Effect of heat treatment on magnetostructural transformations and exchange bias in Heusler Ni48Mn39.5Sn9.5Al3 ribbons

Acta Materialia 103 (2016) 30e45

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Effect of heat treatment on magnetostructural transformations and exchange bias in Heusler Ni48Mn39.5Sn9.5Al3 ribbons P. Czaja a, *, M. Fitta b, J. Przewo
znik c, W. Maziarz a, J. Morgiel a, T. Czeppe a, E. Cesari d
w, Poland Institute of Metallurgy and Materials Science, Polish Academy of Sciences, 25 Reymonta Str., 30-059, Krako The Henryk Niewodniczanski Institute of Nuclear Physics, Polish Academy of Sciences, 152 Radzikowskiego Str., 31-342, Krakow, Poland c AGH University of Science and Technology, Faculty of Physics and Applied Computer Science, Department of Solid State Physics, Al. Mickiewicza 30, 30-059 Krakow, Poland d Department de Fisica, Universitat de Illes Balears, Ctra. de Valldemossa, km 7.5, Palma de Mallorca E-07071, Spain a

b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 14 May 2015 Received in revised form 1 October 2015 Accepted 2 October 2015 Available online xxx

The influence of low temperature annealing and high temperature quenching on martensitic transformation, microstructure, magnetic and exchange bias properties of Ni48Mn39.5Sn9.5Al3 metamagnetic shape memory alloy ribbons have been investigated. It is shown that with increasing thermal treatment temperature the martensitic transformation temperature increases and the exchange bias is reduced. In the case of low temperature annealing this effect is mediated by stress and structural relaxations, whereas in the case of high temperature annealing followed by water quenching this effect is ascribed to the combined influence of grain size enlargement and composition change resulting from metastability of the NieMneSneAl phase at elevated temperatures. Furthermore low temperature thermal treatment leads to the change in the martensite structure from 4O to 10M as is demonstrated by in situ TEM studies. It is demonstrated that low temperature annealing is a preferable method for fine tuning magnetostructural properties of NieMneSneAl ribbons. © 2015 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Metamagnetic shape memory alloys (MSMA) Martensitic transformation Exchange bias Transmission electron microscopy

1. Introduction NieMneX (X ¼ Ga, In, Sn, Sb) Heusler based magnetic shape memory alloys have recently attracted much interest, which is fuelled by their wide applications prospects including actuation, sensing [1], energy harvesting and environmentally friendly magnetic refrigeration [2]. Other interesting properties comprise exchange bias (EB), attributed to the ferromagnetic (FM) e antiferromagnetic (AFM) interfaces present in an EB material, with its potential in spin electronic devices [3]. In general, multifunctional capacity of these systems is based on the coupling between structural and magnetic degrees of freedom resulting in conjunction with the first order, diffusionless, solid e solid and thermoelastic martensitic transformation (MT). In NieMneX alloys MT can be induced by temperature, stress or magnetic field. It goes from the high temperature, high symmetry (Fm3m) cubic austenite phase to the low temperature, lower symmetry (Pmma) and very

* Corresponding author. E-mail address: [email protected] (P. Czaja). http://dx.doi.org/10.1016/j.actamat.2015.10.001 1359-6454/© 2015 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

often structurally modulated martensite phase, which depending on composition and manufacturing conditions can have different configurations: L10, 10M, 14M, 4O [4]. NieMneSn system [5,6] together with NieMneSb and NieMneIn have turned out to be a promising class of materials, since they show lower magnetic moment in the martensite phase than in the austenite phase. This is different to the most studied Ni2MnGa ferromagnetic shape memory alloy (FSMA), for which the magnetic moment of the martensite phase is larger than that of austenite [7]. In turn, it gives rise to a unique phenomenon of the magnetic field induced reverse martensitic transformation (RMT) and for distinction the former alloys are termed metamagnetic shape memory alloys (MSMA) [8]. It further entails an inverse magnetocaloric effect (IMCE) associated with magnetic field induced RMT and which has been proposed, rather than for cooling per se, as complementary to the conventional MCE around the second order paramagnetic (PM) e FM transition at the Curie temperature also found in NieMneX alloys [9e20]. However, prior to large scale engineering applications one of the key issues to be addressed is the MT temperature control. MT in NieMneX alloys can operate at a wide temperature range.

P. Czaja et al. / Acta Materialia 103 (2016) 30e45

Yet, it has been established that in NieMneX Heusler alloys it is very sensitive to the valence electron concentration per atom (e/a) [21]. Therefore one way of MT temperatures tuning in these systems can be realised through stoichiometry adjustments or substitution of alloying elements [22,23]. It has also been shown that MT in some NieMn based alloys is critically linked to the shortest distance between Mn atoms (dMn-dMn) [24], likewise magnetic properties, which in NieMneX alloys arise mainly from the MneMn exchange interactions. This stems from the fact that the magnetic moment in these systems is chiefly confined to Mn, and since these atoms are not in direct contact, the exchange mechanism is indirect mediated through the oscillatory Ruderman-KittelKasuya-Yoshida (RKKY) interaction of conduction electrons. Consequently it is strongly affected by change in interatomic distances [24]. Eventually depending on composition and the change in the lattice symmetry imposed by MT a series of magnetostructural transformations can be observed in these systems. Since MT is influenced by structural characteristics its control can be aided by microstructure refinement such as that obtained by melt spinning of ribbons. Melt spinning has been demonstrated to be an effective, single step production technique of NieMneX multifunctional alloys [25]. It allows for avoiding of the prolonged thermal annealing step and it can be operated on continuous, large mass scale. Considerable microstructure refinement, which it offers can decrease both brittleness of intermetallic materials and lower MT temperature [26]. Since rapid solidification from the melt takes place at high cooling rates it can result in positioning of atoms in non-equilibrium sites. Thus, it may also allow for modification of atomic order, which may be further employed for control of MT temperature and multifunctionality of NieMneX based compounds. The configurational ordering of the constituting elements in the crystal lattice can be modified by an appropriate heat treatment as has been demonstrated in other MSMA alloys [27]. Heat treatment experiments involving long time annealing and high temperature quenching have been conducted for a number of melt spun and bulk polycrystalline NieMneSn based alloys, however the results reported still remain controversial [28e32]. And Therefore this subject desires more attention, especially in the light of recent reports concerning thermal instability of NieMneSn alloys [33,34]. Furthermore microstructure evolution and phase composition upon heat treatment in these alloys were scarcely covered. Also the effects of thermal treatment on the EB properties in NieMneSn based alloys have little been investigated [35,36]. The structural and magnetic studies of Ni48Mn39.5Sn12.5 ribbons showed that with Al substitution for Sn up to 3 at % the TM and the temperature of the reverse transition (TRM) increase to the ambient temperature range, whereas the TCA remains unaffected and the Curie temperature of martensite (TCM ) decreases [37,38]. In particular the Ni48Mn39.5Sn9.5Al3 ribbons appear interesting from the fundamental point of view because of the close proximity of TM and TCA temperatures and the recently proposed argument that the atomic order variations affect MT temperature only in those NieMn-based alloys in which at least one of the structural phases shows magnetic ordering at the transition temperature [39]. The objective of the present study is to address the effect of high temperature quenching and low temperature annealing on microstructure, phase composition, the characteristic transformation temperatures, magnetic and EB properties of Ni48Mn39.5Sn9.5Al3 ribbons, what has not yet been investigated for these alloys. Special attention is paid to the characterisation of phase composition and microstructure evolution using transmission electron microscopy (TEM), including in-situ TEM heating experiments permitting to directly follow the analysed phase transformations.

31

2. Experimental Ni48Mn39.5Sn9.5Al3 ribbons were fabricated by melt spinning technique by ejecting molten alloy with argon overpressure (0.25 MPa) onto a surface of a copper wheel rotating at a linear speed of 25 m s1. The process was carried out under an argon atmosphere. More detailed information on ribbons manufacture can be found elsewhere [37]. In order to evaluate the effect of high temperature treatment on the microstructure and the magnetic properties four ribbon samples were sealed in quartz ampoules filled with argon gas and subsequently annealed for 1 h at four different temperatures, namely 873, 973, 1073 and 1173 K, what was followed by water quenching (WQ) by breaking the ampoules in water container at room temperature. The four ribbon samples are hereafter referred to as AWQ873, AWQ973, AWQ1073 and AWQ1173, respectively. Thermal effects of the high temperature quenching were investigated using a Mettler DSC 823 instrument in the 173e423 K temperature range. Low temperature annealing at 573 K for various lengths of time, ranging from 5 to 320 min, was performed under protective argon gas atmosphere on a Q1000 DSC TA Instrument, which was then used to study the effects of this treatment. The low temperature annealed sample for in overall 635 min is hereafter referred to as A573. The cooling/heating rate used throughout DSC experiments was set at 10 K/min. The microstructure and chemical composition were characterized by Scanning Electron Microscopy (SEM) using FEI Quanta 3D FEGSEM, FEI E-SEM XL30 equipped with an X-ray energy dispersion spectrometer EDAX GEMINI 4000 and with Transmission Electron Microscopy (TEM) employing Tecnai G2 operating at 200 kV equipped with an Energy Dispersive X-ray (EDX) microanalyser and a High Angle Annular Dark Field Detector (HAADF). In situ TEM studies were carried out with the heating rate set at 5 /min. Each recording was taken at an isothermal stop, which lasted no longer than 60 s. Thin foils for TEM examination were prepared by Focused Ion Beam (FIB) technique using FEI QUANTA 3D Dual Beam and with TenuPol-5 double jet electropolisher using an electrolyte of phosphoric acid (20%) and ethanol (80%) at 243 K. X-ray diffraction (XRD) patterns were collected with a Philips (PW1710) diffractometer using Co Ka radiation (l ¼ 1.78896 Å). The data obtained were analyzed using the profile fitting program FullProf based on the Rietveld method [40]. (The background intensity was approximated with a polynome and the peak shape was fitted with a pseudo-Voigt function). The DC mass magnetic susceptibility was measured in the temperature range from 2 K up to 380 K and in the magnetic field of 50 Oe using the Vibrating Sample Magnetometer (VSM) option of the Quantum Design Physical Property Measurement System (PPMS-9). Measurements of AC susceptibility were performed at Lake Shore 7225 AC susceptometer/DC magnetometer. AC susceptibility was measured with the frequency of the oscillating field set to 80 Hz, 240 Hz, 600 Hz, 1200 Hz and 10000 Hz and the amplitude of 1 Oe. Both real c0 and imaginary c” components were determined. For all the thermomagnetic measurements, the ribbons were crushed and pressed in cylindrical polypropylene sample holders. 3. Results 3.1. Martensitic transformations DSC curves obtained for the AWQ873, AWQ973, AWQ1073 and AWQ1173 samples show that on cooling all the ribbons undergo the first order MT and on heating RMT, what is manifested by the corresponding exothermic and endothermic peaks, respectively

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P. Czaja et al. / Acta Materialia 103 (2016) 30e45

(Fig. 1). It is also seen that with increasing annealing temperature the characteristic forward and reverse transformation temperatures increase, however this occurs only up to the 1073 K quench temperature range. Further increase in quenching temperature up to 1173 K brings no significant change to the transformation temperature, but instead it results in broadening of the DSC peaks obtained both on cooling and on heating. This is well visible for the AWQ1073 and in particular for the AWQ1173 sample. This effect may be ascribed to some decomposition processes taking place at this elevated temperature range and it is further disclosed in the following discussion. It is also seen from Table 1 that on contrary to increasing transformation temperatures the width of the thermal hysteresis does not seem to be affected by the increasing quench temperature regime. The values of the MT and RMT temperatures along with the width of the thermal hysteresis (DT) and the values of the transformation enthalpy (DH) and entropy (DSt) changes determined from the calorimetric curves presented in Fig. 1 are also listed in Table 1. The values of DSt for the MT (DStA/M) and RMT (DStM/A) are determined from the calorimetric curves according to the following relations: DStA/M ¼ QA/M/TpA/M; DStM/A ¼ QM/A/ TpM/A, where QA/M, QM/A, TpA/M, TpM/A are the corresponding transformation heats and peak temperatures, respectively [41]. It should be noted that entropy changes calculated in this manner do not deviate appreciably from the values determined by integrating dQ/T over the entire transformation range [42]. It is then noticed that the transformation entropy change values for both forward and reverse MT appear to increase with increasing quenching temperature. In the case of the as melt spun, AWQ873 and AWQ973 samples this increase continues progressively but then it drops for AWQ1073 sample and picks up again for AWQ1173 sample. This behaviour is well illustrated in Fig. 2, which shows the dependence of DStA/M on the relative position of TM and TCA temperatures, and from this perspective it may be well understood in relation to the magnetic contribution to the Gibbs free energy when these two temperatures coincide. According to the observed results the increasing quenching temperature produces an increase in TM with minimal impact on TCA in the case of the AWQ873 and AWQ973 samples, therefore in this instance the increase in transformation entropy change may be due to the shortening distance between the structural and magnetic transitions. On the other hand a decrease in DSt witnessed for AWQ1073 sample as compared to AWQ873 and AWQ973 samples may be associated with the increase of TCA for this sample, suggesting ordering effects or composition change. A

sudden jump of DSt for AWQ1173 sample despite even greater than previously increase of TCA may be related to the decomposition of the primary NieMneSneAl phase resulting in the appearance of additional transformable phases as evidenced by the shape reversibility of the DSC peak measured for this ribbon sample. DSC thermo-grams obtained for the same melt spun alloy ribbons after annealing at 573 K for various times ranging from 5 to 320 min are shown in Fig. 3. It is observed from this figure that the ribbon regardless of heat treatment duration undergoes MT and RMT. However, with increasing annealing time the characteristic transformation temperatures increase. This is well illustrated in Fig. 4(a), which shows the martensite start (Ms), martensite finish (Mf), austenite start (As), austenite finish (Af) temperatures change as a function of annealing time. It is worth noting that the structural transformation temperatures increase more significantly with annealing time extended up to 75 min, after which time additional annealing brings no significant benefits in terms of TM increase and even some kind of saturation limit is reached. The characteristic hump corresponding to the Curie temperature of austenite (TCA ) is not visible on the thermograms (Fig. 3) due to the close proximity between MT and TCA temperatures. The TCA is contained within the large peaks linked to the MT and RMT. Fig. 4(b) shows the transformation entropy (DSt) change vs. annealing time. The MT (DStA/M) and RMT (DStM/A) entropy changes were determined in the same way as in the preceding paragraph. The values of the respective peak temperatures (TpA/M, TpM/A), the width of DT hysteresis and values of enthalpy DH and transformation entropy change upon cooling (DStA/M) and heating (DStM/A) are summarised in Table 2. It is then clearly visible from Table 2 and from Fig. 4(b), which graphically depicts DStA/M and DStM/A dependence on annealing time, that with increasing annealing time both the DStA/M and DStM/A decline. This effect is more pronounced within the initial annealing time interval between 5 and 75 min. Again, as previously, it may be linked to the relative position of TM and TCA temperatures. The latter temperature determined according to thermo-magnetic measurements for the A573 samples annealed for 635 min is 307 K, whereas the as melt spun sample undergoes PM-FM transition at 299 K. Therefore it is reasonable to expect that during the first annealing period for which most significant transformational entropy drop occurs, the Curie temperature of austenite dramatically increases, hence the large temperature interval between TM and TCA giving rise to an abrupt entropy decrease. TM increase is on the other hand more continuous. After the first annealing period TCA value levels off showing no further changes with prolonged annealing time, which is similar to the behaviour of the structural transformation temperature within the same time framework. This effects may point out amongst others to microstrain relaxation and enhancement of the degree of atomic order [43]. Other effects such as grain growth, composition change affecting the e/a ratio and the change in the dMn-dMn interatomic distance should also be taken into account [21,22]. Therefore in the subsequent discussion the potential influence of all of these factors in relation to the current study is briefly reviewed in order to arrive at a clearer picture of the impact of the employed thermal treatment effects on magneto-structural properties of the investigated Ni48Mn39.5Sn9.5Al3 ribbon alloys. The atomic site ordering may have significant impact on the magnetic properties, which relates to the change in the dMn  dMn interatomic distance. In order to evaluate the effect of heat treatment on the magnetic properties of the Ni48Mn39.5Sn9.5Al3 ribbons thermomagnetization measurements have been conducted. 3.2. Thermomagnetization measurements

Fig. 1. DSC thermograms obtained for the as melt spun, AWQ873, AWQ973, AWQ1073 and AWQ1173 ribbons. EXO UP[.

The zero field cooled (ZFC), field cooled (FC) and field heated

P. Czaja et al. / Acta Materialia 103 (2016) 30e45

33

Table 1 Forward (TpA/M) and reverse (TpM/A) MT temperatures given as the corresponding peak temperatures, thermal hysteresis (TpA/M  TpM/A), the Curie temperature of martensite (TCM ) and austenite (TCA ) phases and the absolute values of enthalpy (DH) and transformation entropy changes (DSt) accompanying MT and RMT for the as melt-spun and AWQ873, AWQ973, AWQ1073, AWQ1173 ribbons determined from DSC and thermomagnetic (VSM) curves. Ribbon

Method

TpA/M (K)

TpM/A (K)

DT (K)

TCM (K)

TCA (K)

DH (J$g1)

DStA/M (J$kg

As spun AWQ873 AWQ973 AWQ1073 AWQ1173

DSC VSM DSC VSM DSC VSM DSC VSM DSC VSM

290 291 297 293 301 e 308 e 306 297

304 300 311 e 315 e 322 e 321 e

14 9 14 e 14 14 e 15 e

e 190 e 195 e 197 e 194 e 194

e 299 e 302 e 302 e 305 e 315

1

$K

DStM/A 1

)

Cooling

Heating

Cooling

Heating

13.9 e 15.5 e 16.8 e 15.9 e 19.2 e

10.9 e 13.3 e 13.1 e 12.7 e 15.2 e

48.2 e 52.4 e 55.9 e 51.7 e 62.9 e

36.0 e 42.8 e 41.8 e 39.5 e 47.5 e

(FH) magnetic susceptibilities (c) of the AWQ873, AWQ973, AWQ1073, AWQ1173 and A573 ribbon samples were measured as a function of temperature and are presented in Fig. 5. Before ZFC measurement the sample was cooled from 370 K to 2 K at zero applied magnetic field. Subsequently the ZFC and FC c(T) curves were recorded at an applied magnetic field of 5 mT on heating and on cooling, respectively. The FH c(T) measurement was finally performed at the same applied field during heating up run. It is seen from Fig. 5 that all the ribbons, upon the temperature variation, undergo a change of magnetic susceptibility typical for the first order martensitic transition and at higher temperature a second order ferromagneticeparamagnetic transition at the Curie temperature of austenite, both of which are frequently observed in these systems. In the case of the as melt spun ribbon and according to the thermomagnetic curves the martensitic transition is noted at 291 K, the reverse martensitic transition takes place at 300 K, whereas the magnetic transition of austenite in this sample occurs at 299 K and the Curie temperature of martensite is equal to 190 K (The TM, TCA and TCM temperatures were determined as temperature corresponding to the inflection points of the FC curve, and TRM was determined as temperature corresponding to the inflection points on FH curve). It is observed from Table 1 that all the characteristic structural and magnetic transition temperatures tend to increase with the quenching temperature. The increase is however less

significant in respect to the TCA of austenite, which then leads to the TpA/M and TCA temperature overlap (Table 1). This effect is well evidenced, with the exception of the A573 sample (Fig. 5(a)), with the lowering magnetic susceptibility c accompanying the martensitic transformation as illustrated by the FC curves for the annealed and WQ samples (Fig. 5(c)-(f)). This then may be ascribed to the modification of the degree of atomic order or composition change resulting from the annealing process. Strikingly it is worth to note that the c values increase for all the heat treated samples as compared to the as melt spun sample (Tables 1 and 2). What is more FC and FH curves splitting, which may be associated with the intermartensitic transition, is observed for the latter sample below 180 K and then subsides for the A573 sample and eventually disappears for the annealed and WQ samples. Due to the fact that the onset of the MT in these samples coincides with the Curie temperature of austenite, the c values of the latter are higher than the maximal c values of martensite below its TCM , however it is known that in NieMneSn systems, largely due to its lower magnetic anisotropy, austenite shows more ferromagnetic behaviour than martensite phase. The magnetic state of martensite, in particular in Mn-rich NieMn alloys, is frequently described as a mixture of FMantiferromagnetic (AFM) states. It is believed that the FM state in Mn-rich NieMneX (X ¼ Sb, In, Sn) alloys stemms from Mn atoms at regular sites, whereas the AFM originates in excess Mn atoms located in Ni or X sites. The AFM-FM interactions are usually

Fig. 2. Absolute value of the MT transformation entropy change (DStA/M) as a function of.TCA  TM .

Fig. 3. DSC thermo-grams obtained for the Ni48Mn39.5Sn9.5Al3 ribbons after annealing at 573 K for various times ranging from 5 to 320 min.

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P. Czaja et al. / Acta Materialia 103 (2016) 30e45

Fig. 4. Characteristic temperatures (Ms, Mf, As, Af) change vs annealing time (a), transformation entropy change (DSt) vs annealing time for Ni48Mn39.5Sn9.5Al3 ribbons (b). Annealing was performed at 573 K.

Table 2 Forward (TpA/M) and reverse (TpM/A) MT temperatures given as the corresponding peak temperatures, thermal hysteresis (TpA/M  TpM/A), the Curie temperature of martensite (TCM ) and austenite (TCA ) phases and the absolute values of enthalpy (DH) and transformation entropy changes (DS) accompanying MT and RMT for the as melt-spun ribbon annealed at 573 K for various lengths of time as determined from DSC and thermomagnetic (VSM) curves. Annealing time (min.)

Method

TpA/M (K)

TpM/A (K)

DT (K)

TCM (K)

TCA (K)

DH (J$g1)

DSt A/M

DSt M/A

(J$kg1$K1)

0

DSC

292

302

10

e

e

Cooling

Heating

10.0

9.2

Cooling

Heating 30.7

34.3 5

DSC

294

304

10

e

e

8.8

9.0

29.6 30.1

15

DSC

296

306

10

e

e

9.5

32.5

9.9 32.1

35

DSC

298

307

9

e

e

9.1

30.0

9.2 30.7

75

DSC

299

309

10

e

e

8.9

8.9

28.8 29.7

155

DSC

300

310

10

e

e

9.1

8.8

28.6 30.3

315

DSC

301

311

10

e

e

9.0

8.9

28.7 29.9

635 A573

DSC VSM

302 303

311 311

9 8

reflected by the splitting between the ZFC and FC curves below TCM , above which one can either find a weakly magnetic, paramagnetic or antiferromagnetic state. Such splitting is clearly observed in Fig. 5 for all ribbons. It may be correlated with the EB phenomenon. EB is referred to the shift of the centre of the magnetic hysteresis loop from the origin when the sample has been cooled from the high temperature in the presence of a magnetic field [44]. However EB has also been observed in materials featuring FM-spin glass (SG) and FM-ferrimagnetic interfaces [45]. What is more some investigators argue that excess Mn atoms couple ferromagnetically and this is why no EB was observed in e.g. Mn2NiGa Heusler alloy. In addition AFM and FM phases should be distanced in at least nanometre scale but in the irregular site occupancy model one considers an angstrom scale [46]. Therefore the existence of EB does not give a definite answer to the question of the nature of magnetic state of martensite. In general, whatever is the character of the exchange coupling, it is crucial that the magnetic transition temperature for the nonferromagnetic (NF) phase (AFM, SG, ferrimagnet) TNF is below TC. It is also known that EB vanishes in Ni2Mn1þxZ1x systems above a certain temperature known as the blocking temperature, TB, which coincides with the foot of a step like anomaly observed in the ZFC vs. T curve. This is due to the faster

e 212

e 307

8.7 e

8.8 e

28.2 28.9 e

e

disappearance of NF states with increasing temperature and preservation of the FM character of the high temperature phase [47]. By examining c(T) dependencies in Fig. 5(a) it may be observed that for the as melt spun sample the ZFC and FC curves begin to separate at around 180 K, the ZFC magnetization shows a maximum at this temperature while FC magnetization shows a monotonic increase below this temperature. This may suggest magnetic inhomogeneity in the martensite phase. The TCM in this sample is detected at 190 K, what allows for assuming that at least a FM phase should be present below this temperature. According to Fig. 5(b) to (f) it is noticed that splitting between ZFC and FC in thermally treated samples appears at a higher temperature of about 200 K, which may than suggest the growth of the FM spin structure in these samples [48]. This effect may also be related to a slight shift of TCM to higher values in these samples. Moreover a distinct step like anomaly in the ZFC curve may be seen for AWQ samples just below 125 K, which is indicative of TB typically observed in off stoichiometric NieMn based alloys at about 100 K [3]. Magnetic hysteresis loops of the studied samples at 3 K measured after ZFC and FC (at 2 T) from 380 K are shown in Fig. 6(a) and (b), respectively. All the M-H loops were recorded in the magnetic field, m0$H, of 1 T and þ1 T. Fig. 6(b) shows the M (H) loop restricted

P. Czaja et al. / Acta Materialia 103 (2016) 30e45

35

Fig. 5. Temperature dependence of the FC and ZFC dc mass magnetic susceptibility (c) measured at external magnetic field 5 mT for the as melt spun Ni48Mn39.5Sn9.5Al3 and A573, AWQ873, AWQ973, AWQ1073 and AWQ1173 ribbons.

to 0.2 T  m0$H  þ0.2 T, full loop is shown in the inset. It is seen from Fig. 6(a) that the ZFC loops are symmetric with double shifted hysteresis, typical for EB materials with different magnetic anisotropy phases. It may also be noticed that the martensite phase in thermally treated ribbons is magnetically softer, the maximal magnitude of magnetization of the martensite phase is enlarged by thermal treatment of the as melt spun ribbon. With further increasing the temperature up to 1073 K, the magnetization drops again as is visible for AWQ1173 sample. On the other hand, it is seen from Fig. 6(b) that after FC the magnetization M-H curves at 3 K for all ribbons are clearly displaced toward the negative field axis, revealing EB phenomenon. However it is seen that the M-H loops corresponding to thermally treated samples shift less than the loops for the as melt spun ribbon, which may suggest reduction in the magnitude of EB effect. The values of EB field, HE, is calculated as HE ¼ jH1 þ H2j/2, where H1 and H2 are the coercive fields of the ascending and descending branches of the hysteresis loop, respectively. The mean value of coercive field Hc is defined as

Hc ¼ jH1  H2j/2. The m0$HE values are 45.3, 26.3, 9.6, 3.5, 14.9, 3.9 mT and the m0$Hc values are 27.5, 20.5, 11, 7.5, 14.5, 9.5 mT for as melt spun, A573, AWQ873, AWQ973, AWQ1073 and AWQ1173 samples, respectively. Taking the above into account and in order to shed more light on the effect of thermal treatment on the ground state of martensite detailed analysis of its magnetic properties was performed. Fig. 7 shows full M (H) loops taken after FC (5 T) at different temperatures below and above TCM . It may be seen that the loops taken at 3, 100 and 170 K, below the Curie temperature of martensite, exhibit a sigmoid shape. Magnetization at these temperatures increases with the magnetic field, initially abruptly but later the increase slows down. In general magnetization is low and it is interesting to note that it does not saturate even at the magnetic field of 9 T. It is also evident that the magnetization declines as the temperature goes up. The magnetization reversal is not smooth and it shows a clear hysteresis. At 3 K the M (H) loop is displaced from the origin towards the negative field axis revealing the EB effect. The loop shift

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P. Czaja et al. / Acta Materialia 103 (2016) 30e45

disappears at 100 K, which signifies that TB is around this temperature. The nonzero coercive field still persists at 170 K, however substantially reduced. At 270 K, which is well below TpA/M M (H) shows neither coercivity nor remanence, which should point to a paramagnetic state of martensite at this temperature range. Similarly the M (H) loop recorded at 300 K, which coincides with TpM/A and TCA , shows no hysteresis, yet it displays a sigmoid like shape suggesting superparamagnetic state of austenite at TM. The magnetization difference between martensite and austenite due to the difference in magnetocrystalline anisotropy accounts for the difference in magnetization as visible from the curves recorded in 270 and 300 K, respectively. Furthermore AC susceptibility measurements at different frequencies were carried out for the as melt spun ribbon within the 9e229 K temperature range to confirm the existence of any SG states in the martensite phase. The results of these measurements are shown in Fig. 8. Fig. 8(a) shows the temperature dependence of the real part c0 and Fig. 8(b) shows the temperature dependence of the imaginary part c00 of the AC susceptibility measured at different frequencies f with an AC magnetic field of 1 Oe after ZFC to 9 K. It may be observed that a peak, Tf, appears in each c0 (T) curve, which is associated with the time of measurement (t) being equal to the relaxation time (t) of the system t ¼ t. However it is also evident that Tf is independent of frequency. The temperature at which it appears, determined as the inflection point from c0 (T) (10 kHz) curve, is 194 K, which coincides with the Curie point of martensite. Therefore since Tf ¼ TCM and no additional frequency dependent features are recognised on both c0

Fig. 6. ZFC (a) and FC (b) hysteresis loops of as melt spun Ni48Mn39.5Sn9.5Al3 and A573, AWQ873, AWQ973, AWQ1073 and AWQ1173 ribbons measured at 3 K.

Fig. 7. Full M (H) hysteresis loops recorded for as melt spun Ni48Mn39.5Sn9.5Al3 ribbon at various temperatures (3, 100, 170, 270 and 300 K) after FC at 5 T at the magnetic field ± 9 T. The inset shows part of the hysteresis loop between ±2 T.

Fig. 8. Real part (a) and imaginary part (b) of the AC susceptibility measured at different applied frequencies of the AC signal for the as melt spun Ni48Mn39.5Sn9.5Al3 ribbons. The peak magnitude of the applied AC field is 1 Oe.

P. Czaja et al. / Acta Materialia 103 (2016) 30e45

(T), c00 (T) curves it allows for assuming that no glassines should be present in the low temperature martensite phase in the studied as melt spun Ni48Mn39.5Sn9.5Al3 ribbon. And hence considering all the above discussion the following sequence of magnetic and structural transition in the studied as spun Ni48Mn39.5Sn9.5Al3 is proposed with decreasing temperature (Fig. 5(a)): initially PM austenite begins to order FM at around 299 K, however due to the close proximity to TM at which it starts to transform into martensite, it may show superparamagnetic behaviour at this temperature interval, subsequently superparamegntic/paramagnetic martensite appears and it then begins to order FM at 190 K. Below this temperature FM domains coexist with AFM/ferrimagnetic clusters. At 100 K the AFM/ferromagnetic moments are responsible for exchange pinning of the FM moments giving rise to EB. TB ¼ 100 K. Overall since in general the EB effect is explained in terms of competing FM and AFM interactions resulting from the RKKY interaction between Mn atoms in the regular Mn sites and excess Mn atoms occupying the Sn sites in the off-stoichiometric NieMneSn Heusler systems, the observed EB changes upon thermal treatment may suggest the possible mechanism by which temperature treatment influences magnetic and microstructural properties of the studied ribbons. During the annealing process micro strain relaxation, grain growth, composition change may modify the atomic order and consequently lead to the change in the exchange coupling in the system. Therefore detailed microstructural studies were undertaken to evaluate the potential influence of compositional and structural effects on the observed behaviour. 3.3. Crystal structure Fig. 9 shows the evolution of the XRD patterns taken at room temperature for the as melt spun and A573, AWQ873, AWQ973, AWQ1073, AWQ1173 samples. The XRD patterns for all the ribbon

37

samples can be successfully indexed according to the mixture of the cubic L21 austenite phase and the martensite phase. The lattice parameters and the relative mass contribution of both phases in all the studied ribbons are listed in Table 3. The dependence on the temperature and the protocol of the thermal treatment of the lattice parameter ac and the relative mass contribution pc of the cubic austenite phase is also schematically illustrated in Fig. 10. It can be seen from both Table 3 and Fig. 10 that the lattice parameter of the parent phase increases significantly after annealing at 573 K for 635 min when comparing it to the as spun ribbon sample. At the same time the relative mass contribution of the austenite phase drops after this treatment to 5.85%, which is the lowest determined contribution of the parent phase at RT among the studied samples. Furthermore it is observed that also the high temperature annealing treatment followed by WQ brings about an increase in the ac, however in this case the increase is less pronounced. Moreover it is observable that initially the ac lattice parameter increases for the AWQ873 and AWQ973 samples but then it begins to decrease as the annealing temperature goes up to 1073 and then to 1173 K. Despite this occurrence the ac parameter still remains higher for AWQ1073 and AWQ1173 samples than the lattice parameter in the as spun ribbon sample. Similarly to the low temperature annealing treatment also the high temperature annealing followed by WQ yields a decrease in the relative mass contribution of the austenite phase when comparing with the untreated, as spun sample. It is evident nonetheless that the pc decrease in this instance is less than in the case of the A573 sample. Annealing at 973 K followed by WQ brings insignificant benefits in terms of further reduction in pc but on contrary elevating the thermal treatment temperature to 1073 and 1173 K results in a considerable increase in the amount of austenite phase, exceeding even the relative mass contribution of the parent phase in the as spun sample, in the AWQ1073 and AWQ1173 ribbon samples. This effect is consistent with earlier DSC results and may point out to decomposition processes arising in response to annealing at elevated temperatures. It is also worth noting that frequently the thermal evolution of peak intensities corresponding to the (111) and (200) super lattice reflection peaks, typical for the L21 Heusler structure, is used to evaluate the influence of the heat treatment on the degree of atomic order of the parent phase [32]. In this instance however this task is somewhat difficult, since all the XRD patterns are recorded at RT at which the martensite phase remains the prevailing phase in the studied samples. This may be seen from Fig. 9. From this figure it is evident that the as melt spun ribbons feature very low intensities of these two characteristic peaks, which may be associated with a low volume fraction of the austenite phase in this sample and with a low degree of atomic order. In general the degree of disorder is associated with the very fast cooling rates in the range of 105e107 K/s imposed by melt spinning, which prohibits atomic ordering and therefore results in many atoms being located in nonequilibrium positions. The suggestion that the low intensity of the

Table 3 Room temperature values of the austenite phase cubic unit cell parameter, ac, and its relative mass contribution, pc, and the martensite phase unit cell parameters: am, bm, cm for the ribbons studied. Ribbon

Fig. 9. X-ray diffraction patterns obtained at room temperature for the as melt spun Ni48Mn39.5Sn9.5Al3 and A573, AWQ873, AWQ973, AWQ1073 and AWQ1173 ribbons.

As-spun A573 AWQ873 AWQ973 AWQ1073 AWQ1173

Austenite

Martensite

ac (Å)

Va (Å3)

pc (%)

am (Å)

bm (Å)

cm (Å)

Vm (Å3)

5.949 5.973 5.957 5.963 5.960 5.957

210.644 213.100 211.459 212.029 211.773 211.430

19.98 5.85 9.50 8.91 29.18 31.73

8.570 8.576 8.582 8.579 8.584 8.581

5.628 5.634 5.614 5.612 5.615 5.608

4.343 4.347 4.357 4.354 4.355 4.352

209.522 210.095 209.991 209.678 209.978 209.482

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Fig. 10. Unit cell parameter ac (solid symbols) and the relative mass contribution pc (open symbols) of the cubic austenite phase plotted versus the temperature of thermal treatment. Solid curves are guides for the eye.

(111) and (200) peaks may indicate low degree of order in the as spun ribbons is further supported by the fact that in the case of the AWQ873 and AWQ973 samples the relative mass contribution of austenite phase is less than in the as spun ribbon sample, but on contrary the (111) peak intensity is greater, the peak is easily discernible protruding from the background, in the patterns recorded for AWQ873 and AWQ973 samples than in the patterns obtained for the as melt spun sample. This behaviour may suggest that the atomic order in these ribbons enhances due to atoms reshuffling into their lower energy sites induced by thermal treatment. Considering the A573 sample it is seen that the intensities of the (111) and (200) peaks are not easily distinguished from the background intensity, which may then be primarily attributed to a very low amount of the austenite phase in this sample. On the other hand it is worth emphasising that in the case of AWQ1073 and AWQ1173 samples the intensities of both (111) and (200) peaks increase, which may be ascribed to both the increase in the amount of austenite phase in these samples and to the increase of atomic order, taking into account the close proximity of the ordering temperature when the samples are annealed at elevated temperatures. In addition no peak of other precipitation phases is detected. It should also be noted that the diffraction peaks sharpen after annealing treatment, which is particularly visible for the AWQ1073 and AWQ1173 ribbons. This may suggest crystallite size increase upon heat treatment. 3.4. Microstructure Fig. 11 presents SEM bird view images of cross sections of as melt spun (a), A573 (b), AWQ1173 (c) ribbons and Bright Field (BF) TEM image of the cross section of A573 ribbon (D). Fig. 12 displays SEM images taken at the free surface of the as melt spun (a), A573 (b), AWQ973 (c), AWQ1173 (d) Ni48Mn39.5Sn9.5Al3 ribbons. From Fig. 11(a) it is seen that the as melt spun ribbons show typical, crystalline heterogenous microstructure with a layer of small equiaxed grains at the wheel side and columnar grains running across ribbons' thickness. It is also observed according to Fig. 11(b) and (d) that low temperature annealing does no seem to have much influence over the ribbons microstructure. In Fig. 11(d) martensite plates are visible confined to columnar grains' boundaries. On the other hand some microstructure coarsening when compared to as melt spun ribbon (Fig. 11(a)) is noticed in Fig. 11(c), which depicts the cross section of the AWQ1173 ribbon. The effect of thermal treatment on the average grain size is better visible from Fig. 12. The

average grain size in the as melt spun ribbon is estimated at circa 1.4 mm [37] (Fig. 12(a)). As already remarked the low temperature annealing does not appear to have much impact on the average grain size of the A573 ribbon (Fig. 12(b)). The average grain size in this instance remains at the comparable level. On the other hand annealing at higher temperature followed by WQ results in a remarkable increase in grain size (Fig. 12(c) and (d)), which is ten fold in the case of AWQ1173 compared to the as melt spun ribbon. This agrees well with previously discussed XRD findings. What is more whereas the images of as melt spun and A573 ribbons show no additional signs of surface effects, the SEM micrographs of AWQ973 and AWQ1173 reveal some surface morphology changes most likely attributable to surface oxidation during the WQ process. For this reason EDS analysis was performed in order to confirm the chemical composition and homogenity of ribbons. The results of EDS analysis are given in Table 4. From this table it is seen that the as melt spun, A573 and AWQ873 ribbons retain the initial composition and remain chemically homogenous. Unlike the AWQ973 and AWQ1173 ribbon samples for which a depletion in Mn concentration is noticed. This may be linked to oxidation propensity of this element. This then leads to a decrease in the e/a ratio. Fig. 13(a) shows the TEM bright field (BF) image and the corresponding Selected Area Electron Diffraction Pattern (SADP) taken at RT from the A573 sample. It may be seen from this figure that the microstructure of the A573 sample shows typical for martensite plate like microstructure. It is also observed that the microstructure is predominantly martensitic. The SADP of this phase can well be indexed in the 10 M martensite structure [49] with [010] zone axis. This indicates that annealing at 573 K for 635 min led to the change of martensite structure since in the as melt spun ribbon sample the martensite phase showed the 4O structure [37]. In order to confirm the effect of annealing at this temperature on the structure of the martensite phase in situ TEM studies were undertaken. Fig. 14 shows a series of BF images and corresponding SADPs taken from the as melt spun Ni48Mn39.5Sn9.5Al3 sample while heating it to 613 K and then cooling it back down to RT by switching off the heat supply. Initially at RT the as melt spun ribbon shows twin variant martensitic microstructure (Fig. 14(a)) with SADP (inset) well indexed in the 4O martensite structure [37]. Upon heat treatment the microstructure begins to evolve into austenite phase, which is first evidenced by thinning and gradual disappearance of twin variants (Fig. 14(b) and (c)). It may appear striking at this point that at 318 K the martensitic plates are still visible in Fig. 14(c) since the TpM/A, as discussed above, is determined at 300 K. This may be associated with the fact that during the heating up procedure the system was not allowed sufficient time to equilibrate and therefore the actual temperature on the sample is likely to had been lower at the moment of BF and SADP recording. At 338 K (Fig. 14(d)) and then at 613 K (Fig. 14(e)) the sample has completely transformed into austenite phase. Then the heating was switched off and the sample was cooled to RT. Since the forward transition in the as melt spun ribbon was observed at 291 K after cooling to RT the sample was removed from the microscope and hung for 30 s over a container with liquid nitrogen to allow for sufficient cooling. Subsequently BF and SADP were taken (Fig. 14(f)). It is evident from this figure that the sample has once more completely retransformed into martensite phase, but of the 10 M structure. In the next stage the same sample was heated to 573 K and kept at this temperature for 75 min. The time period chosen for this procedure was motivated by the saturation like effect shown in Fig. 4, from which it appears that extended annealing time above 75 min brings no significant change to the characteristic transformation temperatures. The results of this experiment are shown in Fig. 15. It is seen that after annealing the martensite structure

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39

Fig. 11. SEM micrographs of the cross section of Ni48Mn39.5Sn9.5Al3 ribbons: as melt spun (a), A573 (b), AWQ1173 (c); Bright Field (BF) TEM image of the cross section of the A573 (d).

Fig. 12. SEM micrographs of the free surface of the Ni48Mn39.5Sn9.5Al3 ribbons: as melt spun (a), A573 K (b), AWQ973 (c), AWQ1173 (d).

retains its 10 M feature (Fig. 15(a), (c)). From Fig. 15(b) and (d) it may be observed that annealing at 573 K for 75 min has no effect in terms of the grain size enlargement. Additionally no secondary or precipitation phases have been observed, unlike as in the case of high temperature annealed and WQ samples. Fig. 16 shows Scanning Transmission Electron Microscopy (STEM) e HAADF micrographs of AWQ973 (a) and AWQ1173 samples (b). In the former sample besides the matrix phase secondary phase appears, which according to EDX microanalysis is enriched in Mn. Whereas the

latter shows lamelar like morphology. According to EDX microanalysis the lamelar like features differ in composition and they are denoted as A and B. A has composition of Ni55.1Mn29.4Sn13.6Al1.8 (e/ a ¼ 8.172), while B of Ni65.1Mn31.4Sn2.2Al1.2 (e/a ¼ 8.838). This is consistent with the above DSC and XRD data. Fig. 17 shows BFs and corresponding SADPs taken from the AWQ973 sample from the martensite phase (a) and from the Mn-rich secondary phase (b). The martensite phase can be well indexed in the 4O structure with [140] zone axis. Whereas the Mn-rich phase can be indexed in the

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Table 4 Chemical composition of the Ni48Mn39.5Sn9.5Al3 ribbons determined for the as melt spun, A573, AWQ873, AWQ973, AWQ1073 and AWQ1173 ribbons. Ribbon

Composition EDS (at. %)

As spun A573 AWQ873 AWQ973 AWQ1073 AWQ1173

Ni

Mn

Sn

Al

e/a

48.9 48.7 48.1 47.5 49.0 48.2

38.7 39.0 39.1 37.2 35.6 35.3

9.2 9.1 9.7 10.2 10.4 10.8

3.2 3.2 3.1 5.1 5.0 5.7

8.065 8.059 8.032 7.916 7.958 7.894

14M martensite structure with [111] zone axis. It appears that annealing for 1 h at 973 K followed by WQ unlike annealing at 573 K, does not result in the change of the martensite structure in the whole volume of the sample, nevertheless it leads to the decomposition due to the local composition fluctuations in microareas what may result in the change of martensite structure locally. The extent of decomposition and its impact on the martensite structure is more evident in the AWQ1173 sample. BF images and corresponding SADPs taken from different phases present in this sample are shown in Fig. 18. Fig. 18(a) and (b) show BF image and corresponding SADP, respectively, taken from the lamellar like phase marked as A in Fig. 16(b). Its SADP can be well indexed in the L21 Heusler structure with [101] zone axis. On the other hand the SADP of the second lamelar like phase marked as B in Fig 16(b) and shown along with the corresponding BF image in Fig. 18(d) and (c), respectively, can be indexed in the cubic Mn0.22Ni0.78 phase with [0e13] zone axis. Fig. 18(e) and (f) depict BF image and the corresponding SADP, respectively, taken from the martensite phase contained within the cubic Mn0.22Ni0.78 phase and clearly visible in Fig. 16(b). Its SADP can be well indexed in the L10 structure with [112] zone axis. This is a clear signature of the decomposition processes taking place at this temperature range and is further in accordance with the above presented DSC and XRD studies. The L10 martensite is then likely to constitute the secondary transformable phase suggested by the DSC peak shape reversibility in the AWQ1173 DSC curve. 4. Discussion Overall it was observed that both low temperature post annealing at 573 K and high temperature quenching (873, 973, 1073, 1173 K) yield an increase in the critical temperatures, however the mechanism by which this comes about is somewhat different between these two thermal treatment protocols.

In the case of low temperature annealing at 573 K for 635 min the TM is found to shift from 291 K in the as melt spun sample to 303 K in the low temperature annealed sample and TCA increases from 299 K to 307 K. Also as is visible from the M (H) loops the EB is reduced after 635 min of ageing at 573 K. Therefore it is suggested that in this instance of low temperature annealing the influence of heat treatment on the critical behaviour of the studied ribbons is mitigated mainly by stress relief and possibly to some extent by the change in the degree of atomic order. This is evidenced by SEM and TEM analysis, which unequivocally demonstrated that the annealing at 573 K exerts no influence on the grain size, nevertheless it leads to a change in the martensite structure. The structure changes from 4O, which is present in the as melt spun ribbon, into 10M upon ageing. What is more no secondary precipitation phases appear during this treatment and there is no change in the composition of the ribbons, e/a remains unchanged. This is brought forth in connection to the apparent linear dependence of Ms on e/a in certain NieMn based Heusler alloys [50]. Similar results of low temperature (200e350 K) annealing followed by WQ on the TM and magnetic properties of bulk Ni44Mn45Sn11 and Ni45Co5Mn36.6In13.4 alloys have been reported by Wang et al. [31] and Chen et al. [48], respectively. At this stage one should bare in mind the opposite effects of annealing treatment on the properties of bulk and ribbon NieMn based alloys originating in the different fabrication routes. Whereas in bulk, homogenized alloys heat treatment can lead to atom site disordering in rapidly quenched ribbons it can result in the enhancement of atomic ordering [31]. Both groups found TM decreasing with increasing annealing temperature. Since Chen and his co-authors observed no change to the grain size nor chemical composition upon annealing this effect was primarily attributed to stress and structural relaxation, as in the current letter. In the present scenario these factors may even be more sensitive to thermal treatment considering non-equilibrium conditions of rapid quenching, by which ribbons were manufactured, and which entail high concentration of quenched in vacancies conductive to the ordering process, and an increased lattice strain due to rapid cooling [51]. Stress relaxation, related to stress formed during the quenching process, may decrease the elastic strain energy and henceforth lower the driving force required for MT, which may then necessitate lower undercooling and effectively lead to an increase in TM [52,53]. On the other hand both stress and structural relaxation may alter atom site ordering, the MneMn interatomic distance and the lattice symmetry, the latter is clearly observed here. This then may have an impact on the MneMn exchange coupling and the relative position of the Fermi surface and the Brillouin zone. It is established for a number of Heusler alloys that MT takes place when the Fermi surfaces comes into contact with the Brillouin zone

Fig. 13. TEM BF micrograph and the corresponding SADP taken at RT from Ni48Mn39.5Sn9.5Al3 ribbon subjected to annealing treatment at 573 K for 635 min.

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Fig. 14. TEM BF micrographs and corresponding SADPs (insets) from the as melt spun Ni48Mn39.5Sn9.5Al3 ribbon taken at RT (a), and in-situ while heating the sample to 613 K at 313 K (b), 318 K (c), 338 K (d), 613 K (e) and again at RT (f) after cooling the sample back to RT.

boundary. Therefore it may be assumed that micro-strain relaxation and more equilibrium atom site ordering imposed by ageing at 573 K reduces the strain energy and enhances the degree of the mutual overlap between the Fermi surface and the Brillouin zone boundary promoting in consequence MT and yielding a decrease in EB due to more Mn atoms resuming their second nearest neighbours, lower energy sites. Also owing to greater degree of order FM interactions within the austenite phase are strengthened, what shifts TCA towards higher values. Structural relaxation may also account for the martensite structure change after annealing. It is established that the different types of martensite differ in terms of the twinning stress [54]. And therefore decreased strain energy in a relaxed structure may then alternate towards a more stable martensite configuration. Also the influence on the type of martensite structure by a potential, vacancies assisted, increase in the degree of order should not be altogether discarded. Ito et al. revealed for instance that quenching from the B2 and L21 phases gives rise below the TM to L10 and 6M martensite structures, respectively, in Ni44.2Co5.4Mn37.2In13.2 MSMA alloy [55]. Furthermore Liu et al. studied the effect of ageing at 573 and 673 K on MT temperature and magnetic properties of Ni43Mn46Sn11xSix alloys [56] and found that ageing at 573 K increases the reverse MT, the Curie temperature and magnetization. Also the intensity of the

(111) superlattice reflection peak of the Heusler phase increased and it was suggested that ageing at this temperature enhanced the chemical order of the sample. On the other hand ageing at 673 K led to decomposition of the Heusler phase and the appearance of the second Mn3Ni2Si-type phase with minute amount of Sn. The Sn content then effectively increased in the matrix phase reducing TM. It was further confirmed that solution treatment of this sample from 1073 K restores both MT and TC. Similar results of microstructure evolution under ageing at 670 and 770 K of NieMn-Ga alloys have been discussed by Santamarta et al. [57], what brings the discussion to the second thermal treatment protocol employing annealing for 1 h at 873, 973, 1073 and 1173 K followed by WQ. Unlike in the case of low temperature ageing the critical transformation temperatures increase in the samples annealed at elevated temperatures followed by WQ is chiefly attributable to the combined effect of grain size enlargement, chemical composition change due to losses in Mn concentration as well as decomposition of the Heusler phase and the change in the degree of atomic order. It is clearly shown by TEM that annealing at elevated temperatures leads to the development of a secondary, eutectoid type phase with the composition of Ni65.1Mn31.4Sn2.2Al1.2. This effect is also supported by DSC and magnetic measurements from which it is observed that DSC peaks for samples annealed at 1073 and 1173 K

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Fig. 15. TEM BF micrographs and corresponding SADPs (insets) taken from Ni48Mn39.5Sn9.5Al3 ribbon after the first heating cycle (Fig. 15) at RT (a) and in situ while heating the sample to 573 K at this temperature at 0 min (b), again at 573 K after in situ annealing for 75 min (d) and at RT (c) after cooling the annealed sample back to RT.

Fig. 16. STEM-HAADF micrographs taken at RT from AWQ973 (a) and AWQ1173 (b) samples.

broaden and the TC shifts from around 300 Ke315 K for the AWQ1173 sample, indicating the appearance of a secondary phase. This is consistent with the initial report by Krenke et al. [58], who showed that with lowering Sn content the martensitic transformation temperature increases. In this instance matrix phase decomposition results in the portion of the original Ni48Mn39.5Sn9.5Al3 phase transforming to a composition with lower Sn content and hence leads to the appearance of part of the DSC peak at the higher temperature range. At the same time the matrix phase is enriched in Sn, which lowers TM in that fraction of the sample. This is further consistent with the XRD results, which show that the amount of austenite phase at RT increases with increasing temperature of thermal treatment. Also TEM analysis at RT showed untransformed regions of the AWQ1173 sample. This results are further in agreement with the reports by Yuhasz et al. [33,34] who studied decomposition and the occurrence of the eutectioid like

microstructure in Ni50Mn50xSnx (x ¼ 11, 13,15) alloys subjected to 4 weeks annealing at 1223 K followed by additional annealing at 773 K for various lengths of time. It was concluded then in conjunction with previous studies that at 773 K the solid solution Heusler phase decomposes for x ¼ 10e19 and the range of solid solution extends only from x ¼ 20 to 25, however with no MT occurring within this latter composition range. In this letter it is shown beyond no doubt that the Ni48Mn39.5Sn9.5Al3 melt spun ribbon is metastable over the temperature range 1073e1173 K and decomposes into Ni65.1Mn31.4Sn2.2Al1.2 and Ni55.1Mn29.4Sn13.6Al1.8. Composition modification may then also lead to the change in the martensite structure as recently studied by Zheng et al. in NieMneSn alloys, who observed the appearance of L10 martensite with e/a well above 8.30. And in this instance e/a as calculated based on the results of EDX microanalysis from NieMn phase is equal to ~8.84. It should also be emphasised that decomposition

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43

Fig. 17. TEM BF micrographs and corresponding SADPs taken at RT from AWQ973 sample from the matrix phase (a), (b) and from the secondary phase (c), (d).

does not seem to occur in the entire volume of the sample but rather locally, which may then explain the increase in the amount of austenite at RT in AWQ1073 and AWQ1173 samples as shown in Table 3. Also as demonstrated by XRD studies the martensite structure in AWQ samples remains in the 4O configuration. The fact that decomposition does not seem to take place in the whole volume of the material may be associated with limited annealing time. Yuhasz et al. [33,34] studied decomposition after weeks scale of annealing, whereas in the present study the annealing time was restricted to 1 h at 873, 973, 1073 and 1173 K temperatures. The reduction in EB may furthermore suggest that solution treatment from 873, 973, 1073 and 1173 K temperatures promotes the increase of atomic order. Which is further supported by the appearance of the (111) and (200) superlattice reflection peaks visible for AWQ samples. This is in part due to the increase in the amount of the parent phase at RT when compared to the A573 sample and secondly due temperature induced enhancement in the degree of ordering. One should also note that with increasing solution treatment temperature Mn loses become evident, which may also contribute to EB reduction, considering the current view on the formation of AFM clusters due to AFM MneMn exchange coupling when excess Mn atoms locate in Sn sites. According to EDS analysis Mn content depletion, due to volatilization at high temperature, is more pronounced in AWQ973, AWQ1073 and AWQ1173 samples, which coincides with the lower shift of M (H) loops toward the negative field axis recorded for these samples. The composition change also alters the e/a ratio, which tends to decrease with increasing thermal treatment temperature. Following from the above discussion related to the mutual relationship between TM and e/a ratio in NieMn based Heusler systems it would be expected that this decrease would affect adversely the TM temperature. But on contrary the opposite is observed despite decreasing e/a. This then may suggest that modification of the

electronic structure due to enhanced ordering and strain relaxation exert stronger influence on the TM than the compositional change. This phenomenon is worth pointing out since it is noted that for the NieMneSn systems MT only occurs within the narrow electron concentration range between 8.0 and 8.2 [59]. However this occurs at the expense of EB weakening. On the other hand the effect of grain size on the characteristic martensite start Ms temperature in Ni50Mn37Sn13 polycrystalline alloy ribbons has been more recently studied by Quintana-Nedelcos et al. [60] who observed an inverse relation between Ms and the reduction in the average grain size, which is linked to the elastic strain energy as discussed above. The authors concluded that decreasing of the austenite grain size coupled with the increase in the crystal defects density, limits the size of martensite variants and therefore results in the austenite stabilization. In this work it is shown by explicit SEM and TEM analysis that the grain size increases substantially upon annealing at 973 K and reaches a fifty fold increase in the case of the AWQ1173 ribbon sample. Therefore this effect may account in part for the observed structural transformation temperatures increase, in consistence with the preceeding discussion. 5. Summary and conclusions The influence of low temperature annealing and high temperature annealing followed by water quenching on martensitic transformation, microstructure, magnetic and exchange bias properties has been studied in Ni48Mn39.5Sn9.5Al3 melt spun ribbons. It is found that both thermal treatment protocols lead to an increase in the martensite transformation temperature and exchange bias reduction. In the case of low temperature annealing this effect is associated predominantly with stress and structural relaxations, whereas in the case of high temperature annealing followed by water quenching it is linked to several factors including

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Fig. 18. TEM BF micrographs and corresponding SADPs taken at RT from AWQ1173 sample from the lamellar like phases shown in Fig. 16(b): from the phase at the left hand side of Fig. 16(b) image (a), (b), from the phase seen at the centre of Fig. 16(b) image (c), (d) and from the stripe like phase visible at the bottom right of Fig. 16(b) image (e), (f).

grain size enlargement and composition change due to thermal phase instability. In total it is unequivocally demonstrated that low temperature annealing at 573 K for 75 min is the preferable method for fine tuning TM temperature and magnetic properties of NieMneSneAl based Heusler alloy ribbons affecting neither ribbons microstructure nor phase composition, and as such it may aid the design and development of future materials for multifunctional applications. Acknowledgements The authors would like to gratefully acknowledge prof. Maria Bałanda from the Henryk Niewodniczanski Institute of Nuclear Physics Polish Academy of Sciences for her most kind, prompt and fluent assistance with thermo-magnetization measurements, and also many hearty thanks go to prof. Marek Faryna from the Institute of Metallurgy and Materials Science Polish Academy of Sciences for his expert advice on EDS microanalysis.

This work has been carried out with the financial support of the European Union within the frame of the European Social Fund (Project no. POKL. 04.01.00-00-004/10). References [1] H.E. Karaca, I. Karaman, B. Basaran, D.C. Lagoudas, Y.I. Chumlyakov, H.J. Maier, On the stress-assisted magnetic-field-induced phase transformation in Ni2MnGa ferromagnetic shape memory alloys, Acta Mater 55 (2007) 4253e4269. [2] V.D. Buchelnikov, V.V. Sokolovskiy, Magnetocaloric effect in Ni-Mn-X (X ¼ Ga, In, Sn, Sb) Heusler alloys, Phys. Met. Metallogr. 112 (2011) 633e665. [3] Y.B. Yang, X.B. Ma, X.G. Chen, J.Z. Wei, R. Wu, J.Z. Han, H.L. Du, C.S. Wang, S.Q. Liu, Y.C. Yang, Y. Zhang, J.B. Yang, Structure and exchange bias of Ni50Mn37Sn13, J. Appl. Phys. 111 (2012), 07A916-1-3. [4] J. Pons, E. Cesari, C. Segui, F. Masdeu, R. Santamarta, Ferromagnetic shape memory alloys: alternatives to Ni-Mn-Ga, Mater. Sci. Eng. A 481e482 (2008) 57e65. [5] Y. Sutou, Y. Imano, N. Koeda, T. Omori, R. Kainuma, K. Ishida, K. Oikawa, Magnetic and martensitic transformations of NiMnX (X ¼ In, Sn, Sb) ferromagnetic shape memory alloys, Appl. Phys. Lett. 85 (2004) 4358e4360. [6] T. Krenke, M. Acet, E.F. Wassermann, X. Moya, L. Manosa, A. Planes,

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