsevier B.V. All rights reserved.

1. Introduction

High temperature shape memory alloys specially TiNi-based,are getting more concentration due to the increasing applicationof actuators working in high temperature environment in theautomotive, oil and gas exploration and aerospace industries[1–5].

TiNi-based high temperature shape memory alloys are devel-oped by alloying of third element like platinum (Pt), palladium(Pd), gold (Au), zirconium (Zr) and hafnium (Hf) with binary TiNiand increased the transformation temperatures successfully.Among TiNi-based ternary high temperature shape memory alloys,TiNiPd has got more attention due to its superior qualities like lowthermal hysteresis, better strain recovery under stress free andconstrained conditions, better workability and relatively low castas compared with Pt and Au [6–14].

Like other metallic materials exposed to high temperature, theshape memory behavior of TiNiPd alloy also deteriorates at hightemperature due to thermally driven mechanisms like recovery,recrystallization processes, creep deformation and transformationinduced plasticity [2,15,16] and low critical stress for slip deforma-tion [17]. Due to these effects the functional and dimensional sta-bility of the alloy become poor during operation. There are manytechniques proposed to overcome these problems of high temper-ature shape memory alloys and improve the thermal and dimen-sional stability at high temperature such as: (a) solid solutionstrengthening by quaternary alloying [18–21] (b) precipitationhardening [22,23] (c) cold working and subsequent heat treat-ments [24–28] (d) thermomechanical treatments having thermalcycles under load [29]. Suzuki et al. [19] studied the addition of0.2 at.% boron with Ti50Ni20Pd30 and demonstrated that yieldstrength and ductility are increased at high temperature. Thisincrease was attributed to decrease in grain size due to alloyingof boron and formation of TiB2 particles. Ramaiah et al. [20] studiedthe effects on microstructure and transformation behavior by qua-ternary alloying of 1.0 at.% scandium to Ti50.3Ni24.7Pd25. It was

Table 1Compositional analysis of homogenized Ti50Ni15Pd25Cu10 Alloy.

Element Ti Ni Pd Cu O N

(Weight%) 36.6805 14.0759 39.7061 9.5226 0.00553 0.00937

276 S.ur Rehman et al. / Journal of Alloys and Compounds 616 (2014) 275–283

reported that Mf temperature was decreased by 42 �C while thethermal hysteresis was lowered from 15 �C to 7 �C. It was claimedthat the key factors for obtaining such low thermal hysteresis arethe high purity of alloy, low density of second phase precipitatesand small ratio of twinless/twin martensite. The effects of 5 and10 at.% copper addition in Ti50Ni25Pd25 on the microstructure andshape memory properties were studied. It has been reported thatstrength of B19 martensite and transformation temperature areslightly increased but the thermal hysteresis remained same. Itwas also noted that addition of Cu increases the dimensional sta-bility of alloy at higher stress levels due to increasing the resistanceagainst viscoplastic deformation by solid solution hardening mech-anism [30]. Golberg et al. [31] studied the beneficial effects of coldworking at different percentages followed by annealing at varioustemperatures of Ti50NixPd50�x alloys (x = 10, 15, 20) on the shapememory properties. It was reported that in a thermo-mechanicallytreated Ti50Ni20Pd30 alloy, 100% recovery was obtained in 5.3% totalstrain. The improvement in the shape memory properties of Ti-richTiNiPd was also reported due to creation of fine Ti2Ni precipitates[22].

Many researchers investigated the beneficial effects of precipi-tation hardening for enhancing the functional stability in TiNi-based high temperature shape memory alloys [32–35]. Effects ofaging on the shape memory effect in a Ni-rich TiNiHf aged at550 �C for varying time have been investigated [36] and reporteda better shape memory properties for the time duration of 2 h.Khan et al. [37] studied the effects of cold rolling and precipitationhardening on the cyclic stability. An improvement in the cyclic sta-bility of shape memory properties in Ti50Ni20Pd25Cu5 high temper-ature shape memory alloys has been reported. The effect of Pdcontent in Ti50Ni45�xPdxCu5 (x = 25, 30, 35) on the precipitationbehavior and transformation temperatures was studied [38]. Itwas reported that by increasing the Pd content in the alloy, theprecipitation of nano-scaled precipitates and transformation tem-peratures are increased. The effect of cold rolling ratio on the for-mation of nano-scaled precipitates was investigated by Khanet al. [24]. The formation of precipitates was increased and subse-quently the transformation temperatures and recovery strain werereported to be decreased, as the cold rolling ratio was increasedfrom 10% to 40%. Lin et al. [39] reported the production of twotypes of precipitates (Ti2Pd and Ti(Cu, Pd)2) in Ti50Ni15Pd25Cu10

alloy annealed for 6 h. This investigation was focused only on theeffect of these precipitates on the transformation temperatures inthe stress free condition.

Training at constant tensile stress (thermomechanical treat-ment) is believed to be effective in the improvement of shapememory properties, has been carried out by many researchers.Effects of thermomechanical training and severe plastic deforma-tion on the functional stability in Ti50.5Ni24.5Pd25 high temperatureshape memory alloys were reported [3] and demonstrated theimprovement in shape memory properties like reduction in irre-coverable strain and thermal hysteresis.

Only limited research work has been carried out to explain theeffect of Cu addition to TiNiPd on the microstructure, transforma-tion temperatures and shape memory characteristics like recov-ered and irrecoverable strain at different stress levels. The effectsof aging and thermomechanical training on the transformationtemperatures and other shape memory properties for Ti50Ni15Pd25-

Cu10 have not been studied. Therefore, in the present research;microstructural analysis, evolution of transformation temperaturesand shape memory properties for the solution treated and agedsamples of Ti50Ni15Pd25Cu10 high temperature shape memoryalloys were studied and compared. An improvement in the shapememory properties due to precipitation hardening and thermome-chanical training was presented.

2. Materials and experimental procedure

Ti50Ni15Pd25Cu10 alloy was prepared by vacuum arc melting process using tung-sten electrode and water cooled copper crucible. High purity elemental constitu-ents 99.98 wt.% Ti, 99.98 wt.% Ni, 99.99 wt.% Pd and 99.99 wt.% Cu were used forpreparation of alloy. The cast button was melted 6 times and turned over after eachmelting cycle to ensure alloying homogeneity. The cast button was sealed in quartztube after being vacuumed, filled with argon gas, then homogenized at 950 �C for2 h and water quenched. The homogenized button was sliced into 0.4 mm thickstrips by wire electrical discharge machine (EDM). The 0.4 mm thick strips werecold rolled by 25% and reduced their thickness to 0.3 mm. Samples for differentialscanning calorimetry (DSC), X-ray diffraction (XRD), microstructural studies andshape memory measurement were cut using wire EDM. All the samples were solu-tion treated at 900 �C for 1 h and then some of them were aged at temperature of600 �C for 3 h to compare their properties. All these heat treatments were carriedout in argon filled quartz tubes and quenched in cold water without breaking thequartz tubes.

Stress-free phase transformation temperatures of the solution treated and agedsamples were determined by DSC at a heating/cooling rate of 5 �C/min under vac-uum. Transformation temperatures of the alloy are dependent on the sample prep-aration, as the residual stresses are generated during cutting and grindingprocesses. Thus DSC samples were first cut by wire EDM in the dimensions of2 mm � 2 mm � 0.3 mm and then heat treated. Solution treated and aged sampleswere thermally cycled through complete phase transformation temperature range.Different phases present in the solution treated and aged samples were determinedat room temperature by using XRD with Cu Ka radiation. Microstructural investiga-tions were carried out on metallographically prepared samples using scanning elec-tron microscope (SEM) and transmission electron microscope (TEM) operating at120 kV. Samples for TEM study were prepared by thinning the strips to 100 lmthicknesses by grinding process. Using the spark-erosion process, discs of 3 mmdiameter were cut and then further thinning of the samples were carried out byusing twinjet electro-polisher using an electrolyte consisting of 90 vol.% methanol(CH3OH) and 10 vol.% per chloric acid (HClO4) at �20 �C. Compositional analysisof different phases in the microstructure was measured by energy dispersive spec-troscopy (EDS) attached with SEM on metallographically polished samples in un-etched condition. Compositional analysis of oxygen and nitrogen were analyzedby using oxygen/nitrogen determinator.

Shape memory behavior of solution treated and aged samples were found atdifferent tensile stress levels (100–500 MPa) by ATS creep and rupture testingmachine. Dimensions of all the samples were 15 mm � 2.5 mm � 0.3 mm. Duringshape memory characterization, the tensile samples were first heated up to a tem-perature well above its austenite finish temperature (Af) and then the pre-definedload was applied. The stressed sample was then thermally cycled through full trans-formation range between a temperature 50 �C below its Mr

f and a temperature 50 �Cabove Ar

f (where Mrf and Ar

f are the martensite and austenite finish temperaturesunder applied stress r, respectively). After completing the first thermal cycle at100 MPa, the applied stress was increased with an increment of 100 MPa till theapplied stress was reached to 500 MPa. Transformation temperatures (Ms, Mf, As,Af), recovered transformation strain (erec), irrecoverable strain (eirr), recovery ratio(erec /( erec + eirr)) and work output (product of recovered strain and applied stress)were measured from the strain–temperature curves at different applied stresses.For thermomechanical training, the samples were thermally cycled under constantstress of 500 MPa for 5 times. Heating and cooling rate for thermomechanical cyclingwas kept at 10 ± 2 �C.

3. Results and discussion

3.1. Microstructural characterization

Compositional analysis of Ti50Ni15Pd25Cu10 Alloy was carriedout in homogenized condition and the results are given in Table 1.The contents of oxygen and nitrogen present in the alloy were alsofound out, because these elements affect the transformation tem-perature of TiNi- based shape memory alloys up to great extent[40]. The composition of constituent elements present in the solu-tion treated and aged conditions are given in Table 2.

Fig. 1 shows the backscattered SEM images at differentmagnification of solution treated sample. The microstructure of

Table 2Compositional analysis of solution treated and aged (at 600 �C for 3 h) samples of Ti50Ni15Pd25Cu10 Alloy.

Sample condition Region Ti (at.%) Ni (at.%) Pd (at.%) Cu (at.%)

Solution treated matrix 49.9 14.2 24.7 11.2Ti2Ni 66.8 16.4 12.2 4.6

Aged at 600 �C for 3 h Ti2Pd 60.6 9.3 23.8 6.3TiPdCu 41.1 11.4 26.7 20.8

S.ur Rehman et al. / Journal of Alloys and Compounds 616 (2014) 275–283 277

solution treated sample consists of single phase having self-accommodated martensite structure with low density secondphase precipitates of black colors distributed along the grainboundaries. The average size of these precipitates was measuredto be 0.6–1.2 lm. The average size of the grains was found to be25–40 lm. The microstructure of the aged sample is shown bythe backscattered SEM images at different magnification inFig. 2. In these images two types of precipitates are clearly visiblewhich are randomly distributed along both sides of grain bound-aries. The shapes of these precipitates are mostly lenticular orelliptical.

TEM study was carried out to further investigate the structureof the second phase precipitates and the self-accommodated mar-tensite structure. Fig. 3a and b represent the bright field TEMimages of second phase precipitates formed in solution treatedsample and the corresponding selected area diffraction (SAD) pat-tern respectively. The compositional analysis given in Table 2reveals that the second phase precipitates are Ti2Ni(Pd, Cu). Thecorresponding SAD pattern of precipitate circled A, confirms thatit has the FCC Ti2Ni-type structure with [001] zone axis. The totalvolume fraction of these precipitates was estimated to be less than4%. Such type of precipitates is commonly observed in Ti-rich NiTi-based alloys [41]. It is believed that the Ti2Ni(Pd, Cu) second phaseprecipitates were formed during solidification, because these

5μm

Fig. 1. Backscattered SEM images at different magnification of solution treated sample, sgrain boundaries.

5μm

Fig. 2. Backscattered SEM images at different magnification of sample aged

precipitates have lower melting point than the matrix [42]. Sametype of precipitates were also reported to be formed in ternaryTiNiPd alloys having Pd contents more than 10% [41]. From theseobservations it was concluded that by addition of Cu, the basicmicrostructure of quaternary alloy was not changed. The alloyingelement Cu remained in the matrix of solid solution and did notgenerate any other second phase precipitate, as was reported incase of Sc addition to TiNiPd [7,43].

Fig. 4a and b represent the bright field TEM image of solutiontreated sample and the corresponding SADP respectively. InFig. 4a, a typical B19 martensite structure with plates of two differ-ent martensite variants are clearly identified. The morphology ofthe martensite structure found in this alloy is similar to B19 mar-tensite structure commonly found in ternary TiNiPd-based alloys.It was reported that in such type of structures, long martensiteplates are formed due to the parallel layers of internally twinnedmartensite variants. These long martensite plates are also knownas twin laminates [44]. SAD pattern given in Fig. 4b is takenfrom the twin laminates circled area B represents the existenceof {111} type-I twins with the incident electron beam parallelto [1 �10] axis. It has been reported previously that most ofthe twinning mode in TiNiPd-based alloys was {111} type-I[9,44,45]. This similarity shows that the twinning mode did notchange by addition of Cu in TiNiPd-based alloys.

2μm

howing self-accommodated martensite structure with second phase precipitates at

2μm

at 600 �C for 3 h, showing two types of precipitates (white and black).

(b)

Zone Axis = [001]

A

(a)

200nm

000

040

220

Fig. 3. (a) The bright field TEM image of solution treated sample, showing the second phase precipitates (b) corresponding selected area diffraction pattern of precipitatecircled A in (a).

B

(a)

100nm

(b)

Zone Axis = [1 ]

11 I/II000

00 I 002II

B

(a)

100nm

(b)

Zone Axis = [1 ]

11 I/II000

00 I 002II

Fig. 4. (a) The bright field TEM image of solution treated sample, showing the martensite twin plates (b) corresponding selected area diffraction pattern of the circled area Bin (a), representing {111} type I twins with incident electron beam parallel to [1 �10] axis.

278 S.ur Rehman et al. / Journal of Alloys and Compounds 616 (2014) 275–283

Fig. 5 represents the bright field TEM images at different mag-nification for the 600 �C aged sample. In these images, two typesof precipitates having black and white colors as well as martensiteplates are easily observed. From the backscattered SEM imagesgiven in Fig. 2 and bright field TEM images given in Fig. 5, it canbe confirmed that two types of precipitates are formed when thesample is aged at 600 �C. It has been previously reported that forthe backscattered SEM images, black color represents the existenceof lighter elements while white color represents the existence ofheavier elements [46]. It is well understood that, Ti is lighter thanother constituent elements (Ni, Pd and Cu) due to lower atomicmass, thus black color represents the Ti-rich and white color repre-sents the Ti-lean precipitates. The average sizes of these precipi-tates were measured to be in the range of 400–800 nm. Theresults of EDS analysis given in Table 2 and XRD scans (will be dis-cussed in the next paragraph), have proved that the black andwhite precipitates are the Ti2Pd and TiPdCu respectively. It canbe observed from the compositional analysis of Table 2, that theNi concentration in both precipitates formed in aged sample is lessthan the target composition.

Fig. 6 shows the XRD scans at room temperature for the samplessolution treated and aged at temperature of 600 �C for 3 h. Fromthe XRD scan of solution treated sample, it was confirmed that

B19 martensite phase is present at room temperature. All the peaksavailable were showing the B19 martensite phase and presence ofother second phase precipitate (as were indicated in the SEMimage of Fig. 1 and TEM image of Fig. 3) could not be detecteddue to low volume fraction. By aging at 600 �C temperature for3 h, the XRD profile was significantly changed and peaks other thanmartensite phase were also observed. Two types of peaks, one atthe left side (39.77�) and another at right side (43.67�) of strongestB19 (111) martensite peak, were detected. The former and laterwere showing the formation of Ti2Pd and TiPdCu precipitatesrespectively. For the aged sample, the corresponding peaks of(002) and (022) B19 martensite phases were disappeared, how-ever the other three peaks of B19 (101), (020) and (111) were stillpresent, confirming the presence of martensite phase in the agedsample. The XRD scan obtained for the same alloy which wasannealed at 600 �C after 40% cold rolling reported elsewhere [18]is different significantly from the existing XRD scan. At the men-tioned annealing temperature, only three peaks corresponding toTi2Pd, TiPdCu precipitates and B2 phase, have been shown in theXRD scan. The XRD scan obtained in the present research workafter aging at 600 �C for 3 h shows five peaks representing theTi2Pd, TiPdCu precipitates and martensite phase. The XRD resultsobtained for solution treated and aged samples are strongly

500nm 200nm

Fig. 5. The bright field TEM images of 600 �C aged sample, showing two types of precipitates at different magnification.

30 35 40 45 50 55 60

Intensity

(au)

Angle (2θ)

B19(101

)

B19(111

)

Ti2Pd(103

)

B19(00

2)

B19(020

)

TiPd

Cu(111

)

B19(022

)

Solu�on Treated

Aged at 600oC

B19(111

)

B19(101

)

B19(020

)

Fig. 6. XRD scans at room temperature for the samples solution treated and aged attemperature of 600 �C for 3 h.

100 140 180 220

HeatF

low(m

W/m

g)en

doup

Temperature (oC)

0.2

Hea�ng 5oC/min

Cooling 5oC/min

5th Cycle

5th Cycle

1st Cycle

(a)

MsM As Af

0 50 100 150 200 250HeatF

low(m

w/m

g)en

doup

Temperature (oC)

0.2 Hea�ng 5oC/min

Cooling 5oC/min

1st Cycle

5th Cycle

5th Cycle

(b)

5

15

25

35

45

120

140

160

180

200

0 1 2 3 4 5 6 7 8

ThermalHy

steresis(oC)

Tran

sforma�

onTempe

rature

(oC)

Cycle Number

Solu�on TreatedMsAf

Thermal Hysteresis

AfMs

Aged at 600oC

Thermal Hysteresis

(c)

Fig. 7. DSC curves demonstrating the transformation temperatures of (a) solutiontreated, (b) aged at 600 �C for 3 h and (c) comparison of transformation temper-atures and thermal hysteresis.

S.ur Rehman et al. / Journal of Alloys and Compounds 616 (2014) 275–283 279

confirmed the observations, resulted from EDS, SEM and TEM anal-ysis and proved the formation of two types of precipitates in thesample aged at 600 �C.

3.2. Effect of aging and thermomechanical training on transformationtemperatures

Fig. 7a and b represent the DSC heating and cooling curves forthe evolution of transformation temperatures during five thermalcycles for the samples, solution treated and aged at 600 �C for 3 hrespectively. Transformation temperatures (Ms, Mf, As and Af) forboth conditions were estimated by tangent intersection methodas shown in Fig. 7a. The observed transformation temperaturesand thermal hysteresis (Af – Ms) were plotted in Fig. 7c for compar-ison purposes. DSC curves of solution treated sample show sharppeaks with well estimated phase transformation temperaturesand higher heat absorbed and heat released during heating andcooling curves respectively as compared to DSC curves of agedsample. The transformation peaks in the DSC curves for the agedsample are more broadened and difficult to estimate the As andMf temperatures, thus only the Ms and Af temperatures are com-pared in Fig. 7c.

Phase transformation temperatures for solution treated sample(Ms = 170 �C, Mf = 151 �C, As = 179 �C and Af = 192 �C) in the firstcycle are dropped quickly in the second cycle and then decreasedslightly in the next cycles and become almost stable in the fifthcycle with transformation temperatures (Ms = 166 �C, Mf = 147 �C,

Strain(%

)

Temperature (oC)

2% 500 MPa

400 MPa

300 MPa

200 MPa

100 MPa

50 250150

Cooling

Hea�ng

(b)

200 3000010

Strain(%

)

Temperature (oC)

100 MPa

200 MPa

300 MPa

400 MPa

500 MPa CoolingHea�ng10

%

300100 250200150

Ms

AfAs

Mf

(a)

Fig. 8. Strain-temperature curves showing the shape memory properties ofsamples (a) solution treated and (b) aged at 600 �C for 3 h at stress levels of 100–500 MPa.

280 S.ur Rehman et al. / Journal of Alloys and Compounds 616 (2014) 275–283

As = 170 �C and Af = 183 �C). In solution treated sample the drop inMs and Af temperatures are 4 �C and 9 �C respectively, whereas thedrop in the thermal hysteresis is 5 �C (decreased from 22 �C in thefirst thermal cycle to 17 �C in the fifth thermal cycle). On the otherhand, phase transformation temperatures for the 600 �C aged sam-ple (Ms = 138 �C and Af = 166 �C) in the first cycle are decreasedslightly in the second and third cycles and then remained stable tillfifth cycle with transformation temperatures (Ms = 136 �C, andAf = 164 �C). In the aged sample, the decrease in Ms and Af temper-atures are only 2 �C and thermal hysteresis remained stable at28 �C.

By comparing the transformation temperatures and thermalhysteresis of solution treated and aged samples, it can be easilyobserved that the transformation temperatures significantlyretarded by aging at 600 �C. The reason for this retardation is theformation of Ti2Pd and TiPdCu-types precipitates, which were dis-cussed in Section 3.1. The chemical composition of both precipi-tates present in the 600 �C aged alloy are Ni-lean which in turnincreased the Ni-contents in the matrix and thus resulted in retar-dation of transformation temperatures [47]. The decrease in trans-formation temperatures is also reported [18], when the same alloywas annealed at 600 �C, however this decrease in transformationtemperature is more as compared to the 600 �C aged sample. Forexample, Ms of the solution treated sample was lowered from�192 �C to �122 �C when it was annealed at 600 �C, demonstrateda net decrease of �70 �C. On the other hand, Ms of the solutiontreated sample is decreased from 170 �C to 138 �C when it is agedat 600 �C for 3 h in the present research, with a net decrease of32 �C. This comparison shows that the decrease in transformationtemperatures in the aged condition is less as compared to the sam-ple annealed at the same temperature for the same alloy.

It was noticed that, by increasing the number of thermalcycles the decrease in transformation temperatures occurred inthe solution treated sample is faster as compared to aged sample.The faster decrease in transformation temperatures in solutiontreated sample is due to the generation of dislocations and otherdefects during repeated motion of the austenite–martensiteinterface. This decrease in transformation temperature isunavoidable for quenched and annealed TiNi-based alloys, how-ever, this can be prevented by aging and thermo-mechanicaltreatment [48]. This is why, that the decrease in transformationtemperature in aged sample is less than that of solution treatedsample. It was also observed that the first heating cycle of solu-tion treated sample resulted in higher transformation tempera-tures and then decreased faster in the second cycle. The largestdrop in austenite transformation temperatures after first heatingcycle is attributed to the increase in the formation of dislocationdensities [49]. The dislocation densities have been reported to beincreased remarkably during the first thermal cycle and then itsrate of formation decreases as the number of cycles increases[50]. Therefore due to formation of increased dislocation densi-ties in the first cycle, the austenite transformation temperaturesdropped significantly in the second cycle. This effect takes placein the alloys with low yield strength as observed in the solutiontreated sample. An increase in yield strength by precipitationhardening after aging at 600�C prevents from the dislocation slipand as a result a variation in transformation temperaturebecomes negligible.

By comparing the change in thermal hysteresis of both samples,it was observed that the thermal hysteresis of aged sampleremained stable while the thermal hysteresis of solution treatedsample was continuously decreased. The decrease in thermal hys-teresis after the first thermal cycle is, due to formation of more dis-locations which decreased the As temperature quickly and Ms

temperature slowly and resultantly their difference (thermal hys-teresis) become small.

3.3. Thermal cycling at different stress levels

Thermal cycling at various stress levels was carried out to char-acterize the transformation behavior of solution treated and agedsamples. Fig. 8a and b represent the strain temperature curves atstress levels of 100–500 MPa for the solution treated and 600 �Caged samples, respectively. Transformation temperatures (Ms, Mf,As, Af) as a function of applied stress were measured by the tangentintersection method as shown in Fig. 8a. Recovered transformationstrain, irrecoverable strain, recovery ratio and work output weremeasured from the strain–temperature curves at different appliedstresses and plotted in Fig. 9a and b. Recovered transformationstrain during heating at each cycle demonstrates the amount ofwork output produced by the alloy. Similarly, the amount of irre-coverable strain during heating cycle represents the plastic or per-manent deformation and shows the dimensional instability of thealloy.

Fig. 8a and b represent that all the transformation temperaturesare increasing with increasing stress level. For example the Ms andAf temperatures of solution treated sample is 198 �C and 245 �Cunder 100 MPa, while the same transformation temperatures areincreased to 234 �C and 283 �C under 500 MPa respectively. Simi-larly the Ms and Af temperatures of the aged sample is 162 �Cand 225 �C under 100 MPa, and the same are increased to 205 �Cand 242 �C under 500 MPa respectively. This behavior of increasingtransformation temperatures by increasing the applied stressobeys the Clausius–Clapeyron (Cs–Cl) relationship [51] in bothconditions. By comparing the transformation temperatures of bothsamples, it is observed that the transformation temperatures of the600 �C aged sample are significantly less than the solution treated

0.0

0.4

0.8

1.2

1.6

0

1

2

3

4

5

6

0 100 200 300 400 500 600 700 800

Irrecoverab

leStrain(%

)

RecoveredStrain(%

)

Stress (Mpa)

Solu�on TreatedRecovered StrainIrrecoverable Strain

Aged at 600°CRecovered StrainIrrecoverable Strain

(a)

0

5

10

15

20

25

30

50

60

70

80

90

100

110

0 100 200 300 400 500 600 700 800

WorkOutpu

t(J/cm

3 )

Recovery

Ra�o

(%)

Stress (Mpa)

Solu�on TreatedRecovery Ra�oWork Output

Aged at 600°CRecovery Ra�oWork Output

(b)

Fig. 9. Variation in (a) recovered and irrecoverable strains (b) recovery ratio andwork output as a function of applied stress for solution treated and 600 �C agedsamples.

100 150 200 250 300 350

Strain(%

)

Temperature (oC)

500 MPa (a)5th Cycle

1st Cycle

4%

Hea�ngCooling

30 80 130 180 230 280 330

Strain(%

)

Temperature (oC)

1st Cycle

5th Cycle

0.5%

500 MPa (b)

Hea�ng

Cooling

Fig. 10. Strain–temperature curves showing thermomechanical training of (a)solution treated and (b) 600 �C aged samples at constant stress of 500 MPa.

S.ur Rehman et al. / Journal of Alloys and Compounds 616 (2014) 275–283 281

sample at each corresponding stress level. This behavior is identi-cal to the results obtained by DSC measurements for the stress freecondition. The same reason for the decreased transformation tem-peratures in case of 600 �C aged sample discussed in Section 3.1 issufficient to justify this behavior.

When the applied stress is increased for solution treated andaged samples; the recovered transformation strain is alsoincreased, as shown in Fig. 9a. However the recovered transforma-tion strain increased rapidly in solution treated sample as com-pared to the 600 �C aged sample. At 100 MPa the recoveredtransformation strain in the solution treated and 600 �C aged sam-ples are 3.95% and 0.45% respectively and the same are increased to5.0% and 0.68% at 200 MPa. As the applied stress increases, thenumber of favored martensite variants also increases [52] causingthe transformation strains to increase. At 300 MPa, the recoveredstrain of solution treated sample reached to its highest value of5.2% and then saturated up to stress level of 500 MPa. Howeverthe 600 �C aged sample is not behaving in a similar way like solu-tion treated sample and its recovered strain increases continuouslyby increasing the stress up to 500 MPa. The maximum recoveredstrain obtained for 600 �C aged sample is 1.4%.

It can also be observed from Fig. 9a, that irrecoverable strain forsolution treated sample at each stress level is higher than the600 �C aged sample at corresponding stress level. For examplethe irrecoverable strain in the solution treated sample is 0.1% at100 MPa and increased to 1.6% at 500 MPa while the irrecoverablestrain for aged sample is 0% at 100 MPa and increased to only 0.09%at 500 MPa.

Fig. 9b represents that the recovery ratio of 97% at 100 MPa isdecreased to 76% at 500 MPa and the work output is increased

from 3.95 J/cm3 to 26 J/cm3 for the solution treated sample. Onthe other hand the recovery ratio of 600 �C aged sample remained100% up to stress level of 200 MPa and then decreased to 94% at400 MPa and remains stable at 500 MPa. For the 600 �C aged sam-ple the work output of 0.45 J/cm3 at 100 MPa is increased to 7 J/cm3 at 500 MPa.

By comparing the results shown in Fig. 9a and b, the recoveredstrain and consequently the work output exhibited by solutiontreated sample are higher than the 600 �C aged sample at eachstress level. On the other hand, the irrecoverable strain and recov-ery ratio for the solution treated sample are lower as compared to600 �C aged sample. The higher recovered strain in case of solutiontreated sample is attributed to the presence of more volume frac-tion of martensite phases as indicated in XRD scan of Fig. 6. Thereason for the lower irrecoverable strain and higher recovery ratioobtained for the 600 �C aged sample is the formation of two typesof fine precipitates as discussed in Section 3.1. Due to the presenceof these precipitates, a remarkable resistance against the perma-nent or plastic deformation is induced at all stress levels and dem-onstrates more dimensional stability at higher temperatures. It canalso be concluded that due to the formation of fine precipitateswhich are stable till 600 �C increased the working range of TiPdN-iCu-based alloys. From these results it can be concluded that agingat 600 �C for 3 h significantly improved the cyclic stability of theTi50Ni15Pd25Cu10 at stress level of 500 MPa.

3.4. Thermal cycling at constant stress

Fig. 10a and b represent the thermomechanical training of solu-tion treated and 600 �C aged samples respectively at constantstress of 500 MPa. Each sample was thermally cycled 5 timesthrough complete transformation temperature range. Shape

30

40

50

60

50

70

90

110

ThermalHy

steresis(°C)

Recovery

Ra�o

(%)

Number of Cycles

Recovery Ra�o

Solu�on Treated

Thermal Hysteresis

Recovery Ra�o

Aged at 600°C

Thermal Hysteresis

(b)

0.0

0.5

1.0

1.5

2.0

2.5

0

2

4

6

0 1 2 3 4 5 6 7 8

0 1 2 3 4 5 6 7 8

IrreccoverableStrain(%

)

RecoveredStrain(%

)

Number of Cycles

Recovered Strain

Solu�on Treated

Irrecoverable Strain

Recovered Strain

Aged at 600°C

Irrecoverable Strain

(a)

Fig. 11. Variation in (a) recovered and irrecoverable strains (b) recovery ratio andthermal hysteresis with respect to the number of thermal cycles for solution treatedand 600 �C aged samples at constant stress of 500 MPa.

282 S.ur Rehman et al. / Journal of Alloys and Compounds 616 (2014) 275–283

memory characteristics extracted from the strain temperaturescurves were re-plotted in Fig. 11. Fig. 11a represents the variationin recovered and irrecoverable strains with respect to the numberof training cycles, while Fig. 11b represents the change in recoveryratio and thermal hysteresis as a function of training cycles.

From Fig. 11a, it can be observed that the recovered strain forthe solution treated sample (5.2%) is remarkably greater than thatof 600 �C aged sample (1.6%). High value of recovered strain forsolution treated sample is due to the higher volume fraction oftransforming martensite phase. It can also be observed that theirrecoverable strain for solution treated sample is more than theaged sample at each thermal cycle. The reason for this increasein the irrecoverable strain is mainly due to the greater thermal hys-teresis of solution treated sample [52], as shown in Fig. 11b. It isobserved that the irrecoverable strain in the second cycle isreduced relatively more rapidly in solution treated sample andthen decreased slowly. The rapid decrease of irrecoverable strainin the second cycle is attributed to the generation of internal stres-ses, when the dislocation structure is aligned in the stress direction[30]. Due to these aligned dislocations, the generation of furtherdislocations is reduced. Thus the irrecoverable deformationreduces by increasing the number of cycles due to the trainingeffect [2]. Thus thermomechanical training provides an improve-ment in the decrease of irrecoverable strain in case of solution trea-ted sample as compared to aged sample.

It can be seen from Fig. 11b that the thermal hysteresis for solu-tion treated sample is higher (49 �C) than the aged sample (37 �C)in the first cycle and increases continuously as the number oftraining cycles increase. Contrary to the solution treated sample,thermal hysteresis of the aged sample remains constant at 37 �C.The increasing trend in thermal hysteresis of solution treatedsample is due to the poor resistance against the defect generation

during constrained thermal cycling. Thus aging at 600 �C for 3 hsignificantly improved the thermal stability of the alloy as com-pared to solution treated condition.

Recovery ratio obtained for the 600 �C aged sample is higherthan the solution treated sample at each thermal cycle as shownin Fig. 11b. However the recovery ratio for both samples is improv-ing continuously with the increasing number of training cycles.The recovery ratio obtained for solution treated sample is 69% inthe first cycle and then it is increased to 83%, showing an improve-ment of 14%. Similarly for the aged sample the recovery ratio of94% in the first cycle is improved to 99% in the fifth cycle,indicating an increase of 5%. It can be proved on the bases of theseobservations that thermomechanical training is contributing sig-nificantly toward the improvement of recovery ratio in solutiontreated and 600 �C aged conditions.

4. Conclusion

Microstructure, phase transformation temperatures and shapememory properties for Ti50Ni15Pd25Cu10 high temperature shapememory alloy for solution treated and 600 �C aged conditions wereinvestigated. Thermomechanical training cycles under stress freeand constrained (500 MPa) conditions, were also employed toimprove the shape memory properties and cyclic stability of thealloy. The important findings obtained from the present researchare summarized as follows.

1. Aging at 600 �C for 3 h resulted in the formation of two types offine precipitates i.e. TiPdCu and Ti2Pd. Due to the generation ofthese precipitates, the martensite starts temperature understress free and constrained (500 MPa) conditions are decreasedby 32 �C and 29 �C respectively.

2. Due to precipitation process in the 600 �C aged sample, the irre-coverable strain (1.6%) for the solution treated sampledecreased to 0.09%. Resultantly the recovery ratio of 76% forsolution treated sample is increased to 94% at 500 MPa.However due to decrease in the recovered strain, the workoutput is decreased by 19%.

3. Training for 5 thermal cycles under stress free conditionresulted in the decrease of thermal hysteresis by 5 �C for thesolution treated sample, whereas thermal hysteresis of the agedsample remained stable.

4. Thermomechanical training cycles at 500 MPa for 5 timesincreased the recovery ratio by 14% and 5% for solution treatedand aged samples respectively.

5. It can be concluded that thermomechanical training understress free and constrained conditions are more effective inthe solution treated condition as compared to the sample agedat 600 �C for 3 h.

6. Transformation temperatures and shape memory propertiesobtained for the aged sample, demonstrated stable behaviorduring thermomechanical cycling due to generation of lessdefects. From the stable behavior it can be concluded that agingat 600 �C for 3 h improved the cyclic stability of the alloy.

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