In-situ investigation of the fast lattice recovery during ......diffraction signal varying in response to the progress of the recovery processes are evaluated. Finally, after the treatment

$Page 1: In-situ investigation of the fast lattice recovery during ......diffraction signal varying in response to the progress of the recovery processes are evaluated. Finally, after the treatment$
Glassy, amorphous and nano-crystalline materials: Thermal physics, analysis, structure and

properties, Springer (Germany/USA/UK), editors J. Šesták, J.J. Mareš, P. Hubík (2010), in print.

1

In-situ investigation of the fast lattice recovery during electropulse treatment of

heavily cold drawn Ni-Ti wires

P. Sittner1, J. Pilch1, B. Malard1,*, R. Delville², C. Curfs3

1Institute of Physics, Na Slovance 2, 182 21 Praha, Czech Republic ²EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium 3ESRF, 6 rue Jules Horowitz 38043 Grenoble, France * Now at SIMaP, Domaine Universitaire BP 75 38402 Saint Martin d'Hères, France

Abstract

Lattice recovery processes responsible for evolution of microstructures in 0.1 mm thin cold-drawn Ni-Ti shape memory alloy wire heat treated by DC electric pulse were investigated by combination of in-situ tensile force, electrical resistance and X-ray diffraction measurements. The X-ray data were used to obtain direct experimental information on the evolution of the phase fractions, internal strain and defects in the microstructure evolving through activation of a sequence of lattice recovery processes during the short time electropulse treatment. It is shown that superelastic functional properties of the treated Ni-Ti wire can be precisely set by controlling the progress of the lattice recovery processes by prescribing the temperature T(t) and tensile stress σ(t) in the treated wire. 1. Introduction

Shape memory alloys (SMA) such as the near equiatomic Ni-Ti alloy [1] have attracted considerable attention for their unique functional thermomechanical properties as superelasticity or shape memory effect deriving from the martensitic transformation. Ni-Ti wires are being produced from extruded bars by multiple hot working passes finished by a final cold drawing. In this so called “cold worked” (as-drawn, hard, etc.) state, the alloy possesses a heavily deformed microstructure resulting from severe plastic deformation [2] consisting of mixture of austenite, martensite, and amorphous phases with defects and internal strain [3]. In this state, the wires do not show any functional property.

As-drawn Ni-Ti wires need to be heat treated so their cold worked microstructure changes into an annealed austenite microstructure for which the wire shows the desired functional properties. At the same time, if the shape of a SMA element is constrained during this final heat treatment, it acquires new “parent shape”. The final thermomechanical treatment thus has two purposes – setting the functional properties and setting the new shape of the wire.

Conventionally, this final heat treatment is performed in an environmental furnace [4-6]. Depending on the alloy (chemical composition, cold work) and desired application, heat treatment conditions in the range of temperatures 400o C-500o C and time 10-60 minutes [4] are applied. If the shape of the wire is constrained during the heat treatment [5], the stress generated in the wire affects the functional properties. This is the reason why the tensile force applied during the commercial straight annealing treatment of superelastic Ni-Ti wires is a very important technological parameter of the Ni-Ti wire production process.

Recently, the relatively long time of the conventional heat treatment of Ni-Ti wires became an obstacle for emerging technology of textile fabrics knitted or woven using continuous Ni-Ti filaments. Maximum respooling speed of ~1 m / min achievable with straight annealing treatment in conventional ~6 m long tubular electrical furnaces is still painfully slow for this purpose. To solve this problem the possibility to treat continuously long thin Ni-Ti filaments by passing electric current through it during respooling has been investigated as an alternative method to the conventional treatment. The method has been developed based on a series of dedicated studies [7-10] and called “Final Thermo Mechanical Treatment by Electric Current” /FTMT-EC/. Further experimental details can be found in Refs. [8, 9].



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Figure 1: Microstructures, electron diffraction patterns and tensile stress-strain response of thin Ni-Ti wires just after cold drawing (left) and after electropulse heat treatment 125 W / 12 ms [10] As reported in related works [7-10], thin superelastic Ni-Ti filaments heat treated by the FTMT-EC method display excellent functional properties due to the specific nanosized microstructures found by TEM in the treated wires. This suggests that the lattice recovery processes responsible for the functional property setting (polygonisation, crystallization form amorphous, recrystallization, grain growth, plastic deformation, etc.) are capable to change the cold worked microstructure (Fig. 1) into the annealed one very fast (Fig. 2).

Major difference with respect to conventional furnace heat treatment is that the heat comes from inside the thin wire and the temperature can rise and fall very quickly. The time of heat treatment is in the order of milliseconds instead of minutes. The opportunity to control the fast evolving sample temperature and tensile force by the FTMT-EC method (Fig. 2b) has opened the way for investigations of the progress of the lattice recovery processes closer to the rate at which they actually proceed during the heat treatment. This motivated us to perform dedicated experimental studies during the electropulse treatments by combination of in-situ tensile force measurement, electrical resistance measurement and high speed synchrotron X-ray diffraction with the aim to obtain direct experimental information on the progress of the lattice recovery processes forming nanocrystalline microstructures at heating rates of the order of ~5000o C / s. 2. Experiment and method

2.1 Heat treatment

All experiments were performed on Fort Wayne Metals #1 superelastic as-drawn Ni-Ti wires (56.0 wt.% Ni) having a diameter d = 0.1 mm. The wire is first mounted on the miniature deformation rig especially designed and built for efficient testing of functional thermomechanical properties of thin Ni-Ti filaments in tension. The rig consists of a stepping motor, a 100 N load cell, electrically isolated grips, a Peltier furnace, a laser micrometer for strain measurement and special electronics allowing to send controlled electric power pulse to heat the Ni-Ti wire up to the melting point and perform simultaneously electric resistance measurement.

The initial length l0 (~50 mm) and electrical resistance ρ0 of the wire at room temperature were first evaluated. The wire was then preloaded to a reach starting values of the mechanical constraint σ0 and ε0. In the present experiments, the wire was preloaded to 400 MPa and its length was fixed. The wire was than heated by a controlled DC power pulse P (Fig. 2a). It is essential that the electronics system is capable of controlling the desired power P (t), even if the electric resistance of the wire drastically changes during the Joule heating event. Evolution of the wire temperature, T(t), as a function of time (Fig. 2b) is calculated according to Eq. 1, taking into account the Joule heat supply P and the ambient temperature losses.



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( )( ) ( )( ) 4TATtThPCtTdt

dext ⋅⋅⋅−−⋅−=⋅ σε

(1)

The heat capacity, C, is assumed to be temperature independent, σ is a Stefan-Boltzman constant and A is the surface area of the wire. The specific heat transfer coefficient, h, describing the heat dissipation into air per unit time and the NiTi wire emissivity, ε, needed for the calculation of the radiation heat loss per unit time were indentified from series of calibration experiments. Other effects influencing the wire temperature such as heat conduction losses into grips and latent heats were neglected. In some experiments described below, two parameters of the power pulse (maximum power P and pulse duration t_2) are used to represent the temperature history T = T(t) the wire is exposed to. In other experiments, the temperature profile T(t) is directly set by controlling the P(t) and in accordance with Equation (1).

Figure 2: a) Parameters describing the DC power pulse applied in the FTMT-EC treatment, b) temperature profiles T = T(t) achieved in pulses with parameters P = 2, 3, 4, 5, 6 W; t2 = 0.9 s and P = 10 W, t2 = 0.18 s As the temperature of the wire rises and falls during the treatment (Fig. 2b), thermally activated recovery processes are triggered and proceed with their characteristic intrinsic kinetics which changes with temperature. The electric resistance of the wire, macroscopic tensile force and X-ray diffraction signal varying in response to the progress of the recovery processes are evaluated. Finally, after the treatment is finished, superelastic response of the treated wire in 10 tensile cycles at room temperature is evaluated using the same stress rig. The parameters determined from the superelastic stress-strain curves and their stability during cyclic tensile loading serve as measures of the achieved functional properties of the treated wires. 2.2 Synchrotron X-ray diffraction

Synchrotron radiation enables to extend the X-ray measurements from the static to the dynamic regime, thanks to its unique time structure and very high flux even at high energy, which reduces considerably the data collection time. Very fast time resolution down to the picosecond regime is achievable [11] with stroboscopic studies of reversible phenomena. In case of the irreversible microstructure evolution, the X-ray data have to be acquired on a single shot basis [12] requiring high flux source, very fast and sensitive data acquisition system and a high speed data transfer.

The experiment was performed on ID11 diffractometer at the European Synchrotron Radiation Facility (ESRF) at Grenoble. A monochromatic X-ray beam with energy of 45 keV (λ ~ 0.27552 Å) is obtained with a Laue monochromator. The beam is focussed down to 1 mm in vertical direction by bending the Laue crystal and down to 100 µm in horizontal direction. Frelon2K 2D camera with taper has been chosen as a fast detection system [12]. In order to speed up the data acquisition for the targeted 10 ms time resolution only a radial slice of the 2D detector (Fig. 3a) is used.



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a) b)

Figure 3: a) Setup of the in-situ high speed synchrotron X-ray experiment during the heat treatment by

electric current. The FReLON camera reading was optimized for high speed data acquisition (only part of

the detector – strip of 2048 × 64 pixels – was used to increase the time resolution to ~10 ms); b) The

evolution of the 110 austenitic peak measured during heat treatment experiments P = 2, 3, 6 W / 0.9 s and

P = 10 W / 0.18 s.

The thin Ni-Ti wire mounted between the grips of the stress rig is positioned vertically in the beam (Fig. 3a). The heat treatment experiment is performed with synchronized pulse, force, electrical resistivity and diffraction measurements (Fig. 4). A strip of 2048 x 64 pixels is acquired and 64-binned in order to obtain a line of 2048 x 1 pixels. Then, this line is read out of the CCD and stored before next image is acquired. Afterwards, 300 lines are put together into a single 2D binary image which represents the evolution of the diffraction patterns versus time, as illustrated in figure 3b. To minimise the time needed to read out the line, the 64 active pixels have to be chosen at the bottom edge of the camera. The part of the detection area, which is not active, is masked by a lead mask. By using such kinetic mode of the data acquisition by the FReLON camera, it has been possible to acquire 100 diffraction patterns per second with a readout time of less than 1 ms between 2 images.

Figure 4: Synchronisation between the electric power pulse and X-ray measurements

In order to obtain a diffraction pattern of the intensity versus the diffraction angle for the kinetic mode, a calibration was performed. A very fine capillary of 300 µm with LaB6 powder was placed at the Ni-Ti wire position and 2D diffraction pattern was acquired. The position of each LaB6 peak in pixels was measured using DASH [13] and the corresponding diffraction angle was calculated. The calculated calibration function relating the diffraction angle to the pixel number was used to transform each line of the experimental pattern (intensity versus pixel) into a diffraction pattern (intensity versus diffraction angle). Since the beam centre is in the middle of the detector, the right side of the diffraction pattern is added to its left side to increase the statistics.



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In order to obtain an etalon for stress-free Ni-Ti diffraction pattern, Ni-Ti powder of similar composition as that of the wire was placed in the LaB6 container and diffraction pattern was acquired. Since the Ni-Ti powder existed in a mixture of austenite and martensite phases at room temperature, the diffraction pattern contained both austenite and martensite diffraction lines. The Ni-Ti powder gives diffraction pattern with much better resolution (Fig. 5a) than the thin as-drawn Ni-Ti wires (Fig. 5b). This is due to the complex microstructure of the thin Ni-Ti wires consisting of a mixture of heavily deformed austenite, R-phase, martensite and amorphous phases and frozen in internal stresses. After heat treatment yielding ideal superelastic properties, the microstructure of the wire is essentially austenitic (Fig. 5c). The diffraction lines were analyzed for integrated intensity, peak position and FWHM of the 110 austenite peak using the LAMP software [14] as well as the GSAS software [15] considering the structure of the austenite phase Pm-3m, R phase P-3 and martensitic phase P21/m [1].

Figure 5: Quantitative phase analysis of the diffraction patterns to determine the evolution of the

integrated intensity, position and FWHM of the 110 austenitic peak a) Ni-Ti powder; b) Ni-Ti wire before

heat treatment; c) Ni-Ti wire after heat treatment.

3. Experimental Results

3.1 Heat treatment experiments – macroscopic results

The heat treatment experiments were performed using different power P = 2, 3, 4, 5, 6 W (t2 = 0.9 s) and 10 W (t2 = 0.18 s). Same experiments were performed twice with identical results - once with simultaneous in-situ X-ray diffraction once without. The results of 3 selected experiments are presented in figure 6 showing the variations of the supplied electric power, temperature, tensile stress and electric resistivity during the FTMT-EC pulse on the left side and the obtained superelastic stress-strain response with superimposed variation of the electric resistivity of the wire on the right side. The observed evolution of tensile stress and electric resistivity in response to the prescribed temperature evolution during the treatment (Fig. 2b) is analyzed below..

The tensile stress evolves with time (temperature) due to the several mutually competing processes: i) relaxation of residual stress, ii) thermal expansion, iii) plastic deformation processes, iv) reversible phase transformations. The tensile stress starts to increase with increasing temperature right after the onset of heating (Fig. 6). This is due to the unlocking of elastic deformation held by internal stresses frozen in the heavily cold worked microstructure of the as-drawn Ni-Ti wire coupled with reverse transformation of small fraction of the martensite phase which retransforms back to the austenite phase upon heating. The stress increases with the increasing temperature up to 400o C-500o C, where it reaches a maximum of ~690 MPa (Fig. 6 and Fig. 7b) and falls ultimately down to zero stress depending on the maximum temperature reached. This is due the plastic deformation and thermal expansion processes prevailing at high stress and temperature. Upon cooling, the stress varies basically due to thermal expansion.



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Figure 6: Evolution of power, temperature, tensile stress and electric resistance (left) of the Ni-Ti wire during FTMT-EC treatments with maximum power P = 3, 6, 10 W / 100 mm and resulting superelastic

response and electric resistivity records of the treated wires (right).

Figure 7: Comparison in experiments with maximum power P = 2, 3, 4, 5, 6 W / 0.9 s and 10 W / 0.18 s of

a) the evolution of the tensile stress (left) and the electrical resistivity (right) with time b) the evolution of

tensile stress (left) and electrical resistivity (right) with temperature



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The electric resistivity starts to decrease when the temperature reaches ~200o C (Figs. 6 and 7) and decreases monotonically during heating until it reaches a plateau at ~75% of its starting value. The electric resistivity changes with the temperature due to: i) intrinsic thermal dependence of the resistivity of a metallic wire, ii) reversible phase transformations and iii) progress of the lattice recovery processes causing irreversible microstructure changes. If an already fully heat treated Ni-Ti wire is subjected to the same FTMT-EC pulse [7], the electric resistivity monotonically increases with increasing temperature and returns to its original value after cooling back to room temperature. This is a simple proof that the electric resistance decrease is caused by the lattice recovery processes (iii). Upon cooling, the electrical resistivity again decreases due to intrinsic thermal dependence of the electric resistance of a metal (Fig. 7). Overall, the total decrease of the wire electric resistance due to microstructure recovery reaches ~40% - from 1.45 Ωµm to 0.88 Ωµm for the 6 W / 0.9 s treatment.

3.2 Heat treatment experiments - In-situ X-ray diffraction results

Figure 3b shows the evolution of the diffraction pattern near the 110 austenitic peak as measured by the 2D detector during the experiments using pulses with power P = 2, 3, 6 and 10 W (see also the synchronization in figure 4). It evolves from the broad diffraction pattern corresponding to the complex microstructure of the as-drawn wire to the well-defined narrow peak of the heat treated wire. The intensity of the peak increases, the peak shifts and its FWHM decreases during the electric pulse treatment. Figure 8a shows the evolution of the integrated intensity of the 110 austenitic peak during the experiments. Before the treatment, the austenitic volume fraction is estimated to ~15%, the rest being amorphous with possibly a very small amount of retained martensite. The higher the temperature reached in the treatment, the larger the maximum integrated intensity. The shift in the onset of the increase of the intensity to shorter time with increasing P is due to the fact that the temperature rises faster for more energetic pulses (Fig. 2b).

Figure 8b shows the results of the analysis of the evolution of 110 austenite peak position recalculated into radial strain (Equations 2-4) of the 110 austenitic peak for treatment with P = 2, 3, 6 and 10 W pulses. Radial strain means elastic strain in transverse direction. The axial strain would also be of interest but it could not be evaluated due to the geometry of the diffraction experiment (Fig. 3a). Using the Bragg law (Equation 2) and the definition of strain (Equation 3) an expression for the radial strain of the 110 austenitic peak ε110 is given by equation (4).

λλλλθθθθ .sin.2 nd hklhkl = (2)

hkl

hkl

hkl

hklhkl

hkld

d

d

dd

,0,0

,0 ∆=

−=εεεε (3)

110,0

110110 sin

sin1

θθθθ

θθθθεεεε −= (4)

θ110 is the value of the Bragg angle of the 110 austenitic peak in the wire during the heat treatment and θ0; 110 is the value of the Bragg angle 110 austenitic peak measured on the powder. Because the radial strain is strongly negative before the heat treatment experiment (Table 1), it means that the 110 austenite lattice of the as-drawn wires experienced a tensile stress along its axis. Assuming a Young modulus E = 60 GPa and a Poison ratio 0.33, the stress is evaluated to 2290 MPa. Subtracting the 400 MPa external tensile stress, the tensile internal stress (corresponds to 3.15% tensile elastic strain) existing in the as-drawn wire prior the heat treatment comes down to 1890 MPa. If the as-drawn wire is heated in stress free conditions, it becomes ~3% shorter and the radial strain equals to zero after the treatment. As the temperature increases during the heat treatment, the austenite lattice expands in radial direction due to two processes i) isotropic thermal expansion and ii) directional decrease of the internal tensile stress. Recall that the wire is exposed



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to oscillating external tensile stress (Fig. 7a) which affects the radial strain as well. The radial strain increases and decreases but does not return to its original starting value -12600.10-6 µstrain in any of the treatments (Table 1). Table 1: Key features of the evolution of the peak position in heat treatment experiments

Power 2 W 3 W 6 W 10 W

Onset of the peak position change observed at (ms) 1260 1224 1116 1020 Maximum radial strain observed (ms) 1920 1716 1188 Radial strain value before treatment (µstrain) -12600 -12600 -12600 -12600 Maximum radial strain value (µstrain) 1140 1055 7290 Radial strain after treatment (µstrain) -6225 -5175 -2755 -2145

Figure 8b evidences how internal tensile stress frozen in the wire after the final cold drawing reduction becomes relaxed during the FTMT-EC treatment. The radial strain -2145.10-6 µstrain measured after the 10 W / 0.18 s treatment (Table 1) gives a tensile stress of 390 MPa which nearly equals the external tensile stress. This suggests that the internal tensile stress does not exist in the 10 W / 0.18 s and 6 W / 0.9 s treated wires but it still partially remains in the wires treated with 2 W / 0.9 s and 3 W / 0.9 s pulses.

Figure 8: Evolution of a) the integrated intensity, b) the position (radial strain), c) the FWHM of the

austenite given by the 110 austenitic peak during experiments at P = 2, 3, 6 W / 0.9 s and 10 W / 0.18 s.



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Figure 8c shows the results of the analysis of the evolution of the width of the 110 austenitic peak (FWHM) for experiments with P = 2, 3, 6 and 10 W pulses. FWHM is equal to 0.35° before the heat treatment. It decreases with increasing temperature mainly due to the recovery of the crystal defects and related strain fields existing in the cold worked microstructure, crystallization of amorphous and/or due to the increase of the subgrain size. While in case of the P = 10 W and 6 W heat treatments, the decrease of the FWHM from 0.35° to ~0.1° corresponds to well annealed microstructure, the decrease to only 0.2° in case of the 3 W treatment, suggests that this microstructure probably still contains many defects and related strain fields. In case of the P = 2 W / 0.9 s treatment, the FWHM almost does not change suggesting that the defects and strain fields essentially remained in the microstructure.

4. Discussion The lattice recovery of the cold worked microstructure of the Ni-Ti wire takes place as the temperature increases from 200o C to 700o C during the heat treatment. The question is whether it is possible to control the progress of the lattice recovery processes with a sufficient precision. In conventional furnace treatment, the lattice recovery processes occur in uncontrolled manner as soon as the wire enters the pre-heated furnace. Obviously, precise control of the progress of the lattice recovery processes having fast kinetics (~ms) is not possible by environmental furnace treatment. In the FTMT-EC treatment, the chance is better since the temperature profile T(t) is prescribed.

In addition to the lattice recovery processes, there are diffusion processes characterised by much slower kinetics (min.) such as precipitation or dissolution of precipitates which may proceed during the exposure of the sample to elevated temperatures and affect the functional properties of the wire as well [6]. Hence, in conventional furnace treatments, it is not possible to find out whether the change of functional properties is due to lattice recovery processes or/and to diffusion processes. Since the activity of the diffusional processes is largely suppressed by the short time FTMT-EC treatments, it is possible to establish the heat treatment parameter – microstructure - functional property relationship for heat treated Ni-Ti wires [7, 10].

Obviously, the principal technological goal was to find out the parameters of the FTMT-EC treatment which leads to optimum superelastic functional properties of the treated Ni-Ti wires. It comes out that, among the performed treatments, the optimal ones would be the 5 W / 0.9 s or 10 W / 0.18 s. The 3 W / 0.9 s treatment is not sufficient and the 6 W / 0.9 s treatment results in a loss of stability of the superelastic behaviour due to plastic deformation accompanying the stress induced phase transformation during tensile cycling. Much more detailed results concerning the functional property setting of Ni-Ti wires were reported in [8]. Detailed TEM investigation of microstructures in FTMT-EC treated wires [10] confirmed that the functional properties of Ni-Ti wires can be very well correlated with the microstructures. The microstructure corresponding to optimum superelastic response of the wire is a partially polygonized/recrystallized microstructure with grain size 20-50 nm. Such microstructure was found in a wire treated by 125 W / 0.012 s pulse [10]. The 10 W / 0.18 s treated wire is assumed to contain a very similar microstructure although there was no direct observation. Five lattice recovery processes (Table 2) are expected to become subsequently active during the heat treatment of as-drawn Ni-Ti wires. Though the same processes take place also during conventional heat treatments in a furnace, there are differences that need to be emphasized. First is the fast rate of the FTMT-EC treatment possibly leading to overheating which differently affects the recovery processes with different kinetics, second is the mechanical stress action on the recovery processes and third is the possible direct action of passing electrons on the recovery



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processes. It is assumed that all these three FTMT-EC specific features assist the formation of the desired polygonized/recrystallized nanograin microstructure [7, 10] in the FTMT-EC treated Ni-Ti wires.

The goal of this work was to obtain more detailed information on the lattice recovery processes with the help of the high speed synchrotron X-ray diffraction experiments. Particularly, the goal was to interpret the apparently curious in-situ electric resistance-time and stress-time responses (Figs. 6 and 7) in terms of the activity of various lattice recovery processes subsequently taking place during the heat treatment. To synthesize the obtained results, the macroscopic and diffraction results from 10 W / 0.18 s treatment were plotted together in figure 9. Table 2: Recovery processes expected to occur during electropulse treatment of Ni-Ti wires. Processes take

place in sequence I-V as the temperature increases and those dominating in given temperature range are in

bold. Recovery process Temperature range

I Residual stress relaxation & reverse martensitic transformation 20o C - 500o C II Polygonization & crystallization from amorphous phase & subgrain growth 200o C - 700o C III Recrystallization & grain growth & plastic deformation 500o C - 900o C IV Plastic deformation & grain growth > 800o C V Grain growth > 1000o C

a)

b)

Figure 9: In-situ macroscopic and diffraction results of the P =10 W / 0.18s experiment plot together in

dependence on a) time, b) temperature



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The tensile stress (2) increases from the onset of heating and electric resistance (3) starts to fall at 200° C before any change of the diffraction signal can be detected. However, one has to admit that the experimental error related to 110 peak analysis is too large at temperatures up to 300° C. The increase of the tensile stress is thus taken as a main evidence for the recovery process I. This process is largely reversible if the maximum achieved temperature remains relatively low (see the stress-temperature response during the 2 W / 0.9 s treatment - curve a in figure 7). It can be deduced that the process I dominates over the linear thermal expansion up to ~450o C, otherwise the stress would have to decrease with increasing temperature due to thermal expansion. Upon further heating, the electrical resistivity (3) falls, integrated intensity (6) and radial strain (4) starts to increase at 360o C. It is assumed that the decrease of electrical resistivity starting at ~200o C evidences the beginning of the irreversible microstructure change due to recovery of lattice defects through migration, redistribution and annihilation of point defects and migration of dislocation dipoles into cell walls called polygonization, crystallization from amorphous phase and growth of polygonized domains (recovery process II).

At ~450o C, the thermal expansion and/or plastic deformation accompanying the recovery process II takes over the recovery process I and tensile stress (2) reaches a maximum and starts to decrease. Right after that, at about 600o C, together with the falling tensile stress (2), the FWHM (5) starts to decrease suggesting the accelerating defect annihilation probably due to start of the recovery process III dominated by dynamic recrystallisation. At the same temperature, the rate of change of the integrated intensity also increases. The alloy is extremely susceptible to plastic deformation at this stage. At ~700o C, while the temperature still increases, the rate of the decrease of electric resistivity suddenly changes. This is considered to be an important threshold point in the microstructure evolution which can be roughly associated with the termination of process II. It can be easily recognized as the knee point on the electrical resistivity responses. Upon further heating, the recovery processes III and IV (FWHM decreases, austenite volume fraction increases) continue. The tensile stress had already reached nearly zero terminating the wire plastic deformation. It shall be noticed that, after cooling down to room temperature, the wire has the same length as it had before the test (the tensile stress levels before and after the treatment are similar (Fig. 6 and 7)). On the other hand, recalling that the free as-drawn wire becomes 3% shorter after the 10 W / 0.18 s heat treatment, it is assumed that the plastic deformation did occur even if the wire length did not change. It compensated the relaxed internal strains due to internal stresses and/or martensite phase strains present in the as-drawn microstructure which disappeared during the treatment. The responses upon cooling are much simpler – tensile stress (2) increases and radial strain (4) decreases due to linear thermal contraction taking place under the condition of the constant sample length, electric resistivity (3) decreases linearly due to conventional linear dependence on temperature. As the temperature decreases below ~70o C, the stress (electrical resistivity) starts to decrease (increase) again resulting in local maxima (minima) recorded at temperature 50o C on the respective responses in figures 7 and 9. This is due to the B2-R transformation taking place upon cooling in the already heat treated wires [8]. The situation upon cooling is somehow more complicated in case of lowest power treatments which did not result in complete relaxation of internal stresses (Fig. 7).

Because of heating rates reaching up to 5000 o C / s used and although it was possible to follow the progress of the sequential lattice recovery processes, it is difficult to discuss the kinetics of individual lattice recovery processes, since these are still fast enough to follow the prescribed T(t) profile. Nonetheless, the fact that individual responses in figure 7 do not follow same path upon heating is a clear experimental evidence that the intrinsic kinetics of individual processes matters. In case of the 125 W / 0.012 s treatment [10], in which the heating rate is ~80000o C / s, the intrinsic kinetics of the lattice recovery processes plays much more important role.



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5. Conclusion Functional superelastic properties of the Ni-Ti wires can be precisely set by electropulse heat treatment. Variations of tensile force, electrical resistance and synchrotron X-ray diffraction signal from the 0.1 mm thin as-drawn Ni-Ti wire prestrained in tension were simultaneously recorded during the heat treatment. These data were used to obtain direct experimental information on the evolution of the phase fractions, internal stresses and defects in the microstructure evolving through activation of a sequence of dynamic lattice recovery processes during the short time electric pulse treatment. Main advantage of this method over the conventional furnace heating is the possibility to control the progress of these lattice recovery processes by prescribing the temperature T(t) and tensile stress σ(t) in the wire during the electropulse treatment. Acknowledgments

The authors acknowledge the support of ESRF for performing the in-situ synchrotron experiment (MA-358) and support from projects AV0Z10100520, IAA200100627. References [1] Otsuka K., Ren X.: Physical metallurgy of Ti-Ni based shape memory alloys. Progress in Mat. Sci.. 50, 511-678 (2005) [2] Inaekyan K., Brailovski V., Prokoshkin S., Korotitskiy A., Glezer A.: Characterization of amorphous and nanocrystalline Ti-Ni-based shape memory alloys. J. Alloys and Compounds 473, 71-78 (2009) [3] Schaffer, J.E.: Structure-Property Relationships in Conventional and Nanocrystalline NiTi Intermetallic Alloy Wire. J. of Materials Engineering and Performance. 18, 582–587 (2009) [4]. Duerig T.W, Melton K.N., Stöckel D., Wayman, C.M.: Engineering Aspects of Shape Memory Alloys . Butterworth-Heinemann Ltd., London (1990). [5] Liu X., Wang Y., Yang D., Qi M.; The effect of ageing treatment on shape-setting and superelasticity of a nitinol stent. Materials Characterization, 59, 402-406 (2008) [6] Undisz A., Fink M., Rettenmayr M.; Response of austenite finish temperature and phase transformation characteristics of thin medical-grade Ni–Ti wire to short-time annealing.Scripta Materialia, 59, 979-982, (2008) [7] Malard B., Pilch J., Šittner P., Gartnerova V., Delville R., Schryvers D., Curfs C., Microstructure And Functional Property Changes In Thin Ni–Ti Wires Heat Treated By Electric Current— High Energy X-Ray and TEM Investigations. Functional Mater. Letters, 2, 45-54 (2009) [8] Pilch J., Heller L. and Šittner P., Final thermomechanical treatment of thin NiTi filaments for textile applications by electric current. Proc. of the ESOMAT09, EDP Sciences, 05024 (2009) [9] Czech patent application PV2009-279 [10] Delville R., Malard B., Pilch J., Sittner P., Schryvers D.: Microstructure Changes During Non Conventional Heat Treatment of Thin Ni-Ti Wires By Pulsed Electric Current Studied By Transmission Electron Microscopy, submitted to Acta Materialia, (2009) [11] Plech A.,Wulff M.,Bratos S.,Mirloup F.,Vuilleumier R.,Schotte F.,Anfinrud P.A., Visualizing chemical reactions in solution by picosecond x-ray diffraction. Phys. Rev. Lett. 92, 125505 (2004) [12] Labiche J., Mathon O., Pascarelli S., Newton M, Ferre G., Curfs C.,Vaughan G., Homs A., Carreiras D.: Invited article: the fast readout low noise camera as a versatile x-ray detector for time resolved dispersive extended x-ray absorption fine structure and diffraction studies of dynamic problems in materials science, chemistry, and catalysis. Review of Sci. Instr. 78, 091301 (2007) [13] David W.I.F., Shankland K., Van de Streek J., Pidcock E., Motherwell W.D.S., Cole J.C.; DASH: a program for Crystal Structure Determination from Powder Diffraction Data. J. Appl. Cryst., 39, 910-915 (2006) [14] http://www.ill.fr/data_treat/lamp/front.html. Accessed January 3, 2010. [15] http://www.ccp14.ac.uk/solution/gsas/. Accessed January 3, 2010.

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In-situ investigation of the fast lattice recovery during ......diffraction signal varying in response to the progress of the recovery processes are evaluated. Finally, after the treatment