In situ investigation of the fast microstructure evolution...

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In situ investigation of the fast microstructure evolutionduring electropulse treatment of cold drawn NiTi wires

B. Malard a,d, J. Pilch a, P. Sittner a,⇑, R. Delville b, C. Curfs c

a Institute of Physics, Na Slovance 2, 182 21 Praha, Czech Republicb EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium

c ESRF, 6 rue Jules Horowitz, 38043 Grenoble, Franced SIMaP, Domaine Universitaire, BP 75 38402 Saint Martin d’Hères, France

Received 9 August 2010; received in revised form 5 November 2010; accepted 6 November 2010Available online 10 December 2010

Abstract

Microstructural changes taking place during the heat treatment of cold-worked NiTi alloy are of key interest in shape memory alloytechnology, since they are responsible for setting the austenite shape and functional properties of the heat-treated alloy. In this work,microstructural evolution during non-conventional electropulse heat treatment of thin NiTi filaments was investigated in a uniquehigh-speed in situ synchrotron X-ray diffraction experiment with simultaneous evaluation of the tensile force and electrical resistivityof the treated wire. The in situ results provide direct experimental evidence on the evolution of the internal stress and density of defectsduring fast heating from 20 �C to �700 �C. This evidence is used to characterize a sequence of dynamic recovery and recrystallizationprocesses responsible for the microstructure and superelastic functional property changes during the electropulse treatments.� 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Shape memory alloys; Recovery; Recrystallization; Joule heating; Superelasticity

1. Introduction

Martensitically transforming shape memory alloys(SMA) such as the near-equiatomic NiTi alloy [1] haveattracted considerable attention for their unique functionalthermomechanical properties, such as superelasticity orshape memory effect. Conventionally, NiTi wires are pro-duced from cast ingots by extrusion into bars, followedby multiple hot-working passes finished by a final colddrawing into thin wires. In this so-called “cold-worked”(as-drawn, hard, etc.) state, the alloy possesses a heavilydeformed microstructure consisting of a mixture of austen-ite, martensite and amorphous phases with defects andinternal stress [2–12]. In the cold-worked state, the wiresdo not show any functional properties, since the martens-itic transformation is suppressed. Cold-worked NiTi wiresneed to be heat treated so that their heavily deformed

microstructure changes into an annealed microstructurein which the wire shows the desired functional properties.At the same time, if the shape of a SMA element is con-strained during this final heat treatment, it acquires thisshape as a new “parent austenite shape”. The final thermo-mechanical treatment thus has two goals: setting the func-tional properties and setting the shape of the NiTi wire.

The final heat treatment of NiTi is usually performed inenvironmental furnaces [13–17]. The effect of the final heattreatment temperature and time on the microstructures,martensitic transformations and mechanical properties ofNiTi is now well documented [1,14] and is routinely usedto manufacture NiTi wires with optimized mechanicaland functional properties [3,15]. Depending on the alloy(chemical composition, cold work) and desired application,the typically used heat treatment conditions vary in thetemperature range 400–600 �C and time 10–60 min[14,15]. In addition to microstructure recovery, diffusion-driven processes such as precipitation or dissolution of pre-cipitates [13] affecting the transformation temperatures and

1359-6454/$36.00 � 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.doi:10.1016/j.actamat.2010.11.018

⇑ Corresponding author. Tel.: +42 0266052 657; fax: +42 02286 890 527.E-mail address: [email protected] (P. Sittner).

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Acta Materialia 59 (2011) 1542–1556


functional properties may also take place, depending onthe alloy composition, temperature and time of treatment.If the shape of the wire is constrained during heat treat-ment, stress always appears temporarily [16,18]. In particu-lar, a relatively small constant tensile force is usuallyapplied during straight annealing treatment of superelasticNiTi wires [3]. Although this stress (20–100 MPa [15]) is animportant technological parameter of the NiTi wire pro-duction process, little is known about its effect on the shapeand functional properties of the wire [16].

Recently, several research groups reported [2–12] thatNiTi can be relatively easily brought into an amorphousstate by severe cold working and that subsequent heattreatment yields ultrafine-grained, i.e., nano- or sub-microcrystalline, microstructures which provide NiTi withexcellent mechanical and functional properties. This is alsothe case for thin superelastic NiTi filaments investigated inthis work [18].

The relatively long time of conventional heat treatmentof thin NiTi wires in an environmental furnace has become

Fig. 1. Microstructures and superelastic functional properties of thin NiTi wires evolving during sequential P = 125 W/1, 2, 3, . . . , 18 ms electropulse heattreatment (data taken from Ref. [21]). Microstructures found by TEM in 8, 12, 14, 18 ms treated wires (left) and cyclic tensile stress–strain curves (right);10 cycles in position control at room temperature.

B. Malard et al. / Acta Materialia 59 (2011) 1542–1556 1543


an obstacle for the emerging technology of hybrid textilefabrics knitted or woven using continuous NiTi filaments.The maximum respooling speed of �1 m min�1 achievablewith conventional straight annealing treatment in �6-m-long tubular electric furnaces is painfully slow for textileprocessing of NiTi filaments. To solve this technologicalproblem, a method called “Final ThermoMechanicalTreatment by Electric Current” (FTMT-EC) was devel-oped [19,20].

Fig. 1, adapted from Ref. [21], compares bright-fieldtransmission electron microscopy (TEM) images of micro-structures observed in the same NiTi wire treated by shortelectric current pulses with tensile superelastic stress–strainresponses of these wires in 10 tensile cycles at room temper-ature. Note that, following a single electrical power pulseof 125 W lasting for only 12 ms, the wire shows excellentsuperelastic functional properties, high strength and rela-tively stable cyclic superelastic stress–strain response atroom temperature (Fig. 1c). A dedicated TEM study ofmicrostructures in FTMT-EC-treated NiTi wires wasreported in Ref. [21].

The physical processes responsible for these fast changesin microstructure and functional properties during heatingwere investigated in this work, using high-speed synchro-tron X-ray diffraction with simultaneous tensile force andelectrical resistivity measurements. The ultimate aim wasto establish a knowledge base for the relation betweenthe parameters of the FTMT-EC treatment and to achievesuperelastic functional properties in the treated NiTi fila-ments [19].

2. Experiment and method

2.1. Heat treatment

The experiments were performed using as-drawn super-elastic NiTi wires (56.0 wt.% Ni) with diameter d = 0.1 mm(Fort Wayne Metals #1). The wire was first mounted at theminiature deformation rig especially designed and built forefficient testing of functional thermomechanical propertiesof thin NiTi filaments in tension. The rig consisted of astepping motor, a 10 N load cell, electrically isolated grips,a Peltier furnace to control the environmental temperatureand prevent air circulation, a laser micrometer for strainmeasurement and special electronics which enable a con-trolled electrical power pulse to be sent to heat the NiTiwire up to melting point and perform simultaneously elec-trical resistivity measurement.

n � CpdTdt¼ P � A � h � ðT � T extÞ � A � e � r � T 4 ð1Þ

Cp ¼ 24:647þ 0:002645T � 107:05T �2 ð2Þ

The initial length l0 (�50 mm) and electrical resistance q0of the wire at room temperature were first measured.The wire was then preloaded to the tensile stressr0 = 400 MPa, and its length was constrained. Followingthat, the heat treatment was performed by a single electri-

cal power pulse P(t) controlled in time. Fig. 2a shows theparameters describing the shape of the electrical powerpulse. Only the maximum power P_High normalized for100 mm long wire (henceforth noted P) and the pulse dura-tion t_2 will be used in the following to characterize thetreatments. Alternatively, each treatment can be character-ized by the temperature history T(t) calculated from heatEq. (1) which takes into account the Joule heat supplyP(t) and the ambient temperature losses.

The gram-atomic heat capacity Cp was taken from theliterature [23] in the form of Eq. (2); r is a Stefan–Boltz-man constant, n is the molar amount of atoms and A isthe surface area of the wire. The specific heat transfercoefficient h, describing the heat dissipation into air perunit time, and the NiTi wire emissivity e, needed forthe calculation of the radiation heat loss per unit time,were identified from a series of calibration experiments.Other effects influencing the wire temperature, such asheat conduction losses into grips and latent heats, wereneglected.

Fig. 2b shows the temperature of the wire evolving dur-ing heating by electrical pulses with parameters P = 2, 3, 4,5, 6 W, t_2 = 0.9 s and P = 10 W, t_2 = 0.18 s. Pulse dura-tions t_2 = 0.9 s and 0.18 s were selected owing to the lim-ited reading time of the X-ray detection. They are muchlonger than the pulse durations t_2 = 0.001–0.018 s usedin previous ex situ studies [18,21,22]. Nevertheless, it willbe shown below that one can obtain quite similar func-tional properties of the treated wires.

Fig. 2c shows the evolution in time of a set of parame-ters (temperature T, force F, electrical resistivity q, powerP, voltage V and current I) during electropulse treatmentperformed on an already heat-treated wire. Note that theapplied voltage V and current I (Fig. 2c) vary with the elec-trical resistivity of the wire, so that the power P remainsconstant during the t_2 period as prescribed. Fig. 2d showsthe electrical resistivity q and tensile stress r as functions oftemperature during the same treatment. As expected fortreatment on an annealed wire, the electrical resistivityincreases, and tensile stress decreases linearly with increas-ing temperature, and the responses are completely revers-ible upon cooling. In the case of the cold-worked wire,however, its microstructure evolves towards the annealedmicrostructure during treatment. This brings about addi-tional significant irreversible variation in the electricalresistivity and tensile stress, which superimpose on the lin-ear reversible changes observed in the case of the annealedwire.

In the experiments reported below, the electrical resistiv-ity of the wire, tensile force and X-ray diffraction signalvarying during the treatment were simultaneously evalu-ated during various FTMT-EC treatments. After the treat-ments were finished, the superelastic response of thealready treated wire was evaluated in 10 tensile cycles atroom temperature using the same stress rig. The electricalresistivity of the NiTi wire was also evaluated during tensiletesting. The recorded superelastic stress–strain curves,

1544 B. Malard et al. / Acta Materialia 59 (2011) 1542–1556


including their stability during cyclic loading, serve in thefollowing as measures of the achieved functional superelas-tic properties of the treated wires.

2.2. Synchrotron X-ray diffraction

Synchrotron radiation enables X-ray measurements toextend from the static to the dynamic regime, thanks toits unique time structure and very high flux, even at highenergy, which considerably reduces the data collectiontime. Very fast time resolution down to the picosecondregime is achievable with stroboscopic studies of reversiblephenomena [24]. In the case of the irreversible microstruc-ture evolution in metals, X-ray data have to be acquired ona single shot basis, requiring a high flux source, a very fast

and sensitive data acquisition system and high-speed datatransfer. This has recently become possible with hard syn-chrotron X-ray sources and the development of the latestX-ray camera detector [25].

The experiment in this work was performed on the ID11diffractometer at the European Synchrotron RadiationFacility (ESRF) in Grenoble. A monochromatic X-raybeam with energy 45 keV (k � 0.27552 Å) was obtainedwith a Laue monochromator. The beam was focused downto 1 mm in the vertical direction by bending the Laue crys-tal, and down to 100 lm in the horizontal direction. AFreLoN2K two-dimensional (2D) camera with taper waschosen for the fast detection system [25]. In order to speedup data acquisition for the targeted 10 ms time resolution,a radial slice of the 2D detector (Fig. 3a) was used.

Fig. 2. FTMT-EC treatment by controlled electrical power pulse P(t): (a) parameters describing the pulse shape; (b) calculated temperature profilesT = T(t) achieved in pulses with parameters P = 2, 3, 4, 5, 6 W, t_2 = 0.9 s and P = 10 W, t_2 = 0.18 s; (c) power, voltage, current, temperature, force andelectrical resistivity varying during 6 W treatment of an already 6 W heat-treated NiTi wire; (d) reversible variation in electrical resistivity and tensile stresswith temperature during heat treatment described in (c).



The thin NiTi wire mounted between the grips of thestress rig was positioned vertically in the beam (Fig. 3a).The heat treatment experiment was performed with syn-chronized pulse, force, electrical resistivity and diffractionmeasurements (Fig. 4). A strip 2048 � 64 pixels wasacquired and transformed in order to obtain a line of2048 � 1 pixels. Then, this line was read out of the CCDand stored before the next image was acquired. Afterwards,300 lines were put together into a single 2D binary image,which represents the evolution of the diffraction patterns vstime, as illustrated in Fig. 3b. To minimize the time neededto read out the line down to 10 ms, the 64 active pixels hadto be chosen at the bottom edge of the camera. The part ofthe detection area, which was not active, was masked by alead mask. Using such a kinetic mode of data acquisition

by the FreLoN camera, it was possible to acquire 100 dif-fraction patterns per second with a readout time of


diffraction profile corresponding to a single {0 1 1} austen-itic peak of the fully treated wire (Fig. 5c). The recordeddiffraction patterns were fitted using one austenite {1 1 0}and three martensite peaks (Fig. 5b). The {1 1 0} austenitereflection was analysed for peak height, intensity, positionand width (full width at half maximum (FWHM)) usingLAMP software [28] as well as GSAS software [29] consid-ering the structure of the austenite phase Pm-3m, R phaseP-3 and martensitic phase P21/m [1].

3. Experimental results

3.1. Heat-treatment experiments

The heat treatment experiments were performed usingdifferent powers P = 2, 3, 4, 5, 6 W (t_2 = 0.9 s) and10 W (t_2 = 0.18 s). The same experiments were performed

once with simultaneous in situ X-ray diffraction measure-ment and several times without the X-rays. The reproduc-ibility is very good in the sense that the same functionalproperties of the wires were achieved in treatments per-formed with the same P, t_2 parameters. The in situ resultsare shown in Fig. 6 as the variations in time of the suppliedelectrical power (1, blue), tensile stress (2, black), tempera-ture (3, red) and electrical resistivity (4, green) during theFTMT-EC pulse on the left-hand side, and the obtainedsuperelastic stress–strain response with superimposed vari-ation in the electrical resistivity of the wire on the right-hand side. More detailed results of multiple FTMT-ECtreatments using P = 2–7 W/t_2 = 0.9 s can be found in arelated work [19].

The tensile stress and electrical resistivity responses dueto microstructure recovery are superimposed on the linearthermal dependences of the strain (and hence the tensile

Fig. 5. Quantitative phase analysis of the diffraction patterns yielding information on the evolution of the integrated intensity, height, position and width(FWHM) of the {1 1 0} austenitic peak (in red). Diffraction pattern of NiTi powder (a) and NiTi cold-worked wire before (b) and after (c) heat treatment.(For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)



Fig. 6. Evolution of electrical power, temperature, tensile stress and electrical resistivity of NiTi wires during FTMT-EC treatment with P = 2, 3, 4, 5,6 W/0.9 s and 10 W/0.18 s (left) and resulting superelastic response of the treated wires at room temperature with superimposed in situ electrical resistivityrecords (right).



stress due to the applied constraint) and electrical resistiv-ity (Fig. 2d). Obviously, this makes the interpretation ofthe in situ results very complicated. To allow for mutualcomparison of the results achieved in different treatments,tensile stress and electrical resistivity responses fromFig. 6 are plotted together in dependence on time and tem-perature in Fig. 7. The different colour1- and letter-codedcurves represent the variation in stress (upper graphs)and electrical resistivity (lower graphs) for different electri-cal power.

It is expected that the tensile stress evolves with time(temperature) owing to the several possible mutually com-peting deformation mechanisms, including: (i) reversetransformation of martensite and relaxation of residualstress in the as-drawn wire; (ii) thermal expansion/contraction; (iii) dynamic recrystallization and/or plasticdeformation processes at high temperature and stress;and (iv) martensitic phase transformations B2-R-B190.The tensile stress starts to increase with increasing temper-ature straight after the onset of heating (Fig. 6). This isattributed to the reverse transformation of the retainedmartensite coupled with the relaxation of residual stress(i). Note that mechanism (i) results in wire shortening evenif mechanism (ii) acting simultaneously tends to elongate

the wire. Since mechanism (i) prevails at low temperatures,the cold-worked wire shows negative thermal expansion.Consequently, the tensile stress increases with increasingtemperature up to �350 �C (depending on the used heatingrate), where it reaches a maximum of �650 MPa (Fig. 7a)and then falls to 0 MPa stress if the maximum temperaturereached is high enough. This dramatic change in tensilestress response with increasing temperature is attributedto the termination of the activity of mechanism (i) com-bined with tensile deformation associated with dynamicrecrystallization (iii) and thermal expansion (ii). Uponcooling, the stress varies, basically owing to conventionalthermal contraction (ii) and, possibly, stress-inducedB2-R-phase transformation


in the microstructure (Fig. 1). Upon cooling, the electricalresistivity decreases owing to conventional thermal depen-dence of the electrical resistance of a metal (Fig. 2d).Finally, as the temperature decreases to �50 �C, the electri-cal resistivity may increase again owing to the B2-R-phasetransformation (Fig. 7b, curve c), since the R-phase has ahigher electrical resistivity than the B2 phase [30]. In caseof fully treated wires, total change in the electrical resistiv-ity before treatment (1.45 X lm) and after treatment(0.88 X lm) measured at room temperature reaches �40%(Fig. 7, curves d–f). This change in the electrical resistivityis fully attributed to the microstructural changes caused bythe treatment.

3.2. Heat treatment experiments with in situ X-ray

diffraction

In contrast to the in situ tensile stress and electricalresistivity, the in situ synchrotron X-ray data reflectmainly the irreversible microstructure changes duringFTMT-EC treatments. An exception is the change in peakposition with temperature, which is partially due toreversible thermal expansion/contraction. The evolutionof the profile of the {1 1 0} austenitic peak (Fig. 5b andc) during heat treatment was analysed for integratedintensity, peak height, peak position and peak width(FWHM). Variation in these values with time (tempera-ture) during the electrical pulse treatment represents thein situ X-ray results. In order to present the completeresults obtained from the synchrotron experiment,Fig. 8a–d shows the evolution of the tensile stress(Fig. 8a and b) and electrical resistivity (Fig. 8c and d)evaluated simultaneously with the X-ray data. The resultsare essentially the same as those reported in Fig. 7. Smalldifferences come from different sample environments: noPeltier furnace and a vertical orientation of treated wirewere used in the in situ X-ray experiment.

Fig. 8e and f shows the evolution of the intensity ofthe {1 1 0} austenite peak (peak height) during heattreatments with P = 2, 3, 6 and 10 W pulses. The higherthe temperature reached in the treatment, the larger thepeak height after treatment. The change in peak heightis related mainly to the change in profile shape duringtreatment (Fig. 3b), since the integrated intensity(Fig. 8k and l) is changing marginally during heat treat-ment. The change in peak height by treatment is perma-nent and increases with maximum temperature reachedin the treatment. In the case of P = 10 W/0.18 s treat-ment, the peak height starts and stops increasing earlier,since the temperature rises faster in this shorter pulsetreatment (Fig. 2b).

Fig. 8g and h shows the evolution of {1 1 0} austenitepeak position recalculated into “radial strain”–lattice strainof the {1 1 0} austenitic peak (Eqs. (3)–(5)) for treatmentwith P = 2, 3, 6 and 10 W pulses. Radial strain refers tothe transverse component of the elastic strain in the familyof grains of the austenite phase with h1 1 0i direction ori-

ented perpendicularly to the wire axis. Using the Bragglaw (Eq. (3)) and the definition of lattice strain (Eq. (4)),an expression for the radial strain of the {1 1 0} austeniticpeak e110 is given by Eq. (5).

2dhkl � sin hhkl ¼ n � k ð3Þ

ehkl ¼dhkl � d0;hkl

d0;hkl¼ Ddhkl

d0;hklð4Þ

e110 ¼ 1�sin h110sin h0110

ð5Þ

h110 is the value of the Bragg angle of the {1 1 0} austeniticpeak in the wire during heat treatment, and h0,110 is the va-lue of the Bragg angle of the {1 1 0} austenitic peak mea-sured on the NiTi powder. Because the radial strain inthe wire before heat treatment is strongly negative, it meansthat the grains of the austenite phase with h1 1 0i directionoriented perpendicularly to the wire axis in the as-drawnNiTi wire experience a tensile stress along the wire axis.Recall, however, that there are other austenite grains aswell as the martensite and amorphous phases in the micro-structure of the as-drawn wire. Since the stress must bebalanced, the measured negative radial strain is taken asevidence for the existence of very strong internal inter-granular stresses in the microstructure of the cold-workedwire. Note also that before and after treatments at roomtemperature, the wire is under �400 MPa external tensilestress, which also affects the evaluated radial strain (Figs. 6and 7).

As the temperature increases during heat treatment,the austenite lattice expands in a radial direction, pre-sumably due to three processes: (i) isotropic thermalexpansion; (ii) directional decrease in internal tensilestress; and (iii) decrease in external tensile stress(Fig. 7a). The evolution of the radial strain with temper-ature due to processes (i) and (iii) is reversible. Neverthe-less, the radial strain after cooling to room temperaturenever returns to its original starting value after any ofthe studied treatments (Fig. 8h). The irreversible partof the change in radial strain is thus attributed to thedisappearance of the internal stress (ii) introduced duringthe cold drawing process, and the transformation of themartensite and amorphous phases to austenite before thetemperature reaches �600 �C (Fig. 7). Small internalstresses probably remained in the microstructure afterthe 2 W treatment and, to a lesser extent, after the3 W treatment.

Fig. 8i and j shows the evolution of the FWHM ofthe {1 1 0} austenitic peak for experiments with P = 2,3, 6 and 10 W pulses. In the case of the broadened pro-file of the cold-worked wire (Fig. 5b), the FWHM of the{1 1 0} austenitic peak was evaluated as 0.58�. Itdecreases with increasing temperature, mainly due tothe decrease in the density of defects and amorphousphase in the cold-worked microstructure and relatedstrain fields. Following the P = 6 W treatment, theFWHM falls to the saturated value of �0.1�. The



FWHM decreases to only 0.35� and 0.2� in the case ofthe 2 W and 3 W treatments, respectively, suggesting that

the microstructures of such treated wires still containmany crystal defects and related strain fields.

Fig. 8. In situ X-ray diffraction experiments – evolution of (a and b) tensile stress, (c and d) electrical resistivity, (e and f) intensity, (g and h) radial strain(peak position), (i and j) FWHM (peak width) and (k and l) integrated intensity of the austenite {1 1 0} reflection with time (left) and temperature (right)during heat treatments using P = 2, 3, 6 W/0.9 s and 10 W/0.18 s. The start and finish of the temperature rise are denoted by dashed lines.



4. Discussion

The results of the present in situ experiments supplementthe results of post heat treatment studies of the same NiTiwires performed earlier by ex situ synchrotron X-ray [18]and TEM [21,22] methods. Despite the different electricalpower and time used, there is little doubt that the evolutionof microstructure, texture and functional properties is simi-lar to those reported in the ex situ studies [18,21,22] (Fig. 1).Lower power P and longer pulse time t_2 were used in thepresent experiments to slow down the progress of micro-structure evolution during the FTMT-EC treatment so thatthe in situ X-ray measurements during the selected electro-pulse treatments could be performed with the targeted 10ms time resolution. More detailed evidence on the superelas-tic property setting of NiTi wires by the FTMT-EC methodusing electrical power varying in the range P = 2–7 W wasreported in a closely related paper [19].

4.1. Processes responsible for microstructure evolution

during FTMT-EC treatment

Recording the electrical resistance and force is now rou-tinely used in practical FTMT-EC treatments of NiTi fila-ments prior to their textile processing [19]. In thesetreatments, electrical power parameters and mechanicalconstraint are selected by the user to achieve the desiredsuperelastic properties of the treated filament. Therecorded evolution of electrical resistance and tensile stressduring treatment is saved to obtain the same functionalproperties repeatedly in future use. Although FTMT-ECtreatments can be performed without detailed understand-ing of the physical processes and microstructure evolutionresponsible for the recorded electrical resistance and tensilestress, the lack of this knowledge is a barrier for furtherdevelopment of the method. Direct structural and phase-specific information on the progress of the recovery pro-cesses during heating available from the present X-raystudies as well as related TEM studies [21] is hence of greatvalue in this respect. Table 1 summarizes the physical pro-cesses expected to proceed in the wire during electropulseheating, as discussed below.

Fig. 9 shows the in situ X-ray results, electrical resistivityand tensile force from the P = 10 W/0.18 s treatment plot-ted together in dependence on time (Fig. 9a) and tempera-ture (Fig. 9b). The tensile stress increases right from the

onset of heating, and the electrical resistivity starts to fallat 200 �C, before any significant change in the diffractionprofile can be detected. Such changes, however, might bebelow the detection limit of the in situ X-ray method owingto experimental error given by the restrictions adopted inthe fast detection mode. The increase in tensile stress withincreasing temperature is thus taken as the main evidencefor the activity of process I (Table 1). Process I includesreverse martensitic transformation of the relict martensitecoupled with changes in intergranular stresses in the micro-structure responsible for the negative thermal expansion ofthe cold-worked wire. This process is reversible if the max-imum achieved temperature remains


includes recovery of lattice defects through migration,redistribution and annihilation and rearrangement of crys-tal defects (polygonization), followed by crystallizationfrom the amorphous phase [21].

At �400 �C, the tensile stress (curve 2 in Fig. 9) passesthrough a maximum and starts to decrease with increasingtemperature. The rate of change in electrical resistivity(curve 3), radial strain (curve 4) and FWHM (curve 5) withincreasing temperature (Fig. 9b) is fastest. This is inter-preted as follows. As process I slows down owing to thedisappearance of martensite and internal stress, the ther-mal expansion and plastic deformation due to the startingdynamic recrystallization (process III) begin to have adominating influence on the tensile stress, which starts todecrease. At this stage of the microstructure recovery, theheat-treated wire starts to show superelastic response dueto the phase transformation. An example is the 3 W treatedwire (Fig. 6b), for which the electrical resistivity (Fig. 7,curve b) as well as diffraction results (Fig. 8) are clear evi-dence that pronounced irreversible microstructure changesdue to process II have already taken place. The tensilestress decrease is fastest, and the {1 1 0} austenite peakprofile changes most dramatically at �500 �C. This is inter-preted as being due to pronounced activity of process II(particularly the crystallization from amorphous) and pro-cess III dominated by dynamic recrystallization. The alloy

is extremely susceptible to plastic deformation in this stage.Since the plastic deformation is due to dynamic recrystalli-zation, the strain rate can be very high, and no additionallattice defects are introduced into the microstructure. TheP = 4 W treatment (Fig. 7, curve c) represents an exampleof a treatment in which the dynamic recrystallization haspartially taken place. The superelastic functional propertiesof the 4 W treated wire are not, however, fully developedyet. Such treated wires typically show R-phase transforma-tion during tensile cycling, as evidenced by the hystereticelectrical resistivity–strain dependence (Fig. 6c).

At �600 �C, the rate of decrease in electrical resistivity(Fig. 9, curve 3) with increasing temperature suddenlyslows down, giving rise to the characteristic knee pointon the electrical resistivity–time (temperature) records(Figs. 7 and 9). This is considered to be an importantthreshold point in the microstructure evolution upon heat-ing [21], which is roughly associated with the terminationof process II. Processes III, however, continues furtheron heating beyond 600 �C. The tensile stress (Fig. 9, curve2) decreases to zero, FWHM (Fig. 9, curve 5) approachesits minimum saturated level, and the radial strain (Fig. 9,curve 4) increase slows down. Although the plastic defor-mation stops as the tensile stress reaches zero, the wire con-tinues to elongate with increasing temperature, owing tothermal expansion. This elongation is, however, reversible

Fig. 9. Combined in situ tensile stress (2), electrical resistivity (3) and X-ray diffraction results (radial strain (4), FWHM (5), peak intensity (6)) recordedduring P = 0 W/0.18 s treatment plotted together in dependence on (a) time and (b) temperature.



on cooling. Since the wire is not stretched by the tensileforce for a short time, it can temporarily move out of thebeam. This unfortunately happened in the case of the6 W treatment. This is the reason for the gap in the 6 W dif-fraction data (Fig. 8).

The shape setting of the NiTi wire is completed [32]when the tensile stress reaches zero, which roughly corre-sponds to the 5 W treated wire (Fig. 7, curve d). The func-tional superelastic properties are fully developed (Fig. 6d).However, the peak height of the austenite peak (Fig. 9,curve 6) further increases with increasing temperature. Thisis attributed to the dramatic change in texture evolutionduring heating, related to the completion of the recrystalli-zation and grain growth now at zero external stress (pro-cess IV). Direct evidence of the grain growth and changein texture evolution during the electropulse heat treatmentwas given in a related paper [18]. The growth of recrystal-lized grains beyond 600 �C in the case of the 6 W treatedwire (Fig. 7, curve e) causes a significant deterioration inthe functional properties of the wire: the yield stress forplasticity as well as the strength decrease [19] and the super-elastic response of the wire during cyclic tensile loading(Fig. 6e) [22] become unstable. If the tensile stress persistsup to temperatures >600 �C, plastic deformation due toslip in the newly formed austenite grains takes place [32].This, however, has never been the case in the present exper-iments, so it is not discussed here.

One may ask what kind of shape setting is referred to ifthe wire has the same shape and length after treatment as ithad before treatment. Since tensile stress levels recorded atroom temperature before and after treatment are similar(Fig. 6) and, since the elastic modulus of the wire almostdid not change after treatment, this suggests that the lengthof the wire did not change during treatment in the firstapproximation. Considering the high temperature andstress involved, this is apparently surprising. However, thisis exactly what the shape setting is for. Recall that anunconstrained wire, if treated by P = 10 W/0.18 s, becomes�3% shorter, while the internal stress and retained mar-tensite disappear from the microstructure. The treated wiredoes not show the negative thermal expansion characteris-tic of untreated cold-worked wire. This can be interpretedas if a sort of plastic deformation process occurred duringthe 10 W treatment, even if the wire length did not change(Fig. 9). The plastic elongation due to process III compen-sates for the shortening driven by process I. The P = 10 W/0.18 s treatment is thus sufficient for shape setting of thewire, and at the same time yields optimal superelastic prop-erties. Since the shape (length) of the wire is set at the hightemperature when the stress approaches zero and the wireshortens upon cooling, tensile stresses recorded at roomtemperature after all treatments are similar (�400 MPa).The fact that the wire was preloaded to 400 MPa is onlya coincidence. The shape setting is a remarkable phenome-non and a technologically important process. It deservesspecial attention and is treated separately elsewhere [32].An important conclusion from the present experiments is

that the shape setting process allowing for simultaneouslysetting the shape and functional properties of the cold-worked filament is directly linked to dynamic recrystalliza-tion process III.

4.2. Application and peculiarities of the FTMT-EC method

4.2.1. Control of microstructures and functional properties

Taking into account the results of earlier TEM [21,22]and X-ray [18] studies, it can be concluded that irreversiblemicrostructure changes due to processes I–III (Table 1 andabove discussion of Fig. 9) bringing about the superelasticfunctional properties of the cold-worked NiTi wire takeplace gradually in the present heat treatments as the tem-perature is raised from 200 �C to �700 �C. Hence, in orderto set the functional properties as desired, it is necessary tocontrol the progress of the microstructure evolution duringheating. In conventional furnace treatment, these processesoccur in an uncontrolled manner as soon as the wire entersthe pre-heated furnace. Obviously, precise control of theprogress of the multiple processes with different fast kinet-ics (timescale 10 min). Diffu-sion processes such as precipitation or dissolution ofprecipitates taking place during the longer exposure ofthe wire to elevated temperatures also affect the functionalproperties of medical thin NiTi wires [17]. Therefore, itbecomes very difficult to distinguish whether the functionalproperty change is due to the microstructure recovery or tothe diffusion processes. The situation is different with theFTMT-EC method, in which the rise in the wire tempera-ture can be controlled in the millisecond range. Since theactivity of the precipitation processes is thought to be lar-gely suppressed in the short-time FTMT-EC treatments,it is possible to establish a true heat treatment parame-ter–microstructure–functional property relationship forNiTi wires (Fig. 6 and microstructure details in Refs. [18]and [21]).

It turns out that, among the performed treatments, thebest superelastic properties were achieved using treatmentsP = 5 W/0.9 s or P = 10 W/0.18 s. The P = 3 W/0.9 streatment is not sufficient to allow for the full developmentof the stress-induced transformation and plateau-typesuperelasticity during tensile testing, while the P = 6 W/0.9 s treatment results in a loss of stability of the superelas-tic behaviour owing to plastic deformation accompanyingthe stress-induced phase transformation during tensilecycling. It is known from the earlier ex situ studies per-formed using P = 125 W and pulse time 1–18 ms [21] thatthe microstructure corresponding to the optimum stablesuperelastic response of the wire is a partially polygon-ized/recrystallized microstructure with grain size 20–



50 nm. The P = 10 W/0.18 s treated wire is expected topossess a very similar microstructure to that found in the125 W/12 ms-treated wire (Fig. 1c), although no directTEM evidence is available. As reported in a related paper[22], the superelastic response of such treated wire is stable,since the stress-induced martensitic transformation canproceed in the nanosized microstructure while the disloca-tion generation at the propagating habit plane interfacesand/or slip in the martensite phase are suppressed by thesmall grain size.

4.2.2. Unique features of the FTMT-EC treatment

Based on the results of the in situ X-ray studies reportedin this paper, it became possible to interpret the complexelectrical resistivity and tensile stress responses in termsof the activity of several distinct processes (Table 1) trig-gered at different temperatures during heating. Althoughthe same processes take place also during conventional fur-nace treatments, there are significant differences which needto be discussed.

First is the very fast rate of temperature increase typi-cally employed in the FTMT-EC treatment. It leads toextreme overheating which differently affects individualprocesses with different intrinsic kinetics. Although super-elastic functional properties of the wire can be obtainedin treatments using pulse duration t_2 = 12 ms as well as1 h, there are significant differences in microstructures.

Second is the possible direct action of passing elec-trons [33] on the individual processes, such as pointdefect mobility [34], recrystallization [33,35] as well asplastic deformation [33,36,37]. The peak electrical currentdensities used in the present experiments are in the range104–105 A cm�2. According to the literature, it seems thatthe major effect of the electric current (continuous orpulsed current with densities >103 A cm�2) on the recrys-tallization is an enhancement of the recrystallization ratein the early stages, resulting in a finer recrystallized grainsize [33,35]. Subsequent grain growth is enhanced as well.Electric current applied during plastic deformation doesnot alter the physical deformation mechanism, butmainly reduces the thermal component of the flow stress,causing a significant increase in plastic strain rate atgiven temperatures [33,36,37]. In contrast to the earlierliterature studies, where DC pulsed or AC electric currentwere employed under quasistatic thermomechanical con-ditions, together with external environmental heating, thiswork deals with physical processes occurring duringultrafast heating, performed by a single short DC electri-cal pulse of controlled power P(t). The enhancement ofrecrystallization and plastic strain rates by electric currentis crucial, particularly for the short pulse experimentswhere the current densities used are higher. The threeabove-mentioned FTMT-EC specific features (fast heat-ing rate, dynamic recrystallization under varying stressand direct effects of the electric current) are expected toaffect dramatically the formation of the polygonized/recrystallized nanograin microstructures in the FTMT-

EC-treated NiTi wires. Although further research is nec-essary to recognize their individual effects, there seems tobe a possibility of using the special FTMT-EC treatmentsto produce NiTi wires with special microstructures andspecial functional properties not achievable with furnacetreatments.

Third is the role of wire thickness. The thinner the wire,the less efficient is the Joule heating, since more heat is lostinto the environment owing to the larger surface to volumeratio of the thin wire. This disadvantage can be beneficiallycompensated using a shorter electrical pulse. Note that,although the same energy is supplied to the wire in theP = 2 W/0.9 s and P = 10 W/0.18 s treatments, only thelatter treatment is sufficient to develop superelasticity.The maximum wire temperature reached is about twice ashigh in the latter treatment (Fig. 1). At the same time,the heating rate 3500 �C s�1 used in the 10 W/0.18 s treat-ment is �10 times greater than that used in the 2 W/0.9 streatment. The greater heating rate is also more effectivein suppressing the slow processes and gives rise to thesmaller temperature gradient from the wire core to the sur-face. This possibly affects the homogeneity of the nanosizedmicrostructure formed in the wire. It should be noted thatheating rates approaching 50,000 �C s�1 were used insequential FTMT/EC treatments reported in Refs. [18]and [21]. Finally, the electric current density is greater incase of shorter treatments, which, according to Conrad[33], further increases the kinetics of the processes, withall the consequences discussed above.

There are basically two practical ways of performingFTMT-EC treatment of as-drawn NiTi wire, leading tothe desired microstructure and functional superelasticproperties. Either the parameters of the electropulse treat-ment (P, t_2 and constraint resulting in the desired T(t),q(t) and r(t); see Fig. 6) are determined from a preliminaryseries of experimental tests [19,20]. This is used for shapesetting of wire segments embedded in structures or inte-grated in hybrid textiles. Or as an alternative option, it ispossible to heat treat continuous NiTi filaments duringrespooling while controlling the supplied electric power ina closed loop with respect to the change in the electricalresistance of the filament caused by the heat treatment[20]. This is particularly useful for heat treatment of contin-uous NiTi filaments prior to their textile processing [38].

5. Conclusions

Functional superelastic properties of cold-worked NiTiwires can be precisely set using non-conventional heattreatment by pulsed electric current. In contrast to the con-ventional furnace treatment, this method allows for precisecontrol of the rise in wire temperature T(t) and tensilestress r(t) during fast heating (3500 �C s�1). This enablescontrol of the progress of the dynamic recovery and recrys-tallization processes during heating, which are responsiblefor the fast evolution of the microstructure and settingthe superelastic functional properties of the treated wire.



A unique combined in situ experiment was carried out inwhich evolution of the parameters of the synchrotron X-ray diffraction profile of the {0 1 1} austenite reflection,tensile stress and electrical resistivity of the wire weresimultaneously recorded during heat treatment by thepulsed electric current. The in situ results obtained providedirect experimental evidence on the evolution of the inter-nal stress (radial strain) and density of defects (electricalresistivity, peak width) in the fast-evolving microstructure.

Based on analysis of the in situ results and earlierreported microstructure studies, the evolution of functionalsuperelastic properties of the wire during fast heating from20 �C to 700 �C was attributed to the sequence of dynamicprocesses involving subsequent activation of reverse mar-tensitic transformation and residual stress relaxation,polygonization and crystallization from amorphous,recrystallization and grain growth.

The fact that the austenite shape and optimal superelas-tic functional properties of the cold-worked filament can besimultaneously set by the fast electropulse treatment isattributed to the controlled dynamic recrystallization pro-cess, which provides sufficient plastic deformation at veryhigh strain rates and at the same time yields nanosizedmicrostructure of the treated wire.

The decrease in electrical resistivity of the wire withincreasing temperature was used to follow the evolutionof superelastic functional properties during treatment andstop it when the desired properties were reached. Thiscan be beneficially used to heat treat continuous NiTi fila-ments during respooling while controlling the suppliedelectrical power in a closed loop with respect to the changein the electrical resistance of the filament caused by heattreatment.

Acknowledgments

The authors acknowledge the support of this work byEC within the integrated AVALON (NMP2-CT-2005-515813) Project, the support of ESRF for performing thein situ synchrotron experiment (MA-358) and supportfrom Czech National Research Projects GACR P108/10/1296, AV0Z10100520, IAA200100627 and P107/10/0824.

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