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Intermetallics 42 (2013) 1e8

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Intermetallics

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Evolution of microstructure and texture in Ni49.4Ti38.6Hf12 shapememory alloy during hot rolling

K.S. Suresh a, Dong -Ik Kim b, S.K. Bhaumik c, Satyam Suwas a,*

aDepartment of Materials Engineering, Indian Institute of Science, Bangalore, Karnataka 560 012, IndiabHigh Temperature Energy Materials Research Center, Korea Institute of Science and Technology, Seoul 136-791, Republic of KoreacMaterials Science Division, Council of Scientific and Industrial Research (CSIR), National Aerospace Laboratory, Bangalore 560 017, India

a r t i c l e i n f o

Article history:Received 30 August 2012Accepted 14 April 2013Available online 30 May 2013

Keywords:B. TextureB. TwinningC. RollingD. MicrostructureD. Grain-boundary character distribution

* Corresponding author. Tel.: þ91 80 22933245; faxE-mail addresses: satyamsuwas@gmail.com, satyam

(S. Suwas).

0966-9795/$ e see front matter � 2013 Elsevier Ltd.http://dx.doi.org/10.1016/j.intermet.2013.04.009

a b s t r a c t

This paper deals with the evolution of microstructure and texture during hot rolling of hafnium con-taining NiTi based shape memory alloy Ni49.4Ti38.6Hf12. The formation of the R-phase has been associatedwith the precipitation of (Ti,Hf)2Ni phase. The crystallographic texture of the parent phase B2 as well asthe product phases R and B190 have been determined. It has been found that the variant selection duringthe B2 / R phase transformation is quite strong compared to the case of the B2 / B190 transformation.During deformation, the texture of the austenite phase evolves with strong Goss and Bs components.After transformation to martensitic structure, it gives rise to a [011]jjRD fiber. Microstructure and texturestudies reveal the occurrence of partial dynamic recrystallization during hot rolling. Large strainheterogeneities that occur surrounding (Ti,Hf)2Ni precipitates are relieved through extended dynamicrecovery instead of particle stimulated nucleation.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

NiTi based shapememoryalloys (SMAs)find their applications inmedical and engineering fields because of two unique characteris-tics, namely, shape memory effect (SME) and super-elasticity (SE),also known as pseudoelasticity. The origin of these effects is due tothe reversible crystallographic transformation between the hightemperature austenite phase B2 (space group: Pm3m) and the lowtemperature martensite phase B190 (space group: P21/m) [1e5].Owing to the B2 4 B190 transformation temperatures the applica-tions of binary equi-atomic NiTi are limited to w100 �C. However,many engineering applications of SMA, particularly in aircraft andautomobile engines aswell as sensors and actuators inpowerplants,require the shape memory effect (SME) at temperatures beyond150 �C. This has necessitated the development of high temperatureshape memory alloys (HTSMAs) [6,7]. Addition of ternary alloyingelements such asHf, Zr, Nb, Pd, Pt, andAuhas been found to increasethe transformation temperature (TT) in these alloys [7e18].Amongst these elements, the addition of Hf to NiTi is attractivebecause of its relatively lower cost and higher transformationtemperature. Moderate levels of Hf addition (8e15 at pct.) can raise

: þ91 80 23600472.suwas@materials.iisc.ernet.in

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the TTofNiTi alloys considerably [12]. Further, the addition ofHf alsoincreases the transformation hysteresis up tow100 �C fromw40 �Cin binary NiTi alloys [12e14]. Another characteristic structuralchange in Hf based NiTi alloy is the evolution of R phase (Spacegroup: P3) during martensitic transformation at the coherent in-terfaces between thematrix phase B2 andprecipitate (Ti,Hf)2Ni [15].In this context, the evolution of texture in R phase vis-à-vis B190

phase is worth investigating, in order to understand the variantselection during different martensitic transformation.

The addition of Hf to NiTi alloys not only affects the phasetransformation but also alters the deformation characteristics ofthe alloy significantly. The deformation behavior is influenced bysolid solution strengthening as well as by the presence of secondphases in the microstructure. It is well known [16e19] that themicrostructural features and the crystallographic texture evolvedduring the processing of NiTi alloys have a significant influence onthe transformation behavior as well as on the recoverable strain.However, only a few reports [20e22] are available in which theeffect of addition of Hf to NiTi alloys has been studied in relation tothe evolution of microstructure and texture during hot working.

In the present study, an attempt has been made to examine theevolution of microstructure and texture in both the austenite andmartensite phases of a Hf based NiTi alloy. The role of the R-phase onthe evolution of texture during hot rolling has been particularlyemphasized.Analloywith thenominal compositionofNi49.4Ti38.6Hf12has been chosen to address the issues mentioned above.

Fig. 1. X-ray diffraction patterns for (a) as-cast, (b) solutionized, (c) annealed, and (def) hot rolled conditions; (d) 60 pct., (e) 80 pct. and (f) 90 pct. thickness reduction.

K.S. Suresh et al. / Intermetallics 42 (2013) 1e82

2. Experimental procedure

2.1. Material and process

An alloy with nominal composition Ni49.4Ti38.6Hf12 was pre-pared by vacuum arc melting using a non-consumable tungstenelectrode. The as-cast alloy was solutionized in a vacuum sealedquartz tube at 1000 �C for 24 h. The solutionized alloy was hotrolled at 900 �C to the thickness reductions 60, 80 and 90 pct. Thetransformation temperatures were determined after differentstages of processing using a Mettler Toledo Differential ScanningCalorimetry (DSC) at a heating/cooling rate of 10 �C/min.

2.2. Characterization methods

Phase analyses were performed by X-ray diffraction method. APANalytical XPert Pro system with Cu-Ka was employed for thispurpose. The rolled samples were subjected to textural andmicrostructural investigations. The deformation texture of themartensite phase was measured by X-ray diffraction method usinga Bruker D8-Discover X-ray texture goniometer based Schultzreflection geometry.

Microtexture was measured on electropolished samples using aHitachi S-4300SE FEG-SEM attached with an HKL Channel 5 EBSDsystem. All the EBSD measurements in the austenite phase werecarried out at 280 �C using a Kammrath & Weiss GmbH heatingstage. The microscope was operated at 30 KV and a beam current of4 nAwas used to obtain good quality patterns at high temperature.The EBSD data were acquired with a step size of 2 mm. In certaincases, a smaller step size of 0.3 mm was used to capture the finerdetails. For texture analysis, a line scan was performed, and thescanned area included almost 2000 grains. The measured orienta-tions including the texturewere calculatedusingREDS software [23]as well as TSL OIM Analysis 5. The grain boundaries were charac-terized as very low angle boundaries (VLAB) with misorientationrange 2e5�, low angle boundaries (LAB) with misorientation range5e15�, high angle boundaries (HAB) with misorientation greaterthan 15�, and Coincidence site lattice (CSL) boundaries having aspecific axiseangle pair. The grain sizewas estimated using the areafraction method, inwhich the diameter with an equivalent area of acircle is reported. To estimate the amount of strain levels withindifferent orientations, microstructural parameters like GrainOrientation Spread (GOS) and Kernel AverageMisorientation (KAM)were determined. GOS was calculated by averaging out the differ-ences between the average grain orientation and the orientation ateach point. For calculating KAM, a kernel is defined and the averagemisorientation between the center of a kernel and all the points atthe perimeter of the kernel are determined. The first nearestneighbor orientations were considered.

Transmission electron microscopic (TEM) study was carried outon a few selected samples using FEI F-30 FEG-TEM. The samples forthe TEM study were prepared by twin jet electro-polisher using asolution consisting of 5 pct. HNO3 and 95 pct. Ethanol at 20 V inwhich the temperature was maintained at �40 �C.

3. Results

3.1. X-ray diffraction studies

Fig. 1 (aec) show the X-ray diffraction patterns of the hot rolledsamples after 60 cpt., 80 pct., and 90 pct., thickenss reduction. TheXRD patterns of the hot rolled samples reveal the presence of thematrix phase B190, the R-phase, and the (Ti,Hf)2Ni precipitates. Hotrolling results in the broadening of the peaks corresponding to B190

and R phases with an increase in the value of full width half

maximum (FWHM), while that of (Ti,Hf)2Ni phase remains almostunaltered. The integrated intensity of the peaks corresponding tothe (Ti,Hf)2Ni phase increases with the rolling reduction.

3.2. Differential scanning calorimetry

Fig. 2 shows the DSC thermograms for the solution treated andhot rolled samples. The transformation temperatures for the sol-utionized and hot rolled samples obtained from the DSC analysisare presented in Table 1. It can be seen that although the Ms tem-perature does not change, the Mf, As, Af temperatures decrease inthe hot rolled material. The transformation hysteresis decreases onhot rolling. Another important observation from the DSC curves isthat in the solutionized sample, both the forward and the reversephase transformations B2 4 B190 are a single-step process. How-ever, in the hot rolled samples, this transformation is a two-stepprocess. Moreover, the phase transformation in the hot rolledsamples is sluggish compared to the solutionized sample.

3.3. Microstructural features

3.3.1. Microstructure of the room temperature martensite phase(B190) of the hot rolled material

Fig. 3 (a, b) shows the SEM micrographs of some selected hotrolled samples. At lower rolling reduction, the grain boundary(Ti,Hf)2Ni precipitates fragments at several regions. During thecourse of hot rolling, the precipitate size decreaseswith deformationand it is no longer populated along the grain boundaries; rather theyare uniformly distributed both at the grain boundary aswell as in thegrain interior region (Fig. 3a). In addition to the precipitate, the

(a)

(b)

Fig. 2. Differential scanning calorimetric (DSC) curves for (a) solutionized sample, (b)hot rolled sample exhibiting the transformation hysteresis.

K.S. Suresh et al. / Intermetallics 42 (2013) 1e8 3

matrix microstructure gets modified and results in the generation ofequiaxed grains (Fig. 3b). Hot rolling also results in the refinement ofthe martensite phase in the microstructure. Although the thick-nesses of the twins remain the same, their aspect ratios change.

The TEM micrographs of a few selected rolled samples areshown in Fig. 3(cee). After 10 pct. rolling reduction, predominantdislocation activity is observed in one of the twin variants (Fig. 3 c),the dislocation activity in other twin variants is almost negligible. Ascanning transmission electron micrograph of the region displayedin Fig. 3 (d) clearly depicts differential dislocation activity amongstthe twinned variants. As the amount of deformation increases, thetwin length reduces and leads to a large gradient in orientation ofthe martensite twins within a prior austenitic grain (Fig. 3e). As aresult, finer martensite laths lead to formation of moiré fringes.

3.3.2. Microstructure of the austenite phase (B2) of the hot rolledmaterial

The microstructural features of the austenite phase after 60, 80,and 90 pct. deformation levels have been recorded above the Aftemperature by electron back scattered diffraction with a heating

Table 1Transformation characteristics of processed alloys derived from the DSC thermo-gram (Temperatures given in �C).

Transformation temperature Ms Mf As Af

Solution treated 200 160 224 305Hot rolled 90 pct. 200 115 170 250

stage in SEM. The microstructural features are represented usingEBSD generated maps, namely, the inverse pole figure (IPF) map,image quality (IQ) map with a superimposed grain boundary andthe GOS maps (Fig. 4). The micrographs clearly reveal that above 60pct. deformation, no noticeable grain size reduction occurs in thematerial. It is to be mentioned that the grain size of the startingmaterial is w500 mm. The GB map indicates an increase in thefraction of VLAB and CSL boundaries with an increase in the degreeof hot deformation. Most of the CSL boundaries observed in themicrostructure are characterized by S3, S9, S13, S21 and S33.Majority of the grains show high fraction of VLAB and a few grainsalso appear without any VLAB. The intra-grain misorientation isrepresented through GOS map in Fig. 4 (c). With an increase indeformation, an increase in the average GOS value and a corre-sponding shift of the GOS distribution to the higher angle side couldbe seen. After 90 pct. deformation, fine equiaxed grains with lowerintra-grain misorientations are observed at the prior grainboundaries as well as in the grain interiors.

The grain size distributions and grain aspect ratio of the hightemperature austenite phase are presented in Fig. 5 (aec) as a func-tion of rolling reduction. The grain size distribution maps indicategrain refinement during hot rolling, however, the extent of refine-ment is limited. The distribution becomes more homogeneous andthe standard deviation is reduced with an increase in deformation.The grain boundary character distribution given in Fig. 5 (d) shows ahigher fraction of LAB compared to HAB at all rolling reductions. Thedistribution of LAB and HAB are similar at all rolling reductions. TheKAM distributionwidens and a shift of KAM peak to a larger angle isobservedwith an increase in the amountof hot deformation (Fig. 5 e).

3.4. Texture evolution

3.4.1. Texture of room temperature martensite (B190) phaseThe experimental (001) and (101) pole figures for the B190 phase

and the (300) pole figures for the R-phase, as measured using X-raydiffraction, are shown in Fig. 6. The R-phase shows a strongertexture development than the B190 phase.

For the R-phase, till 80 pct. deformation, a weak <001>jjNDfiber is observed. With an increase in strain, this fiber strengthens.However, the location of the maxima continuously changes. Themaximum strength is observed along the <104>jjND fiber axis. Inbetween these two fibers, certain components like (104)[421] and(124)½611�, which are in the vicinity of <104> fiber, exhibit strongintensity.

For the B190 martensite phase, a (011) fiber is observed for all thedeformation conditions, and it shows a spread up to (013). After 80pct. deformation, the (011) fiber exhibits the maximum strengthbeyond which it weakens. However, compared to the R-phase, thestrength of (011) fiber in the B190 phase is insignificant.

3.4.2. Texture of high temperature austenite (B2) phaseThe (100) and (110) pole figures of the high temperature

austenite phase are shown in Fig. 7. Aweak {111} fiber is observed inall the rolled samples. Albeit aweak intensity, this fiber strengthenswith deformation. The main texture components observed in thetexture of the austenite phase are {111}<112> (Cu), {123}<634> (S),{112}<110> and {110}<001>. These components are present in allthe rolled samples. The locations of some important texture com-ponents aremarked in the (100) and (110) pole figures (Fig.7). A fewadditional texture components like {110}<112> (Bs), {112}<111>and {013}<100> (rotated cube) with strong intensity are alsoobserved in the samples deformed to 90 pct. reduction.

During the initial stages of deformation, the {111}<112> and{112}<111> components are located in their ideal positions withthe highest intensity. As the deformation proceeds, the spread of

Fig. 3. SEM micrographs of samples rolled to (a) 10 pct. and (b) 80 pct. rolling reduction. (c) TEM and (d) STEM micrograph of the sample rolled to 10 pct reduction. (e) TEMmicrograph of sample rolled to 60 pct. reduction. (Moiré fringes after 60 pct. reduction indicate a high dislocation activity).

K.S. Suresh et al. / Intermetallics 42 (2013) 1e84

these components increases resulting in an overlap after 90 pct.deformation. To have a better quantitative estimation of evolutiontexture, the volume fractions of important texture componentshave been plotted as a function of rolling reduction (Fig. 8). It can beseen that with the increase in deformation, the volume fractions of{111}<112>, {112}<110> and {110}<112> components increases

while the strength of {123}<634> and {112}<111> components getsaturated. The rate of evolution of texture is, however, not the samefor all the texture components. The components {110}<001> and{111}<112> show either a linear increase or decrease in volumefraction, however, in the case of {110}<112>, {112}<110> and{123}<634>, the rate of change is abrupt.

Fig. 4. Inverse pole figure maps, image quality plus grain boundary maps and grain orientation spread maps of the austenite phase of the samples rolled to (a) 60 pct., (b) 80 pct.,and (c) 90 pct. reduction (for the colours associated with the figures, please refer to the online version of the paper).

K.S. Suresh et al. / Intermetallics 42 (2013) 1e8 5

4. Discussion

The salient features pertaining to microstructural evolutionduring various stages of thermo-mechanical processing are asfollows: (i) microstructural modification during hot rolling, and (ii)the evolution of the R phase in the microstructure. These aspectswill be discussed in the following subsections.

4.1. Microstructural modification during hot rolling

The microstructures presented in Figs. 4 and 5 indicate a higherfraction of VLABs than the LABs for all the deformation levels.VLABs can form either during deformation or annealing. Thegeometrically necessary dislocations generated during deforma-tion, to maintain the strain compatibility, could manifest in theform of VLAB. On the other hand, the dislocation annihilationduring recovery could also give rise to the formation of VLAB. In thecase of deformation, the increase in the fraction of VLAB should beassociated with a concomitant increase in LAB. However, theanalysis of EBSD scans clearly reveal that increase in the fraction ofLAB is not so significant as that of the VLABs. Therefore, it could beinferred that the large increase in VLAB is mostly associated with

recovery. The microstructural features also indicate that there is aconsistent increase in the fraction of HAB, which is a clear indica-tion of the occurrence and progress of dynamic recrystallization. Anincrease in misorientation values displayed in the GOS map (Fig. 4)corroborates the increase in lattice strain with rolling reduction. Inspite of the occurrence of dynamic recrystallization, majority of thegrains still possess deformation characteristics.

The absence of completely coherent (Ti,Hf)2Ni precipitates afterrolling indicates that they are metastable and can induce a largercoherency strain. Most of the plate like structures get distorted andare converted to lenticular and spherical (Ti,Hf)2Ni precipitates. Forlenticular precipitates, only the ends are incoherent. Partial co-herency is maintained along the length. Although coherent(Ti,Hf)2Ni precipitates are absent after hot rolling, the partlycoherent (Ti,Hf)2Ni precipitates are still observed, which could actas a nucleation site for the R phase.

4.2. Development of texture

4.2.1. Texture of the austenite phaseThe hot rolling of the Ni49.4Ti38.6Hf12 alloy resulted in high in-

tensity Bs and Goss components with a large spread between the

(a) (b)

(c) (d)

(e)

Fig. 5. Grain size distribution for the austenite phase of (a) 60 pct. (b) 80 pct. and (c) 90 pct. rolled samples. (d) grain boundary character distribution and (e) Kernal averagemisorientation distribution for the rolled samples (for the colours associated with this figure, please refer to the online version of the paper).

K.S. Suresh et al. / Intermetallics 42 (2013) 1e86

components. Amongst the various texture components, Bs, S androtated Cu orientations show large intra-grain misorientation(Fig. 5). Goss, Cube and rotated Cube orientations develop smallintra-grain misorientation. This suggests that the former compo-nents are due to deformation whereas the later ones are related todynamic recrystallization. The presence of a weak [111]jjRD fiberand the presence of strong Bs, Goss and rotated Cu components

indicate that the addition of Hf can significantly modify the hotrolling texture. This observation is quite new compared to theearlier findings in binary NiTi alloys. Inoue et al. [16] have reportedthe evolution of texture componentsð332Þ½110� and ð211Þ½111�during rolling and ð111Þ½110� and (110)[001] during annealing inthe NiTi alloy without Hf. In the present case, these components arequite weak.

Fig. 6. (001) and (101) pole figures of B190 phase and (300) pole figures of R phase of the martensite for 60 pct., 80 pct. and 90 pct. reduction (for the colours associated with thisfigure, please refer to the online version of the paper).

K.S. Suresh et al. / Intermetallics 42 (2013) 1e8 7

4.2.2. Texture of the martensite phaseFormation of R phase has been associated with the different

precipitates in NiTi based alloys [19,24e26]. In the present study,evolution of the R-phase has been associated with (Ti,Hf)2Ni

Fig. 7. (100) and (110) pole figures of austenite phase of the samples rolled to (a) 60 pct.superimposed on (100) and (110) pole figures for the austenite phase (for the colours asso

precipitates. The high intensity (300) pole figures of the R phasesuggest that variant selection during the martensitic trans-formation to R phase is very strong, whereas for the case of B190

there is no particular preference for variant selection. The (002)

, (b) 80 pct. and (c) 90 pct. reduction. (d) Location of important texture componentsciated with this figure, please refer to the online version of the paper).

{110}<001>{111}<112>{112}<111>{123}<634>{112}<110>{110}<112>0

1

2

3

4

5

6

7V

olum

e Fr

actio

n

Texture Components

60 pct. 70 pct. 80 pct.

Fig. 8. Volume fractions of important texture components in austenite phase of therolled samples.

K.S. Suresh et al. / Intermetallics 42 (2013) 1e88

pole figure of martensite and (110) pole figure of austenite corre-sponding to 90 pct. rolling reduction represent some texture cor-respondence. The texture component (001)[100] of martensite isrelated to the component (110)[001] of the austenite. Earlier studies[27,28] on binary NiTi have reported such correlation between theaustenite and martensite texture components. They also correlatedthe texture components and the types of twins.

5. Conclusion

In the present investigation, the evolution of microstructure andtexture in the austenite and martensite phases of the shapememory alloy Ni49.4Ti38.6Hf12 has been investigated with a partic-ular emphasis on the R phase. The experimental results have led tothe following conclusions:

(i) The microstructure of the martensite phase at lower rollingreduction reveals preferential dislocation activity amongstdifferent twin variants.

(ii) The elongated grain morphologies of the as-cast microstruc-ture transform to a homogeneous and equi-axed microstruc-ture after 90 pct. rolling reduction. Large strain heterogeneitiesdevelop surrounding the (Ti,Hf)2Ni precipitates, which relaxthrough extended dynamic recovery during furtherdeformation.

(iii) The observations pertaining to microstructure and textureevolution indicate the occurrence of partial dynamicrecrystallization.

(iv) Hot rolling of the experimental alloy results in a texture withstrong {111}<112>, {123}<634> and {110}<001> components.

(v) Transformation to R phase is associated with planar (Ti,Hf)2Niprecipitates. Further, the R phase also reduces thermal hys-teresis of transformation. Strong texture in R phase confirms apreferential variant selection during martensitictransformation.

(vi) Texture correspondence has been observed between theaustenite and martensite phases.

Acknowledgment

Use of Central facility for X-ray and Advanced facility formicroscopy and microanalysis (AFMM) at IISc are acknowledged.

One of the authors KSS would like to acknowledge the support thathas been provided by Korea Institute of Science and Technology,Seoul. (Grant No. 2E22131).

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