Mechanically induced disorder and crystallization process in Ni-Mn-In ball-milled alloys

1 Mechanically induced disorder and crystallization process in Ni-Mn-In ballmilled alloys V. Sánchez-Alarcos1,2, V. Recarte1,2, J. I. Pérez-Landazábal1,2, S. Larumbe1,2, R. Caballero-Flores1,2, I. Unzueta3,4, J.A. García4,5, F. Plazaola3 and J. A. RodríguezVelamazán6,7 1Departamento Física. Universidad Pública de Navarra, Campus de Arrosadia, 31006 Pamplona, Spain 2Institute for Advanced Materials (INAMAT), Universidad Pública de Navarra, Campus de Arrosadía, 31006 Pamplona, Spain 3Elektritate eta Elektronika Saila, Euskal Herriko Unibersitatea UPV/EHU, p.k. 644, 48080 Bilbao, Spain 4BC Materials (Basque Centre for Materials, Applications and Nanostructures), 48080 Leioa, Spain 5Fisika Aplikatua II Saila, Euskal Herriko Unibersitatea UPV/EHU, p.k. 644, 48080 Bilbao, Spain 6 Instituto de Ciencia de Materiales de Aragón, Departamento de Física de la Materia condensada, CSIC–Universidad de Zaragoza, E-50009 Zaragoza, Spain 7 Institut Laue Langevin, 71, Avenue des Martyrs, 38042 Grenoble Cedex, France *Corresponding author: Tel.: +34 948 169582; Fax.: +34 948 169565 E-mail address: [email protected] Abstract. High mechanical deformation has been induced in a Ni-Mn-In metamagnetic shape memory alloy by means of ball milling. The evolution of both the martensitic transformation and the magnetic properties associated to the microstructural variations has been characterized. The as-milled nanometric particles display an amorphous structure with a frustrated magnetic state compatible with a canonical spin-glass. On heating, an abrupt crystallization process occurs around 500K leading to a cubic B2 structure, which, in turn, does not show martensitic transformation. Modified Arrott plots point to competing longand short-range magnetic couplings in the B2 structure. On further heating, a relaxation process takes place above 700 K concurrently with a B2-L21 atomic ordering, giving rise to an anomalous two-step thermal expansion. The combined effect of both processes makes possible the subsequent occurrence of a martensitic transformation, which takes place at the same temperature than in the bulk. The large relativecooling-power linked to the magnetocaloric effect at the martensitic transformation in the annealed powder makes it interesting for practical applications of magnetic refrigeration at nanoscale. Keywords: Ni-Mn-In, ball-milling, crystallization, atomic order, nanoparticles This document is the Accepted Manuscript version of a Published Work that appeared in final form in Journal of Alloys and Compounds 689(1) : 983-991 (2016). To access the final edited and published work see https://doi.org/10.1016/j.jallcom.2016.08.068 © 2016 Elsevier under CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/) 2 1. INTRODUCTION Ni–Mn-based Heusler alloys exhibiting both long-range magnetic ordering and thermoelastic martensitic transformation (MT) are being intensively investigated over recent years, from both fundamental and applied points of view, due to the unique properties they show linked to the occurrence of a first-order structural transformation between magnetically ordered phases [1-3]. The magnetism in these alloys mainly arises from the coupling between the Mn atoms, in which the magnetic moment is chiefly confined, so the magnetic exchange interactions (which can be treated in the framework of a Ruderman-Kittel-Kasuya-Yoshida model) strongly depend on the Mn–Mn distance [4, 5]. Therefore, different sequences of magnetostructural transformations can be observed depending on both the third alloying element and the change in interatomic distances caused by the MT. In particular, in Ni-Mn-Z (Z=In, Sn, and Sb) alloys, the so-called metamagnetic shape memory alloys, the MT takes place between a ferromagnetic austenite and weak magnetic martensitic phase, in such a way that the large magnetization drop occurring at the MT allows the induction of the MT by an applied magnetic field [6]. This phenomena gives rise to multifunctional properties, namely magnetic shape memory, giant magnetoresistance and large inverse magnetocaloric effect (MCE), of great technical interest for practical applications in sensing and magnetic refrigeration [7-14]. One of the main drawbacks of the metamagnetic shape memory alloys, in turn, is the poor mechanical properties they show. All these alloys are very brittle and difficult to handle, and cracks form very easily when the crystal is being thermally cycled through the phase transformation [15, 16]. Alloys in form of films, ribbons, wires or foams [17-20] have been studied as an alternative to overcome the mechanical limitations of the bulk material. In particular, the integration of powder alloys into a polymer to form composites has been found to be especially interesting for the development of magnetically-controlled dampers and actuators. Such composites may show low cost, large mechanical energy absorption and good mechanical properties [21, 22]. In this case, a previous complete characterization of the powder alloy is needed in order to properly tune the response of the composites. 3 The magnetostructural properties of metamagnetic shape memory alloys have been deeply studied in polycrystalline bulk alloys. Nevertheless, up to now, much less work has been devoted to the analysis of the MT and the magnetic properties at reducing sizes tending to nanoscale. Even though the grain size and the state of internal tensions are known to be two parameters highly influencing the characteristics of the MT through the variation of the elastic energy term [23]. In this sense, it has been recently proposed that below a critical size, shape memory alloy nanoparticles show non hysteretic behavior [24]. Taking into account that the transformation hysteresis linked to a first order transition is an important drawback for magnetocaloric applications, nanoparticle alloys seem to be a promising alternative procedure to improve simultaneously the mechanical behavior and the MCE required for refrigeration systems. Ball milling technique is one of the simplest and cheapest methods to produce nanostructured materials. The use of ball milling as a reduction method to nanoscale size has been studied in multiple systems, and interesting physical and chemical phenomena have been found to appear. A considerable modification of the magnetic behavior and the synthesis of new phases have been reported [25-27]. For instance, the reduction in the size of particles by ball milling may optimize the MCE in different magnetic systems, either by increasing the temperature range [28] or by reducing the hysteretic losses linked to a magnetostructural first order transformation [29]. On the other hand, together with grain size reduction, ball milling produces a huge amount of defects (vacancies, dislocations, chemical disorder, strains,…) in the material. A proper analysis of the milled powder may therefore provide valuable information concerning the influence of those defects on both the MT and the magnetic properties. The effect of ball milling on the magnetostructural properties has been studied in Ni-Mn-Ga, where the effect of thermal treatments has been widely characterized and different sequences of structural transition (different from those in the corresponding bulk) have been found in the achieved nanoparticles [30-33]. Likewise, nanoparticles of Ni-Mn-Sn have been recently obtained by ball milling [34, 35]. As in Ni-Mn-Ga, atomic disorder and lattice strain inhibit the MT in the as-milled alloys, while postmill annealing treatments are needed to restore the MT. A significant influence of annealing on 4 the ferromagnetic exchange interaction, exchange bias and MCE, has been also reported in these alloys. On the contrary, the use of ball milling in the quest for functional nanoparticles has been scarcely studied in the Ni-Mn-In alloys [36, 37]. In fact, a complete study on the correlation between microstructure and magnetostructural properties in Ni-Mn-In nanoparticles is still absent. This work is devoted to the study of the evolution of both the martensitic transformation and the magnetic properties with microstructure in ball-milled Ni-Mn-In nanometric particles. It is found that the as-milled sample display an amorphous structure with a frustrated magnetic state compatible with a canonical spin-glass. The crystallization process taking place on heating the milled alloy and the magnetic properties of the resulting B2 structure have been characterized. A relaxation process occurs on annealing above 700 K, concurrently with a B2-L21 atomic ordering, giving rise to an anomalous two-step thermal expansion. The combined effect of both processes makes possible the subsequent occurrence of a martensitic transformation, which takes place at the same temperature than in the bulk. The large relative cooling power linked to the magnetocaloric effect at the martensitic transformation in the annealed particles makes them interesting for practical applications of magnetic refrigeration at nanoscale. 2. EXPERIMENTAL A Ni50Mn34In16 alloy was prepared from high purity elements by arc melting under protective Ar atmosphere. The as-cast ingot was homogenized at 1073 K during 15h and then quenched into iced water. The composition of the elaborated alloy was analyzed by EDS in a Jeol JSM-5610LV Scanning Electron Microscope (SEM). The alloy was subjected to a ball milling process in an argon atmosphere during an effective milling time of 40h at room temperature. In order to prevent overheating of the sample and grinding jars, the samples were milled during 5 minutes (effective milling time) and then stopped for 10 minutes, and so on. Ball milling was performed at 300 rpm using a Retsch PM4 with a ball:powder ratio of 7:1 and 7 balls of 5 mm diameter. Both jars and balls were of tungsten carbide. The structural transformations and the possible recovery processes 5 taking place on annealing the as-milled powder samples were characterized by differential scanning calorimetry (Q-100 DSC, TA Instruments). The magnetic characterization (magnetization and susceptibility measurements)) was performed by SQUID magnetometry (QD MPMS XL-7). Particles size and shape were observed by conventional transmission electron microscopy (Hitachi H600 100 kV TEM). The evolution of crystal structure, long-range atomic order, crystallite size and microstrains were determined from in-situ powder neutron diffraction measurements performed on the high-flux D1B two-axis diffractometer, at the Institute LaueLangevin (Grenoble, France). The diffraction patterns were measured on heating from room temperature (RT) up to 1173K at 1K/min using a neutron wavelength of 1.28 Å. The structures were refined by the Rietveld method using the FullProf package programs [38]. 3. RESULTS AND DISCUSSION Figure 1a shows the DSC thermogram performed on cooling/heating the reference bulk alloy between 350K and 165 K. The occurrence of a first-order MT is evidenced by the presence of the exothermic and endothermic peaks corresponding to the forward and reverse MT, respectively. The transformation temperatures, taken as the temperature of the peak maximum, are TMforward = 226 K and TMReverse = 241 K, which means a therrmal hysteresis of ∆T = 15 K for the MT. A second-order magnetic transition taking place in the austenitic phase can also inferred from the baseline inflection observed above the MT, around 300 K. The complete sequence of magnetostructural transitions have been determined from the low-field (100 Oe) magnetization versus temperature curve, M(T), shown in Figure 1b. The magnetization increases on cooling below 300 K, due to the ferromagnetic ordering of the austenitic phase at the Curie temperature, TCaust, and then it suddenly decreases on further cooling to the almost paramagnetic martensite. The subsequent increase of magnetization is linked to the magnetic ordering of the lowtemperature phase at TCmart. As shown in the inset, the magnetization drop associated to the martensitic transformation occurs at lower temperatures on increasing the applied magnetic field, H, according to the Clausius-Clapeyron equation 6 (1) where TM is the MT temperature, and ∆M and ∆S the magnetization and entropy change at TM, respectively. Since the different exchange interactions in each structural phase promote a discontinuity in the magnetic entropy at the MT temperature, the application of a magnetic field at temperatures close to TM may result in the induction of the transformation, and therefore in a magnetically induced entropy change (that is, in a MCE). The entropy change in isothermal conditions can be calculated by numerical integration of the derivative of the magnetization with respect to temperature following the expression (2) The obtained MCE values are shown in Figure 1c as a function of both temperature and applied magnetic field. A positive peak (inverse MCE) is observed linked to the MT, which increases with the increasing magnetic field up to a maximum value of 12 J/kgK at 6 T, a value similar to that found in the literature for this compound [12]. Once characterized, the reference bulk alloy was subjected to high energy ball-milling in order to produce nanoparticles. No martensitic transformation was detected from the DSC thermograms performed on the as-milled sample, suggesting a critical impact of the ball-milling on the crystallographic structure of the alloy. Figure 2a shows the room temperature neutron diffraction pattern of the as-milled samples. The absence of well-defined Bragg reflections seems to confirm the achievement of an almost amorphous phase, as a result of the structural distortion and the associated introduction of structural defects caused by milling. In fact, the presence of some small and broad peaks would indicate some marginal degree of crystallinity in the alloy. As shown in the TEM image in the inset, the milled powder mainly consists on spheroidal nanometric particles (tens of nm) together with some bigger clusters. In this respect, the slight peaks on the diffractogram could be also thought to be due to crystalline order inside nanoparticles, that is, to size effects on crystalline powder. As we will see, the clear occurrence of a crystallization process discards this latter hypothesis. The magnetic properties of the as0 M dT M dH S µ D =- D ( ) ( ) 0 ,,0 H iso H M S S T H S T dH T ¶ æö D=-= ç÷ ¶ èø ò 7 milled alloy have been analyzed from the temperature dependence of magnetization and AC susceptibility measurements. Figure 2b shows the zero-field cooled/ field cooling/ field heating (ZFC/FC/FH) curves obtained on the as-milled sample under 100 Oe. The splitting between the ZFC and FC curves at low temperature (sharp jump in ZFC), typical of frustrated magnetic systems, suggests a spin-glass state, as recently proposed in similar alloys [36]. As shown in the upper inset, the low magnetization values at 6 T and the high magnetic anisotropy (the system does not saturate) are indeed in agreement with a spin-glass behavior. Furthermore, a frequencydependent maximum linked to the freezing temperature (temperature of the maximum in the ZFC curve, Tf), appears in the measurements of the real part of susceptibility. The corresponding frequency dependence of the freezing temperature, Tf(f), perfectly fits to the Vogel-Fulcher law [39], again in agreement with a spin-glass state. Further measurements delving into the low temperature magnetic properties of the milled alloys point to a canonical spin-glass state [40]. The microstructural evolution and the recovery processes taking place on heating the amorphous sample have been studied by means of ‘in-situ’ powder neutron diffraction experiments. Figure 3a shows the thermodiffractograms obtained on heating from RT up to 1220 K at a 1 K/min constant rate. It can be seen that the amorphous structure remains stable up to 500 K. At this temperature crystallization occurs, as evidenced by the sudden appearance of Bragg reflections. The occurrence of this process does indeed confirm the amorphous-like state of the as-milled alloy. Specifically, the sample crystallizes to a disordered B2 cubic structure with nearest-neighbors long-range atomic order. On heating around 700 K, the appearance of L21 superstructure reflections (those in which the h, k, l indices are all odd, according to the corresponding unit cell structure factor) indicates the occurrence of an ordering process from B2 to a L21 cubic structure with next-nearest-neighbors long-range atomic order. A similar B2-L21 ordering process is observed on heating bulk Ni-Mn-In alloys previously quenched from high temperatures. In that case, in turn, the ordering temperature (which indeed depends on the degree of disorder retained by quenching) typically lies between 500 K and 600 K [41, 42]. On further heating above 900 K, the intensity of the L21 reflections gradually decreases until it completely 8 vanishes as a consequence of the L21-B2 order-disorder transition, which takes place at 1050 K, a temperature again almost 100 K higher than in bulk alloys [41, 42]. The temperature dependence of the integrated intensity of the (111) reflection shown in Figure 3b illustrates the evolution of long-range atomic order of the nanometric particles with temperature. The two consecutive ordering processes are clearly identifiable from the appearance and subsequent disappearance of the superstructure reflection. The microstructural relaxation can be also explored from the analysis of the width of the Bragg peaks of the diffraction patterns. In this respect, the integral breadth method has been used to determine the microstructural effects from the analysis of the diffraction peaks’ shape. This method allows us to obtain separately the crystallite size and microstrain contributions to the fullwidth at half maximum (FWHM) of the diffraction peaks from its fitting to a Thompson-CoxHasting pseudo-Voigt profile function. The profile instrumental resolution function used for the analysis of the diffraction patterns has been obtained from the RT diffractogram performed on a Na2Ca3Al2F14 sample (typical calibration compound). As an example, Figure 4 shows the diffractograms obtained for the B2 structure at three different temperatures, together with the respective fits by Rietveld method. It can be seen that the broadening of the reflections significantly decreases (and simultaneously the integrated intensity increases) on heating up to 1173 K, pointing out a relaxation of the microstructure. It is also worth noting the presence of small additional peaks, especially one at around 2θ = 29º, probably ascribed to the appearance of a small amount of manganese oxide. In any case, just a very small mass fraction of the alleged oxide is inferred from the relative intensity of the peaks, so no significant effect are expected neither on the magnetic properties nor on the alloy composition. The evolution of both microstrain and crystallite size with temperature is shown in Figure 5a. The main microstructural relaxation takes place between 700 K and 800 K, where the microstrain abruptly decreases from 0.35% to 0.1%. This is precisely the temperature range in which the B2-L21 ordering process occurs (see Figure 3b). Taking into account that atomic ordering implies atomic diffusion, as long as it is meditated by vacancies, it points to vacancy elimination and/or vacancy-assisted 9 dislocation annihilation as the main recovery processes. On the other hand, the crystallite size increases exponentially with the increasing temperature, from 20 nm at RT (in agreement with TEM observations) to 150 nm at 1100 K, approximately. As expected, the higher growth rate is reached above 800 K, just after the recovery process (microstrain drop). Likewise, it is interesting to note that the temperature dependence of the lattice parameters (Figure 5b) shows a marked change in slope between 700 K and 800 K, concurrently with both the B2-L21 ordering and the relaxation processes, thus giving rise to an anomalous two-step thermal expansion. The crystallization process observed from thermodiffraction has been also analyzed from DSC measurements. Figure 6 shows the thermogram obtained on heating the amorphous powders up to 700 K at a 1 K/min rate (same heating rate than in neutron diffraction experiment). A narrow exothermic peak (P1) overlapped to a much broader peak (P2) is observed between 500 K and 600 K. The temperature range of appearance and the width of the P1 peak are in agreement with the sharp crystallization process revealed by diffraction measurements. In turn, no significant structural variation is detected linked to the P2 peak. Nevertheless, the enthalpy of the whole double-peak (∆H ≈ 70J/g) is similar to that obtained for the crystallization process in similar alloys, so the same origin could be attributable to both peaks. The kinetics of the associated processes has been analyzed using the well-known Kissinger’s method, from which the activation energy of a thermally activated process can be obtained from the variation of the DSC peak temperature (TP) as a function of the heating rate ( 𝜙 ), through the expression (3) where kB is the Boltzmann constant, Ea is the activation energy and B is a constant. The activation energies, determined from the linear fit of ln$ ( 𝜙/𝑇! " ) vs 1/𝑇! shown in the inset of Figure 6, are Ea = 1.7 eV and Ea = 1.8 eV for the P1 and P2 peaks, respectively. 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Stanley, Introduction to phase transitions and critical phenomena, London Oxford University Press, London, 1971. [46] J. S. Kouvel, M.E. Fisher, Detailed magnetic behavior of Nickel near its Curie point, Phys. Rev. 136 (1964) A1626. [47] K.A. Gschneidner Jr, V.K. Pecharsky, Magnetocaloric materials. Annu. Rev. Mater. Sci. 30 (2000) 387–429. [48] N.M. Bruno, C. Yegin, I. Karaman, J.-H Chen, J.H. Ross Jr., J. Liu, J. Li, The effect of heat treatments on Ni43Mn42Co4Sn11 meta-magnetic shape memory alloys for magnetic refrigeration, Acta Mater. 74 (2014) 66-84. Figure captions FIG. 1: (a) DSC thermogram performed on cooling/heating the reference bulk at 10 K/min. (b) Temperature dependence of magnetization for the bulk alloy at 100 Oe (inset: M(T) for different applied fields ranging from 100 Oe up to 60 kOe). (c) Magnetically-induced entropy change in the bulk alloys as a function of temperature and applied magnetic field. 18 FIG. 2: (a) Room temperature neutron diffraction pattern of the as-milled samples (inset: TEM micrograph). (b) ZFC/FC/FH curves obtained on the as-milled sample under 100 Oe (Upper inset: field dependence of magnetization. Lower inset: linear fitting to the Vogel-Fulcher law). FIG. 3: (a) In-situ neutron powder thermodiffractograms on heating between 330 K and 1220 K at 1 K/min. (b) Integrated intensity of the (111) reflection as a function of temperature. FIG. 4: Measured neutron diffraction pattern (dots), calculated profile (full line) and difference between the measured and calculated profiles (dashed line) for the milled alloy at 538 K, 630 K and 1170 K. FIG. 5: (a) Microstrain and crystalline size as a function of temperature. (b) Evolution of lattice parameters (in terms of the B2 cell) with temperature. FIG. 6: DSC thermogram on heating the milled sample at 1 K/min. Inset: plot of ln$ ( 𝜙/𝑇! " ) vss 1/𝑇! for the P1 and P2 peak temperatures. FIG. 7: Temperature dependence of the magnetization (cooling curves) at 100 Oe for milled samples heated up to 538 K and 630 K (just above P1 and P2 peaks, respectively). Inset: corresponding M(T) curves under 60 kOe applied field. FIG. 8: Magnetic-field dependence of magnetization between 175 K and 275 K for applied fields up to 60 kOe for (a) sample heated up to 538 K and (b) sample heated up to 630 K. FIG. 9: Temperature dependence of spontaneous magnetization (MS) and inversely initial susceptibility (χ0-1) together with their fitting to equations (4) and (5), respectively. (a) Sample heated up to 538 K, (b) sample heated up to 630 K. 19 FIG. 10: Kouvel-Fisher plots for the MS(T) and χ0-1(T) data. (a) Sample heated up to 538 K, (b) sample heated up to 630 K. FIG. 11: (a) Temperature dependence of magnetization at 100 Oe for the nanometric particles heated up to 1170 K (inset: M(T) for different applied fields ranging from 100 Oe up to 60 kOe). (b) Magnetically-induced entropy change in the milled alloy heated up to 1170 K alloys, as a function of temperature and applied magnetic field.