Relaxation mechanisms in martensitic NiTi(Cu ...

Relaxation mechanisms in martensitic NiTi(Cu): internal friction measurements correlated to in situ TEM straining

Published by Maney Publishing (c) IOM Communications Ltd

M. Parlinska-Wojtan*1,2, R. Schaüblin3 and R. Gotthardt1 The microstructural origin of the internal friction response in B19’ martensitic NiTiCu polycrystals is investigated by in situ TEM straining experiments. Internal friction measurements showed that the transient contribution Q 21 Tr , determined as a function of temperature in martensitic state, vanishes when the heating or cooling is stopped. In situ TEM experiments evidence defects moving across the martensitic variants along the microtwin boundaries. They are identified as partial dislocations (PDs). They move under applied stress as well as under temperature gradients. In the present 21 study a correlation between their movement and the changes of Q 21 Tr is proposed. The Q Tr is found to be sensitive to the density of PDs depinned from the martensite variant boundaries. Their movement appears to be a reversible process, as when stress is released the PDs return into the variant interface and consequently Q 21 Tr value drops. Keywords: Martensitic transformation, Shape memory alloys, Internal friction, In situ TEM deformation, Dislocation structure

Introduction The martensitic phase transformation is a displacive and diffusionless, first order phase transformation, which can be induced either by variation of temperature or by application of stress.1–5 It occurs without diffusion by a cooperative movement of atoms over distances below the interatomic ones. The martensitic transformation is at the origin of properties such as shape memory effect, superelasticity and high damping capacity (internal friction, IF). The most known shape memory alloys are NiTi binary and ternary alloys with additions of Cu, Fe,6 Hf,7 or Pd,8 ternary Cu based alloys like CuZnAl,9,10 or CuNiAl and less common noble metals shape memory alloys such as AuCd or AgCd.1 NiTi alloys are among the most interesting ones for applications.11 The high temperature parent phase called austenite has a B2 structure and the product of the phase transition upon cooling is martensite which has a B19’ monoclinic structure.12 The substitution of Cu affects the transformation behaviour of the NiTi alloys. The hysteresis width is decreased13 and the transformation temperatures are raised slightly above room temperature.14 At Cu contents up to 5 at.-% the alloy transforms directly from the B2 austenitic phase to the

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IPMC-FSB, Swiss Federal Institute of Technology Lausanne (EPFL), 1015 Lausanne, Switzerland Laboratory for Nanoscale Materials, Empa, Ueberlandstr. 129, 8600 Du¨bendorf, Switzerland 3 Fusion Ecole Polytechnique Fe´de´rale de Lausanne (EPFL), Centre de Recherches en Physique des Plasmas, Association EuratomConfe´de´ration Suisse, 5232 Villigen PSI, Switzerland 2

*Corresponding author, email [email protected]

ß 2008 Institute of Materials, Minerals and Mining Published by Maney on behalf of the Institute Received 29 June 2007; accepted 14 November 2007 DOI 10.1179/174328408X302594

B19’ martensitic phase, omitting the R phase.15,16 At Cu concentrations .5 at.-% a two steps transformation from B2 to B19 orthorhombic and then to B19’ is observed.14 The transformation temperatures depend on alloy composition.14 The physical processes that occur in the material during the phase transformation can be divided into two classes. First, there is the formation of the new phase nuclei and, second, there is the growth of the new phase. The growth process can be studied by IF experiment, whereby the transformation is induced by an applied stress whose amplitude and sign change periodically in time at a selected frequency f. The IF peaks show, among other characteristics, a dependence on the amplitude of the applied oscillatory stress.17–19 It should be noted that results obtained for single crystals differ from those obtained with polycrystalline material.20 Moreover, the thermomechanical history of the material strongly influences the IF spectrum.21 In addition, the IF peaks depend on the heating or cooling rate.21,22 This so called T˙ effect23–25 refers to the difference in damping between a measurement performed under isothermal conditions and one performed at a constant increasing (heating) or decreasing (cooling) temperature rate. Various interpretations exist on the transformation mechanisms in NiTi and NiTiCu alloys in relation to IF measurement.25–30 During cyclic torsional deformation in the transition regime, the IF, or dissipated energy Q21, results from the mobility of microstructural features. In addition, the evolution of the free oscillation frequency f is related to changes in the elastic shear modulus of the specimen. In the martensitic phase, below the temperature Mf at which the transformation is finished upon cooling, it is assumed that the relatively

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Internal friction measurements correlated to in situ TEM straining

high IF signal in the fully transformed material is mainly due to the stress dependent movement of interfaces.31 However, to the authors’ knowledge, no direct evidence or quantification of the dependence of the intrinsic part of the IF on the microstructrure, and in particular on the density of martensite variants, has been provided. The aim of the present study is to identify the microstructural mechanisms underlying the martensitic phase transition by performing IF measurements and transmission electron microscopy (TEM) observations. In a previous study9,10 performed on CuZnAl the relationship was established between the microscopic hysteresis effect observed in the IF behaviour that is involved in the microstructural evolution visualised by TEM and the macroscopic hysteresis effect observed during electrical resistance measurements. It was shown that the transient IF peak is dependent not only on the volume fraction that is transformed during oscillation, but also on the density of mobile martensite/martensite interfaces. In the present study it is attempted to perform the same type of measurement and correlate the microstructure mechanisms observed during the in situ TEM straining to the energy dissipated during the IF measurement in the B19’ martensitic phase of NiTiCu. In particular, these two techniques are used to identify the microstructural mechanisms responsible for the drastic decrease of Q{1 Tr below Mf when the temperature rate is stopped during the IF measurement. The difficulty in the in situ TEM straining test of NiTiCu, relative to CuZnAl, resides in the much lower spatial scale at which these mechanisms occur.

Experimental Ni44?2Ti50?6Cu5?2 polycrystals were heat treated for 1 h at 750uC in order to grow large, internally twinned martensite variants. The Cu content (5?2 at.-%) was determined with the microprobe analyser ARL-SEMQ (15 kV, 25 nA) and with an X-ray energy dispersive spectrometer (12 kV, 100 nA). The start and end transformation temperatures of the samples, determined by differential scanning calorimetry, are 65 and 45uC respectively. To prepare the electron transparent specimens for in situ TEM straining the bulk material was first laminated to strips with a thickness of y260 mm, then heat treated and cut into shape by electroerosion and then mechanically polished down to y100 mm. The final thinning down to electron transparency was carried out by electropolishing in a solution of methanol and 20% sulphuric acid at a temperature of 26uC and a voltage of 20 V. The in situ tensile experiments were performed at room temperature in a JEOL2010 high tilt lens TEM with a LaB6 cathode operated at an acceleration voltage of 200 kV and having a point to ˚ . The GATAN low temperapoint resolution of 2?34 A ture straining sample holder was a single tilt holder, with a maximum tilting angle of ¡30u and a maximum displacement of 1500 mm. The GATAN sample holder allows for a constant deformation rate and is therefore suitable for the observations of relaxation mechanisms. The selected applied straining speed was 100 nm s21. All the observations were performed at an end maximum magnification of one million times in weak beam dark field conditions and were filmed with a light amplified TV camera GATAN 622. The acquisition of the film was made directly on a computer, Macintosh G4 400 MHz,


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1 Bright field TEM image of thermoelastic self-accommodated martensite in heat treated NiTiCu polycrystal

equipped with a digitising acquisition card, Aurora Igniter, with a compression rate of 1?5, through Adobe Premiere for software control. The IF was measured in a forced, inverted torsion pendulum with an oscillation frequency of 1 Hz, under a He pressure of 10 mbar, a strain amplitude of e5261025 and a temperature rate of 0?25 K min21. The measured temperature range was 300 to 380 K. The specimen was a flat bar of 506461 mm approximately, and its effective length was 40 mm. In situ TEM experiments allowed the observations of the movement of defects inside the martensitic microstructure. The resulting applied stresses appeared to be sufficiently low so as to impede martensite reorientation.

Results and discussion Martensite: microstructure and IF Figure 1 shows a typical bright field TEM micrograph of the martensitic microstructure of the NiTiCu polycrystal. The number density of interfaces estimated from the TEM micrographs such as in Fig. 1 is 45 mm22. Only the interfaces between the martensite variants and the microtwins in NiTiCu are counted because, as will be shown later in the paper, which directly influence the IF response. Dislocations in the martensitic NiTiCu polycrystal b~0 rule under were analysed in the TEM using the ~ g:~ two beam conditions. There is to the authors’ knowledge no information on the possible Burgers vector in this particular monoclinic phase. Therefore only the direction of ~ b could be identified. Two types of dislocations were analysed: the ones that are in the boundaries and the ones that are seen in the interior of the twins, not attached to boundaries. In Fig. 2 micrographs taken for ~ gb ~(2¯2¯2) and ~ gc ~(11¯1¯) with the respective ga ~(022¯), ~ diffraction patterns are shown. The analysis shows that the Burgers vector ~ b of the dislocations in the boundaries is parallel to [011¯], indicating that the observed dislocations are edge in character. The few long grown-in dislocations seen in the interior of the twin have a Burger’s vector parallel to ~ g for the visibility condition, indicating that they are screw in character. As




2 Bright field TEM images of dislocations in martensite taken at two beam conditions for a ~ g a5[022¯] b ~ g b5[222¯] and c ~ g c5[111¯] with respective diffraction patterns

a heating; b cooling 3 Internal friction measured as function of temperature: T˙ effect is indicated by dots and intrinsic part of IF is depicted by broken line

the edge dislocation core has generally a more coplanar configuration than the screw one, in particular in noncubic structures,32 edge dislocations have a higher mobility than screw ones. These grown-in dislocations seen in the interior of the twins being screw are thus stopped at the end of the sample production process, while edge dislocations, highly mobile, are moving easily and are stopped at the interfaces. Figure 3a and b shows the IF measured for the NiTiCu polycrystal during heating and cooling respectively. In both graphs the Q{1 Tr part is indicated by the black dots and the broken line corresponds to the situation of constant temperature (dT/dt50) the Q{1 dT=dt~0 values. The start and end transformation temperatures As and Mf are indicated on the curves. As expected, the alloy exhibits a low damping capacity in the B2 austenitic phase, a peak in the transition region and a high level of IF in the B19’ martensitic phase.

In situ TEM straining experiments Movement of partial dislocations (PDs) under stress application

Figure 4 shows the in situ TEM straining experiment performed under constant strain rate. The movement of PDs and the subsequent appearance of a change in the

atomic stacking sequence are clearly observed. The PD gliding plane is almost parallel to the observation direction. Dislocation A, which is in the lowest variant interface indicated by the arrow in Fig. 4a, starts to move under low applied stress and can be seen as a change of contrast in Fig. 4b. A dark contrast, corresponding to a strain field, at the head of the dislocation denoted B is observed in Fig. 4a. The dislocation moves down and creates a diffraction contrast after its passage (Fig. 4b). The same event is observed for dislocation C in Fig. 4b and c. These trailing contrasts are the signature of a change in the atomic stacking sequence generated by the passage of a PD. In Fig. 4d a schematic view of the movement of dislocations A, B and C is presented. It corresponds to the 2D representation in Fig. 4e. The partial twin dislocation inducing a change in the atomic stacking sequence is shown schematically in 2D and in plane view as observed in the TEM. In Fig. 4, the dislocations move along their gliding planes and their movement is fluent through the entire trajectory. Figure 5 shows the in situ TEM straining experiment performed in a region in which the dislocation glide plane is slightly more inclined than the one in Fig. 4. In the thin electron transparent samples, the trace of the


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a t50 s; b t52 s; c t53 s; d initial state; e final state 4 a–c TEM weak beam micrographs of PDs moving along twin boundaries under application of stress extracted from film recorded during in situ TEM tensile experiments (see text for details) and d,e schematic representation of initial and final state respectively

dislocations at the surface remain parallel and the moving dislocation is curved due to pinning points at the surfaces: one pinning point is on the upper surface while the other is in the lower one, as indicated schematically in Fig. 5e. The dislocation loop segments (Fig. 5), become mobile when additional stress is applied and the loop segment grows. Their movement is observed to be step wise between pinning points, as seen in the sequence Fig. 5a–c. When there is enough energy supplied by the applied stress, the dislocation loop segment depins and continues, as shown in Fig. 5d. At stresses corresponding supposedly to the beginning of pseudoplastic region in the stress–strain curve the dislocations start to depin collectively, and a ‘rain’ of dislocations moving along the twin planes is observed. This might indicate the onset of the martensite variant reorientation. The above described collective movement of dislocations contributes to the high IF in martensite. Owing to the applied stress during the IF measurement these dislocations would move along twin planes, causing a high IF response. It should be mentioned, that there are still some dislocations that are pinned, which remain immobile, shown for example by the black marker R in Fig. 5, even at maximum deformation during the experiment. Reversible movement of PDs: effect of stress relaxation

Figure 6a–d shows micrographs extracted from the sequence recorded on a computer during in situ TEM straining experiment. A movement of dislocations indicated by arrows towards the bottom of the images can be observed. The gliding planes of the dislocations can be seen as white, vertical traces. Figure 6e–h was recorded when the strain rate was stopped and the sample started to relax. The dislocations

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activated during stress application moved backward along the same trajectory. During the initial stage of relaxation the dislocation in Fig. 6e and f is moving slowly backwards and then rushes back to its initial position at the twin boundary as indicated by the arrow in Fig. 6g. In the same figure already the next relaxing dislocation appears. Although a large activity of dislocations under stress and during the relaxation period could be observed, the variant interfaces remained immobile (see arrow in Fig. 6).

Discussion The TEM observations of dislocations movement under low applied stress can be compared with the temperature influence on the Q{1 Tr for the following reason. The changes in temperature below Mf cause stress in the martensite, due to the thermal expansion or contraction coupled to the anisotropic evolution of the elastic constants below Mf.33 During the in situ TEM experiment the stress is imposed by tensile deformation of the martensitic NiTiCu polycrystal. As the stress applied in the TEM experiment activates dislocation movements as the stress resulting from temperature changes, the influence of the stress on the microstructure observed in the TEM can therefore be compared with the influence of the temperature during the measurements of Q{1 Tr . During the measurement of the T˙ effect the temperature rate is stopped and an important decrease of Q{1 Tr until its disappearance is observed (Fig. 3). The IF decreases fast, just after stopping the temperature rate, until it reaches the level corresponding to the Q{1 Int , and then remains constant. When the temperature rate starts to change, it takes two or three degrees until the initial level (Q{1 Tr ) is reached again. The same effect was observed in the behaviour of dislocations under application of tensile




a t50 s; b t52 s; c t516 s; d t528 s 5 Transmission electron microscopy weak beam micrographs of moving PD with gliding planes nearly perpendicular to observation direction: a–c dislocations appear as growing loops (indicated by the black arrow) between pinning points, d loop depinns at one point and continues to move stepwise as stress increases and e schematic representation

stress in the TEM experiment. The dislocations were depinned from the variant interface and started to move forward along the twin boundary as shown in Fig. 6a–d. When the deformation was stopped, the dislocations moved back immediately along the same trajectory. However, the longer the stress remained constant, the less dislocations moved back and some of them even stopped halfway between the variant interfaces. When additional strain is applied, the dislocations are not immediately reactivated. It appears then that the movement of dislocations due to stressing and relaxation is a reversible process as long as an important plastic deformation is not induced. During in situ TEM tensile experiments a large quantity of dislocations was observed in the martensitic phase, which became mobile when stress was applied to the specimen. In the same conditions, in terms of microstructure, and assuming equivalent thermally generated stresses relatively high values of the Q{1 Tr were observed. Indeed, according to the authors’ finite element modelling18 elastic stresses in the thin foil at the edges parallel to the acting force can reach at maximum 1 GPa during straining. During the experiment the sample was observed a few minutes after the straining was stopped as it induces too much image drift. It can therefore be assumed that the stresses remaining in the sample are a fraction of the 1 GPa, i.e. y100 MPa. In the IF measurement the thermal stresses are due to thermal expansion for a temperature difference of y100 K over a length of 40 mm (the specimen). These stresses are in the order of 66 MPa, which means that their value is comparable to the stresses generated in the TEM experiment. The model of Dejonghe19,20 is expressed by the following relationship : ( " 3 #) dn : T k 2 : : dn sc {1 {1 dT z 1{ QTr zQPT ~ s0 v J 3p ds s0

where Q{1 PT is the phase transition contribution to IF. This model predicts that during the martensitic transformation the Q{1 Tr has two contributions: the Ln=Lt, which comes from the transformed volume fraction n of martensite in time, and the second one Ln=Ls originating from the transformed volume fraction n of martensite due to stress changes. Moreover, one should consider the non-negligible contribution of the individual movement of twin dislocations. Below the Mf temperature, Ln=Lt and Ln=Ls do not contribute any more, therefore the elevated values of Q{1 Tr should not be observed anymore. Nevertheless, the present study shows that the internal stresses generated by the temperature changes imposed during the IF measurement are present in the anisotropic martensite and therefore cause a movement of the twin dislocations. Indeed, they could give a nonnegligible contribution to the Q{1 below Mf. Tr Consequently, it seems that the Q{1 is sensitive to the Tr density of activated, mobile dislocations in the B19’ phase. The activation originates from local stress concentrations causing the depinning of dislocations. This mechanism is clearly observed during in situ TEM experiments (Figs. 5 and 6). Additionally, it appears to be a reversible process, because during relaxation the activated dislocations move back to the variant boundaries. Therefore, the relaxation process is directly related to the dislocation movement in the B19’ phase of NiTiCu. It should be noted that the intrinsic contribution of 17,30,31,34 the IF (Q{1 to originate from Int ), is known the movement of the variant interfaces. It results from the collective activation of dislocations, which form the variant interfaces, by the measurement strain during the IF experiments. In another study,9,10 it was shown that the displacement of martensite/martensite interfaces together with the motion of PDs observed in situ in the TEM could be related to the application of an


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a t50 s; b t51?1 s; c t53 s; d t53?1 s; e t525 s; f t525?05 s; g t525?1 s; h t526 s 6 Transmission electron microscopy weak beam micrographs of PDs moving along twin boundaries a–d under application of stress movement of dislocation towards bottom of image is observed and e–h during relaxation movement with opposite direction is observed

external stress and the creation of a local elastic (internal) stress.

Conclusions In situ TEM observation is an excellent technique to explain details at the nanometre level of the IF evolution in the martensitic phase of NiTiCu. In situ TEM clearly showed that when stress is applied dislocations are emitted from the twin interfaces. These mobile dislocations stop moving when the applied stress is maintained constant and move back to their initial position when the stress is released. Assuming that the change in temperature below Mf provokes also stress and subsequent emission of dislocation from the twin boundary, when maintaining T constant the dislocations would return to their boundary of origin explaining the Q21 evolution measured by IF as there is less mobile dislocations when T is maintained constant. The present study shows that the IF signal finds its origin in edge dislocations emerging from the interfaces.

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Acknowledgements The authors are thankful to the European Commission for the financial support of the project through the Brite/ EuRam project ‘ADAPT’, no. BE 97 – 4134. The Paul Scherrer Institute is acknowledged for the overall use of facilities.

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