Journal of Alloys and Compounds 645 (2015) 213–222

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Journal of Alloys and Compounds

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A systematic investigation on the role of microstructure on phasetransformation behavior in Ni–Ti–Fe shape memory alloys

http://dx.doi.org/10.1016/j.jallcom.2015.04.2240925-8388/� 2015 Elsevier B.V. All rights reserved.

⇑ Corresponding author. Tel.: +91 784 0022252.E-mail address: basu.ritwik@gmail.com (R. Basu).

Ritwik Basu a,⇑, Mostafa Eskandari b, Lalit Upadhayay a, M.A. Mohtadi-Bonab b, Jerzy A. Szpunar b

a Department of Mechanical Engineering, ITM University, Gurgaon 122017, Indiab Department of Mechanical Engineering, University of Saskatchewan, Saskatoon S7N5A9, Canada

a r t i c l e i n f o a b s t r a c t

Article history:Received 6 April 2015Received in revised form 28 April 2015Accepted 30 April 2015Available online 8 May 2015

Keywords:Ni–Ti–Fe shape memory alloysThermo-mechanical treatments (TMTs)Enthalpy of transformationTransformation temperature (TT)Differential scanning calorimetry (DSC)Electron backscattered diffraction (EBSD)

A systematic study on the phase transformation behavior in a polycrystalline Ni–Ti–Fe shape memoryalloy is reported. The investigation was carried out through series of differential scanning calorimetry(DSC) tests on three different samples with distinctly different microstructures of the alloy processedthrough different thermo-mechanical routes. The applied procedures involved cold rolling and marform-ing (rolling in liquid nitrogen) followed by short annealing treatments to recover the cold workedmicrostructure. These treatments altered the microstructure in terms of grain size, misorientation(defects) and texture. Three different microstructures were generated through adopted deformation pro-cesses (i) lower degree of texture, (ii) textured and (iii) refined grains with texture. The influence of themicrostructural parameters on the phase transformations in terms of transformation temperature (TT)and associated energy of absorption/release was investigated. Temperature of transformation (TT)between different samples was affected predominantly by the grain size difference while enthalpy oftransformation (austenite M martensite) was mainly due to differences in texture, stored energy andin-grain misorientation. Though the differences in TT and enthalpy were not very significant, theseresults form a guideline for microstructure tailoring of bulk scale manufacturing of these alloys.

� 2015 Elsevier B.V. All rights reserved.

1. Introduction

Among many class of shape memory alloys (SMAs) identified todate, Ni–Ti based SMAs are notable for attractive characteristics inaddition to good shape memory and pseudoelastic properties, suchas strength [1–3], plasticity [4] and corrosion resistance [5–7] mak-ing them useful in applications ranging from aerospace to medical.There is a potential demand for SMAs with a stable cyclic response.Unstable transformation temperatures and degradation of shapememory or pseudoelasticity upon thermal and mechanical cyclinglimits the use of these alloys in several applications which requirehigh fatigue strength. Fatigue resistance has been shown to begreatly improved by different thermo-mechanical treatments(TMTs) [8]. Significant development in microstructures occursthrough different TMTs. All these TMT routes aims to strengthenthe austenite matrix to reduce cyclic degradation of the shapememory properties [9]. Some of the important microstructural fac-tors leading to austenite strengthening are grain size refinement,

formation of coherent precipitates, development of desired crystal-lographic texture and increased dislocation or defect densities dur-ing TMTs [10].

The different thermo-mechanical routes adopted for processingof bulk Ni–Ti alloys are high temperature deformation, ausforming,severe plastic deformation and marforming, to list a few. A limitednumber of work exists on the effect of ausforming on martensitictransformation behavior and fatigue properties of Ni–Ti alloys[8,9,11]. ‘Ausforming’ comprises deformation in the austenite phaseso that defects are introduced in the austenitic grains. Severe plas-tic deformation (SPD) processes such as high pressure torsion (HPT)or equal channel angular pressing (ECAP) have been adopted toproduce ultrafine grained alloys [12,13]. Marforming involvesdeformation of the material in the martensitic condition [14–17].All these deformation processes are accompanied by short anneal-ing treatments to preserve the footprints of deformationmicrostructure. However in all these studies, the effect of TMTsand microstructure developments on phase transformation behav-ior has been left unexplored. These studies were mostly restrictedto relating the resulting microstructures to conventional mechani-cal properties such as strength and ductility, but did not preciselyaddress the effect of microstructures on shape memory properties

Fig. 1. Schematic of the adopted deformation processes. Route A is cold rolling in austenite phase and Route B is marforming. Deformations through both routes were carriedthrough 10% reduction in thickness. Marforming involves treating the sample with LN2 prior to deformation so that deformation is carried out in martensite phase. The finalannealing treatments for samples A and B were kept identical; 730 �C for 0.5 h for both deformed samples.

214 R. Basu et al. / Journal of Alloys and Compounds 645 (2015) 213–222

and phase transformation behavior in Ni–Ti alloys such as the TTsand enthalpy.

As mentioned earlier the applications of these alloys are wide.The transformation behavior of these alloys can be tuned accord-ingly depending on the nature of application. Both TT and enthalpyof transformation are largely dictated by the presence of secondaryphase (precipitate), chemical composition (trace elements addi-tions) and previous thermomechanical treatments. For instanceactuators and sensors require small transformation hysteresis withcharacteristic sharp and intense peaks which is a signature of fasttransformation with larger enthalpy changes. Higher enthalpymeans transformation occurs in higher volume of material. The pre-sent study therefore attempts to relate different microstructuresgenerated through combination of cold rolling and maforming tochanges in the phase transformation behavior in terms of transfor-mation temperature and enthalpy. This particular work would be ofinterest to researchers who are constantly hunting for processes todesign shape memory components with improved transformationcharacteristics. It has been shown that subtle changes in themicrostructure can lead to changes in the transformation charac-teristics. The results constitute a fundamental understanding todiscuss the microstructural mechanisms possibly responsible forchanges in the observed transformation behavior. The microstruc-tures for all the conditions of TMTs were investigated through elec-tron backscattered diffraction (EBSD) techniques. The observedmicrostructures were related to calorimetric measurements fromdifferential scanning calorimetry (DSC). The details of microstruc-tural developments through the adopted deformation processeswere beyond the scope of the present study (Subject of a separatestudy: PhD thesis of R. Basu, IIT Bombay; 2012).

2. Experimental

2.1. Materials and preparation

Alloy of Ni47–Ti50–Fe3 (in at.%) were supplied by Hlmet Co., Ltd., China in theform of hot rolled cylindrical bars of dimension 18.0 mm(Ø) � 50 mm (l). Theas-received alloy exhibited significant heterogeneity in microstructure. To introducehomogeneity, the alloy was further cold rolled through 50% reduction in thickness in3 passes with intermediate annealing at 800 �C for 25 min followed by final anneal-ing at 800 �C for 1.5 h. This was the starting material for all subsequent deformationprocesses; here onwards would be referred as ‘‘S’’. The material was deformed toanother 10% reduction through two different routes A and B. The route A was thenormal cold rolling process, where the material was deformed in austenitic condi-tion in a single pass. The route B was a complex mixed deformation process where

the starting material S was ‘‘marformed’’ under liquid nitrogen (LN2). This was doneby ‘‘dipping’’ the sample in LN2 (about 5 min) for complete martensitic transforma-tion and then subjecting to rolling to 10% reduction in single pass. The rolled samplewas then allowed to heat up to the ambient temperature. The schematic of theadopted deformation processes are summarized in Fig. 1. Samples obtained throughroute A and B were generically termed as A and B respectively. The final annealingtreatment for samples A and B were kept identical, 730 �C for 0.5 h.

Samples used in this study were polished to mid-width thickness through con-ventional metallographic polishing routes. For EBSD measurements, the final stageof sample preparation was followed by polishing the surface with 0.04 lm colloidalsilica solution for 40 min to obtain a relatively strain free surface.

Specimens for calorimetric characterization (DSC) were cut using slow speeddiamond wafer blade to avoid any deformation that may arise in the microstructure.

2.2. Experimental procedures

A SU 6600 Hitachi field emission scanning electron microscope (FE-SEM)equipped with an Oxford Instruments Nordlys Nano EBSD detector was used tocarry out EBSD measurements. The patterns were acquired using the AZTEC 2.0 dataacquisition software compatible with the EBSD detector with a binning of 4 � 4 pix-els. The software indexes the diffraction patterns to evaluate the crystallographicorientation of the selected region. A typical step size of 0.6 lm was used for obtain-ing large area (420 lm � 210 lm) scans for all the deformed samples. The condi-tions of beam and the video were kept identical for all the sets of scans. The EBSDraw data were further treated for analysis using the Oxford Instruments Channel5 post processing software. This software allows the identification of the grainsand grain boundaries. Grain boundaries (GB) were defined as continuous regionsof misorientation (misorientation angle >5 deg), whereas misorientations were typ-ically represented as average point-to-point misorientation inside an identifiedgrain: local average misorientation. Average misorientation represented the averagemisorientation of each measurement point with respect to all eight (as in a squaregrid) neighbors (kernel): provided any of these did not exceed 5 deg.

Calorimetric measurements were performed by differential scanning calorime-try (DSC) using a TA Instrument’s Q2000 DSC unit under an Ar atmosphere with aLN2 cooling attachment. DSC scans for all specimens A, B and S were recorded ata heating and cooling rate of 3 �C/min. The heating–cooling cycle was performedin the temperature span of �50 �C to 50 �C.

The stored energy was estimated from EBSD misorientation criterion using theformula [18]

Eb ¼ SV �c ¼ 3�cdECD

ð1Þ

�c ¼X62:8

5

½cðhÞf ðhÞ� ð2Þ

cðhÞ ¼ cmhhm

� �1� ln

hhm

� �� �; h 6 hm ð3Þ

cðhÞ ¼ cm; h P hm ð4Þ

Fig. 2. (a) Microstructural developments for all the conditions of deformations, included are the EBSD measured inverse pole figure (IPF) map, texture IPF and Euler space plot(ODFs) at 0 and 45 deg sections of samples S, A and B. Inset shows the direction of Rolling (RD) and transverse direction (TD) marked and IPF color triangle representing thecrystal directions parallel to the normal direction (ND) of the sample reference plane and (b) area fractions as estimated from EBSD ODFs of important fibers for all thedeformation conditions. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

R. Basu et al. / Journal of Alloys and Compounds 645 (2015) 213–222 215

Eb is the stored energy due to boundary dislocations in J mole�1 obtained by multi-plying the average energy per unit boundary area (�c) to the area per unit volume

SV � 3dECD

� �[18] where, dECD is the average grain size measured from equivalent circle

diameter considering all misorientations between 5 deg 6 h 6 62.8 deg [18]. f(h) rep-resents the boundary fraction for a given misorientation, cm ¼ 0:617 J m�2 [19] is theenergy per unit area of a HAGB, h is the boundary misorientation and hm=15 deg is themisorientation angle above which the energy per unit area is independent of misori-entation angle. Here boundary misorientations between 5 and 15 deg are categorizedas low-angle boundaries (LAGBs) while misorientations P15 deg are high-angleboundaries (HAGBs).

3. Results

Fig. 2a summarizes the microstructural developments for sam-ples A, B and S. Included are the EBSD-inverse pole figure (IPF)map, the EBSD measured IPFs and the orientation distribution func-tions (ODFs). The colors in the IPF maps are representative of differ-ent z-normal directions (NDs) to the sample plane. The differencesin crystallographic NDs were noticeable in the IPF plots. The impor-tant texture components are clearly indicated on the ODFs. The

Fig. 3. Significant differences in microstructural parameters between samples A, B and S. Included are (a) austenite grain size distribution, (b) average austenite grain size, (c)local misorientation distribution (kernel), (d) local average misorientation and (e) grain boundary fraction plot for the cold rolled (A), marformed (B) and parent (S) samples.

Table 1Transformation temperatures (TTs) and enthalpy of transformation (H), as estimated by differential scanning calorimetry used in the current study.

Ms (�C) Mf (�C) As (�C) Af (�C) Enthalpy (H) (austenite ? martensite) (J/gm) Enthalpy (H) (martensite ? austenite) (J/gm)

Sample A �21.3 �24.7 �21.8 �17.87 3.26 3.89Sample B �21.85 �25.78 �20.00 �17.2 3.92 3.34Sample S �21.5 �25.0 �22.84 �18.2 2.77 5.06

216 R. Basu et al. / Journal of Alloys and Compounds 645 (2015) 213–222

contour levels as represented in both IPF and ODF are decided fromthe continuous orientation distribution in the Euler space repre-senting the strength of the texture as the number of times randomoccurrence. The parent sample S exhibits a weak texture with few

grains oriented along ND//h110i as seen in the IPF. Significant dif-ferences in the microstructures were brought out through normalcold rolling and marforming, performed separately on the parentsample S through equal reduction in thickness. Room temperature

Fig. 4. DSC results for samples A, B and S after annealing. The annealing treatmentwas necessary, at all deformation stages, to regain the calorimetric signatures of thephase transformation.

R. Basu et al. / Journal of Alloys and Compounds 645 (2015) 213–222 217

rolling led to appearance of ND//h111i orientations in sample A,suggesting that the orientation of most of the grains has rotatedtoward ND//h111i or in its vicinity. With marforming ND//h111iwas slightly strengthened along with some presence ofND//h110i orientation as apparent in measured IPF plots.Furthermore, marforming introduced a significant grain refine-ment. The intensity changed from the lowest of 2.68 to a maximumof 4.51 times random. While IPFs bring out orientation of crystalsalong one component (ND), ODF plots provide more completeinformation on the crystallographic orientation or fibers wrt to

Fig. 5. (a) transformation temperature (TT) as a function of average austenite grain sizaverage misorientation and (c) elastic stored energy in Ni–Ti–Fe samples A, B and S.

the rolling direction or the sample reference frame. No significantfiber was observed in Sample S except that a small spread of orien-tation around ND//h110i was present. However with cold rollingand marforming c-fiber became distinct. Marforming also showedpresence of weak Goss orientation. Fig. 2b presents the relativepresence of different fiber orientations for all the conditions ofdeformation.

More importantly, combinations of cold deformation and mar-forming followed by short annealing treatments introduced signif-icant differences in grain size, misorientation and grain boundaryfractions. This point is illustrated in Fig. 3a–e. Both cold rolledand the parent microstructure exhibited distribution of grain sizeover a wide range. Though samples A and B were produced throughequal degree of deformation, the average grain size of marformedsample B was reduced by 2.5 times from the starting structure Sas opposed to 1.25 times reduction in the cold rolled sample A.Similar reductions in the grain size through normal cold rolling inthe austenite phase requires higher loads. Furthermore, rollingin the austenite phase led to severe cracking which was absent insample B. The average in-grain misorientation of sample A washigher than the sample B with a significant shift in the misorienta-tion distribution. On the whole, ND//h110i orientations showedhigher in grain misorientation as compared to ND//h111i orND//h112i for all the conditions of deformation. Both cold rollingand marforming led to a decrease in the fraction of LAGBs. Nonoticeable differences were observed for the HAGB fraction.

The objectives of the present study as indicted were to investi-gate microstructural effects on the calorimetric signatures ofaustenite M martensite transformation. The DSC measurements interms of net heat flow and temperature of transformation are

e; enthalpy of transformation (H) as a function of (b) in-grain misorientation/local

Fig. 6. Schematic representation of different martensite variants growing inside parent austenite grains. The growth of the variants take place in a way so as to minimize theelastic strain to accommodate the shape change. The grain boundaries act as barriers for propagation of martensite plates. Adopted from PhD thesis of Dizertacni Prace, BrnoUniversity of Technology, Czech Republic.

Fig. 7. Plot of chemical free energies of austenite and martensite phase as a functionof temperature [29].

218 R. Basu et al. / Journal of Alloys and Compounds 645 (2015) 213–222

summarized in Table 1. Fig. 4 represents the plot for the calorimet-ric signature for the samples S, A and B. Martensite transformationcan initiate during cooling the material below Ms, defined as thetemperature at which the martensitic transformation starts. Mf isthe temperature at which the martensite transformation ends.The transformation is recoverable during heating, with As thetemperature at which the reverse austenitic transformation(martensite ? austenite) begins and Af the temperature at theend of the reverse austenitic transformation.

Fig. 5a summarizes the austenite grain size (AGS) measuredfrom EBSD as a function of the experimental values of Ms, Mf, As

and Af obtained from DSC for samples A, B and S. Results suggestthat these transformation temperatures are sensitive to small vari-ations in AGS. The grain size estimated for the samples however donot show pronounced effect in the transformation enthalpy.Though sample S and A presented nearly the same average grainsize, however the enthalpy of transformation in both forward andreverse transformation changed considerably illustrated inTable 1.1 Among all microstructural parameters, the enthalpy oftransformation fitted best with elastic stored energy and misorienta-tion. This point is supported by Fig. 5(b and c): increasing storedenergy correlated with an increase in the enthalpy of martensitictransformation and corresponding decrease in the enthalpy ofaustenite transformation. It is to be noted that the typical accuracy

1 For clarity, the enthalpy terms in both forward and reverse transformations wereconsidered as absolute values.

of estimated enthalpy and transformation temperatures are consid-ered to be within 4% of the mean values – this probable ‘inaccuracy’depends on the quality of the baseline calibration. The measured Hand TTs are marked with suitable error bars in Fig. 5a–c.

4. Discussions

The austenite ? martensite transformation is characterized by ahigh rate of nucleation and growth with an interface front thatapproximates to an incoherent or semi-coherent interface betweenparent and product phases i.e. an interface exhibiting an incompat-ibility in the crystal structure of the two adjoining phases [20]. Thisstructural mismatch in the interface plane between parent and pro-duct phases contribute to the glissile nature of interface, i.e. peri-odic accommodation of dislocations in order to minimize theelastic strain energy associated to the interface [20,21]. The growthof the martensitic transformed region is governed by the disloca-tion motion and proceeds as long as the transformation interfaceremains glissile. Besides the glissile behavior of the interface, theparent lattice undergoes distortion in a cooperative manner (shearlattice distortion) into product lattice by moving atoms by less thanone interatomic distance [22]. In addition, the growth of martensiteplates is impeded when it encounters a strong obstacle typically agrain boundary or a previously formed martensite plate [23].Fine-grained austenite microstructure offers greater resistanceand acts as barrier for martensite plates to grow. During thermalinducement of martensite in polycrystalline material, the crystallo-graphic variants of the martensite phase forms plate-like mor-phologies growing within an austenite grain in order to minimizestrain energy associated with the shape and the volume change[24]. This is schematically represented in Fig. 6. The evolution ofthe different transformation temperatures, especially the Ms, hasbeen usually interpreted in terms of grain boundary strengtheningof the austenitic phase or increase in the non-chemical free energyopposing the transformation, when grain size reduces [25,26].Brofman and Ansell [27] suggested that the Hall–Petch strengthen-ing of the austenite affects the nucleation of martensite, by provid-ing a greater resistance to the motion of dislocations involved in thenucleation process [28]. As shown in the schematic in Fig. 7, the dif-ference in the free energies between austenite and martensite is thecritical driving force for the transformation to initiate [29]. Thechange in free energy associated with martensitic transformationis given by

DGA!Mnet ¼ DGA!M

ch þ DGA!Mnon ch ð5Þ

Fig. 8. (a) Inverse pole figure (IPF) and band contrast (BC)-grain boundary (GB) maps obtained through high resolution (HR) EBSD scans for sample A. The voltage wasmaintained at 10 k eV and a scan step size of 0.2 lm was used. The boundary misorientation are represented on BC map in different colors. The corresponding BC map showsincreased dislocation structures represented by high concentration of low angle misorientation (red); (b) difference in orientation spread of grains 1 and 2 with near ND//h110i orientation seen in the (110) pole figure plot. A significant spread of orientation in grain 2 is clear due to increased strain localization; (c) misorientation distribution ofgrains 1 and 2. Though average in-grain misorientation were identical, presence of dense dislocation boundaries (strain localization) in grain 2 caused a shift in themisorientation distribution to the right. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

R. Basu et al. / Journal of Alloys and Compounds 645 (2015) 213–222 219

Fig. 9. Enthalpy of transformation as a function of low angle misorientation forsamples A, B and S.

Fig. 10. Strain localization was sensitive to certain grain orientation. (a) Grains withND//h110i show a wide spread in the orientation as clear in the IPF triangles. (b) Thecorresponding point-to-origin misorientation profile plot for two ND//h110i grainsshows sharp spikes at a regular intervals indicating large misorientation changes atdislocation boundaries.

220 R. Basu et al. / Journal of Alloys and Compounds 645 (2015) 213–222

where DGA!Mch is the difference in chemical free energy proportional

to the amount of masrtensite formed and DGA!Mnon ch is the nonchemical

free energy that tends to oppose the transformation (consisting ofelastic strain). For a given microstructural condition, the onset oftransformation occurs when the DGA!M

net for the system takes a min-imum (maximum negative), if the temperature is decreased. As dis-cussed before, fine grain microstructures offer greater resistance tothe dislocation mobility for nucleation and propagation of the phasefront leading to a greater difference in the free energies of the twophases resulting in a shift of the Ms temperature to the left corre-sponding to minimum free energy difference.

During reverse transformation of martensite ? austenite, thedifferent martensite variants transforms to the parent phase in itsinitial orientation. The internal elastic energy stored in the selfaccommodating thermally induced martensite variants becomesthe driving force for the reverse transformation. The conclusionderived from few earlier studies [30–32] is that the grain size doesnot affect As of the reverse martensite transformation, however thepresented results show a pattern; a decrease in As temperature withincrease in the AGS. What could be speculated from the trend isthat the relatively high elastic strain in the martensite phase dueto misfits between different variants of martensite arising fromthe fine grained austenite microstructure (B) limits the reversemartensite to austenite transformation by introducing additionalinternal stress at the phase interface resulting in higher As.

The net molar enthalpy of transformation, DhA!Mnet corresponding

to the enthalpy of transformation in calorimetry upon transforma-tion from the parent austenitic phase A to the martensitic phase M(forward transformation), is governed by the enthalpy balanceequation [33]:

�DhA!Mnet ¼ �DhA!M

ch þ DhA!Mel þ DhA!M

d ð6Þ

where DhA!Mch is the chemical molar enthalpy that arise as a result of

the difference in entropy between the parent austenite and marten-sitic crystal structures, represents the total energy available for the

transformation. However, part of this molar enthalpy, DhA!Mel , is

stored as elastic energy associated to grain boundaries and the other

part, DhA!Md takes into account the defects or dislocations in the

crystals. The reverse transformation (martensite ? austenite) isendothermic and the heat balance Eq. (6) is modified to

DhM!Anet ¼ DhM!A

ch � DhM!Ael þ DhM!A

d ð7Þ

The energy components corresponding to defects during bothforward and reverse transformation are positive, therefore an

increase of defect enthalpy reduces Dhnet , upon forward transfor-mation but increases it upon reverse transformation. Assuming thatthe defect enthalpy scales with the average in-grain misorientation,one can expect a relationship of the measured enthalpy with themisorientation as represented in Fig. 5b. On the other hand,increase in the elastic stored energy, Dhel alone will reduce thenet enthalpy during both forward and reverse phase transforma-tion. However, the increasing nature of the forward transforma-tional enthalpy in the present samples suggests defect energy tobe much higher than the elastic stored energy. The contribution

of first term DhM!Ach is assumed to remain constant. During the

reverse transformation however, the defect energy contribution isnot pronounced, so the drop in the enthalpy is purely due toincrease in the elastic stored energy.

R. Basu et al. / Journal of Alloys and Compounds 645 (2015) 213–222 221

The present study also reveals some key microstructural featuresuch as dislocation banded structures present in few differentgrains that largely dictate the enthalpy of transformation. Suchdeformation heterogeneities (banded structures) in the grainsincrease with the degree of strain. A high resolution EBSD scan(step size of 0.2 lm) was obtained for sample A to bring out suchcharacteristic features. The degree of banding appears to vary fromgrain to grain. As shown in Fig. 8a, such dislocations bands are char-acterized by misorientation arrays with misorientation anglebetween 1 and 4 deg seen as low angle misorientation (red) inthe combined band contrast (BC) map. Regions consisting of tan-gled arrays of dislocations correspond to diffused Kikuchi patternsand could not be indexed. So those regions appeared black as rep-resented in the BC map. It is worth mentioning that determiningthe extent of rotations across these banded structures were notpossible at that point using measured EBSD data, only a smallregime that underwent recovery or partial recrystallization couldbe mapped with an acceptable mean angular deviation (MAD)value. Such dislocation arrays could result from inhomogeneousinternal stress fields introduced during deformation of polycrys-talline Ni–Ti alloys and are observed as strain localization. Thegrains labeled 1 and 2 in Fig. 8b, have nearly identical orientationand the same average misorientation. However the presence ofstrain localization in grain 2 has increased the orientation spreadand widened the distribution of misorientation presented inFig. 8c. The increased strain localization can substantially reducethe free energy of a martensite nucleus; correspondingly the driv-ing force or the enthalpy of transformation for martensitic transfor-mation is reduced. Fig. 9 presents the enthalpy of transformation asa function of fraction of low angle misorientation that correspondsto these dislocation structures. It clearly indicates that withincrease in the strain localization (low angle boundary fractions)the enthalpy of transformation for martensite formation decreases.It is therefore not the misorientation or dislocations in themicrostructure alone but the distribution of dislocations that deci-des the enthalpy of transformation. Furthermore, strain localiza-tions were sensitive to certain orientations; this is presented inFig. 10. In general, ND//h110i had the highest presence of strainlocalization that tend to change the orientation spread within thegrains. It could be therefore speculated that significant texturingin austenite (preferably ND//h110i) can substantially cause lower-ing of enthalpy of transformation in these alloys. This measurementhas been performed for few different grains of known orientations.This constitutes the prime novelty of this work.

Though a complete understanding on the pattern of change inenthalpy of transformation with increasing misorientation or elasticstored energy is speculative; the experimental observations are stillstatistically reliable and appears to be the only plausible explanationavailable at this point. A calculation of the defect energy from themeasured misorientation data would have indeed been interesting.

5. Conclusions

In this work, three different room temperature austenitemicrostructures were generated from samples of Ni47–Ti50–Fe3(in at.%) alloy processed through cold rolling and complex maform-ing routes. The idea was to carry out a systematic study on thetransformation behavior during phase transformation(austenite M martensite) in relation to microstructure and texture.The following conclusions have been gathered after EBSD andcalorimetric measurements.

(1) The transformation temperatures were purely function of thegrain size. The lowering of Ms with decreasing grain size hasbeen attributed to increase in the non-chemical free energy

or grain boundary strengthening opposing the transforma-tion. Conversely lowering of As temperature with increasein the austenite grain size was due to the high dislocationdensity in the martensite phase that limits the reversemartensite to austenite transformation by imposing frictionstress at the phase boundaries.

(2) The enthalpy of martensitic transformation increased withelastic stored energy and decreased with misorientation.The reverse transformational enthalpy showed the oppositetrend. Although increase in the stored energy alone corre-sponds to decrease in the martensitic transformationenthalpy, this trend was suppressed by higher defect energywhich increases the transformation enthalpy.

(3) Lastly, strain localization seemed to play an important role inlowering of martensite transformation enthalpy. These areseen as dislocation bands that tend to minimize the transfor-mation driving force during martensite transformation. Thedeciding factor for enthalpy is not the misorientation alonebut distribution of misorientation is important.

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