Keywords: Heusler alloys, melt-spinning, magnetic shape memory.
Abstract. The most extensively studied Heusler alloys are those based on the Ni-Mn-Ga system. However, to overcome the high cost of Gallium and the usually low martensitic transformation temperature, the search for Ga-free alloys has been recently attempted, particularly, by introducing In, Sn or Sb. In this work, Mn50Ni40In10, Mn50Ni34In16, Ni50Mn36-xIn14+x (x = 0, 0.5, 1, 1.5) and Ni50Mn37Sn13 ribbons has been obtained by melt spinning. We outline their structural and thermomagnetic behavior. Columnar grains and preferential orientation has been obtained. The martensitic, Tm, and the Curie, TC, temperatures of the ribbons are lower than those of the bulk samples with similar compositions. This effect is probably due to the ribbons small and constrained grains. For it, a large under-cooling is necessary for the martensitic transformation. The decrease of TC in the ribbons could be associated with the increased degree of quenched-in short-range disorder around defects.
Introduction
Ferromagnetic shape memory alloys exhibit ferromagnetic and shape memory effect simultaneously. The ferromagnetic shape memory effect can be controlled by temperature and stress, as well as by magnetic field. Their potential functional properties are: magnetic superelasticity [1], large inverse magnetocaloric effect [2] and large magneto-resistance change [3]. These properties make them of noteworthy interest for developing new thermal or magnetically driven actuators, sensors and magnetic coolant for magnetic refrigeration [4].
Shape memory alloys exhibit a first-order martensite phase transition. By lowering the temperature a cubic high-temperature parent austenite phase transforms into a tetragonal, orthorhombic or monoclinic structurally modulated martensite ordered by domains. The transformation temperatures of shape memory alloys strongly depend on the composition and their values spread in a very wide range [5]. The mobility of the martensitic domains allows inducing large macroscopic deformations of the sample by applying an external stress. This deformation does not cost much energy since the crystal structure remains unmodified; only the domain walls move [6].
Among the alloys that exhibit magnetic shape memory effect, the most extensively studied are the shape memory Heusler alloys. A century ago, Heusler first reported that Mn-Cu bronze can be alloyed with other elements: Sn, Al, As, Sb or B [7]. Nowadays, Heusler alloys are defined as magnetic ternary intermetallic systems with L21 or B2 crystal structure. Their generic formula is X2YZ. Here, X is usually a transition metal 3d (Fe, Co, Ni, Cu, Zn), 4d (Ru, Rh, Pd, Ag, Cd) or 5d (Ir, Pt, Au). The position of Y is usually occupied by 3d (Ti, V, Cr, Mn), 4d (Y, Zr, Nb), 5d (Hf, Ta) or by lanthanides (Gd, Tb, Dy, Ho, Er, Tm, Yb, Lu) or actinides (U). The Z is a group-B element:
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III-B (Al, Ga, In, Tl), IV-B (Si, Ge, Sn, Pb) or V-B (As, Sb, Bi). In Figs. 1 and 2 are shown the L21 and an example of ordered martensite [8] structures, respectively.
Fig. 1 The Heusler L21 structure consisting of four interprenetrating face-centered cubic sublattices A, B, C and D. B and D shows positions of X atoms, A shows positions of y atoms and C shows
positions of Z atoms.
Fig. 2 Unit cell corresponding to the (_
55 ) stacking sequence model of the 10-layered martensite [8].
144 Solid Phase Transformations II
Usually, these alloys were bulk polycrystals obtained by arc or induction melting followed by a high temperature annealing, or single crystals grown by Czochralski method [9-11]. Actually, rapid quenching by melt-spinning offers three potential advantages for the fabrication of practical devices: the ribbon shape, the avoiding (or reduction) of the annealing to reach a homogeneous single phase and the fact that is easy to obtain highly textured polycrystalline ribbons. Melt-spinning conditions: the geometry of the device, the ejection gas pressure, the melting temperature, the clearance between a nozzle head and the roller surface and wheel rotation speed influence the structure of the resulting ribbon as well as their dimensions. These parameters affect the heat transfer and the crystallization kinetics. Melt spinning has been performed in several systems exhibiting shape memory: Fe-Cr-Mn-Si-Tb-B [12], Ni-Mn-Ga [13], Ni-Mn-In-Co [14] and FePd [15]. One of the first times reported in a Heusler type alloy was in 1999 [16]. In this work, one step martensitic transformation was detected accompanied by the formation of modulated martensite crystal structures, which differ from those, observed for single crystals of the same compositions.
The most extensively studied Heusler alloys have those of the Ni-Mn-Ga system. However, to overcome some of the problems related with practical applications (such as the high cost of Gallium and the usually low martensitic transformation temperature) the search for Ga-free alloys has been recently attempted. In particular, by introducing In or Sn. Martensitic transformation in ferromagnetic Heusler Ni50Mn50-xSnx alloys with 10 ≤ x ≤ 16.5 was first reported by Sutou et al. [17]. Later, Krenke and co-workers studied phase transformations, magnetic and magnetocaloric properties of the Heusler Ni50Mn50-xSnx alloy series with 5 ≤ x ≤ 25 [18-19]. Samples with x = 13 and 15, are ferromagnetic in the martensitic state undergoing a first order martensitic-austenitic structural transition at a temperature below the respective Curie points of both phases. Brown et al. [20] and Koyama et al. [21] reported on the structural and magnetoelastic behavior of the alloy Ni50Mn36Sn14. Ni-Mn-Sn system is therefore of prospective importance as ferromagnetic shape memory alloy. In all the cases, alloys were produced as bulk polycrystalline ingots by arc melting followed by high temperature homogenization annealing.
Concerning to In compositions, since Sutou et al. [17] reported the occurrence of martensitic transformation in the ferromagnetic Heusler system Ni50Mn50-xInx, considerable attention has been dedicated to study magnetism, magnetic shape memory effect [22-24], magnetic entropy change [24-28] and magneto-transport properties [29-31] of these alloys. Nevertheless, ferromagnetism in both phases is only observed in the narrow composition range of 15 ≤ x ≤ 16 [22]. The characteristic temperatures of the reversible first order structural transformation between both phases, referred as martensitic and austenitic starting and finish temperatures (i.e. MS, Mf, AS and Af, respectively), strongly vary upon small changes in the chemical composition. The crystal structure of austenite and martensite depends on composition and the transformation can be also induced by applying a magnetic field [22,24]. Additionally, a large inverse and direct magnetocaloric effect has been measured in Ni50Mn34In16 [26-28]. It has been also reported in Ref. [32] a magnetic-field induced strain by a reverse phase transformation in Ni-Mn-In-Co. Therefore, Ni-Mn-(In, Sn) Heusler alloys are of significant prospective importance for applications in both, magnetically driven actuators due to magnetic shape memory effect, and as working substances in magnetic refrigeration technology.
Our recent contribution in this field deals with the synthesis of these alloys by rapid solidification using the melt spinning technique and their basic magnetostructural characterization. Previous results of the research carried out can be found elsewhere [33-36]. Hence, we outline the microstructural and phase transition characteristics of melt spun ribbons of Heusler alloys in the Ni-Mn-(In, Sn) systems.
Experimental
As-cast pellets of nominal compositions, Mn50Ni40In10, Mn50Ni34In16, Ni50Mn36-xIn14+x (x = 0, 0.5, 1, 1.5) and Ni50Mn37Sn13 were prepared by Ar arc melting from 99.98% pure Ni, 99.98% % pure Mn, 99.999 % pure Sn and 99.999 % pure In, using Bühler MAM-1 compact arc melter. Ingots were
Solid State Phenomena Vol. 150 145
melted four times to ensure a good starting homogeneity. The samples were melt-spun in argon environment.
X-ray diffraction (XRD) experiments were performed in a Siemens D 500 (S2) diffractometer coupled with a TTK temperature chamber. The ribbons flakes (about 1 cm large) were put parallel to the plane defined by incidence and diffracted beam using a special device for low temperature patterns (see Fig. 3). Scanning was carried out in the interval 20o ≤ 2θ ≤ 80o with a step increment of 0.05o.
Fig. 3 Device designed to perform XRD patterns at low temperature.
Microstructure and elemental composition of ribbons were examined by JEOL JSM-6100 scanning electron microscope (SEM) equipped with a microanalysis system based on the Oxford Instruments INCA-energy model 200 EDX silicon detector.
Magnetization measurements were performed in the temperature interval of 4.2 – 350 K, using a Quantum Design PPMS-14T platform with a vibrating sample magnetometer (VSM) module. Zero-field cooled (ZFC), field cooled (FC) and field heated (FH) thermomagnetic curves were recorded at Hext = 50 Oe and 50 kOe, at a temperature heating or cooling rate of 2 K/min. The magnetic field was applied along the ribbon axis. Curie point (TC) was inferred from the minimum in the dM/dT vs. T curve. Hysteresis loop were performed up to 10 kOe in a Quantum Design MPMS-5T with a SQUID module.
Results and Discussion
The analysis of ribbons morphology was followed by SEM. The images corresponding to the fracture cross-section and the free-surface of ribbons are similar in all samples. As an example, Fig. 4 shows micrographs corresponding to Ni50Mn37Sn13 composition. As the micrographs demonstrate, samples are fully crystalline showing fracture surfaces of a cleavage type revealing the fast crystallization and growth kinetics of the alloy. The ribbon thickness was around 7-15 µm. Ribbons are brittle and easy to cleave in such direction. A thin layer of small equiaxed grains crystallizes in the ribbon surface in contact with the wheel. Further, an abrupt change takes place by the formation of ordered columnar grains growing perpendicularly to ribbon plane. This suggests that the rapid solidification process induces a directional crystalline grain growth. At the free-surface another thin
146 Solid Phase Transformations II
layer of small grains appears. It cuts and covers the oriented columnar structure. The granular microstructure exhibited by the free surface is shown in Fig. 4 (up).
Fig. 4 SEM micrographs of the free surface (up, bar: 20 µm) and the fractured cross section (down,
bar: 10 µm) of as-spun Ni50Mn37Sn13 ribbons.
A careful study by EDX microanalysis was carried out to estimate the average elemental chemical composition. After numerous (> 30) analyses performed on the cross section and both ribbon surfaces of different samples, we found a nearly homogeneous distribution of the chemical elements but a shift with respect to the original composition. For example, the average composition found were Mn49.5Ni40.4In10.1 and Ni50.6Mn36.3Sn13.1 whereas the original compositions were Mn50Ni40In10 and Ni50Mn37Sn13. SEM examinations in backscattering emission mode also confirmed that the chemical elements are homogeneously distributed in the alloy. It has been found, in the Ni-Mn-Co-In system, that the ribbons formed by melt spinning has an increased macroscopic chemical homogeneity if compared with bulk samples [37]. Furthermore, segregation of minor or secondary phases was not observed.
20 µm
10 µm
Solid State Phenomena Vol. 150 147
The crystal structures evolution with temperature were investigated by XRD. Figs. 5-7 correspond to Mn50Ni40In10, Mn50Ni34In16 and Ni50Mn37Sn13 compositions, respectively. The patterns here presented were obtained at 250 and 160 K. In all compositions, at 250 K the crystalline phase is the ordered cubic L21. At 160 K, martensite was found in In10 and Sn13 ribbons, but In16 maintains the L21 structure.
20 40 60 802 Theta / deg
Intens
ity / a.u.
1 0 11
0 0 14
1 2 7 1
2 -12
2 1 11
3 0 6
1 1 5
2 0 -1
160 K
14M
2 0 0
2 2 0
4 2 0
4 0 0
4 2 2
250 K
L21
1 1 -7
2 0 2
3 2 7
1 1 1
3 1 1
0 4 0
1 1 -5
Fig. 5 XRD diffraction pattern for as-quenched Mn50Ni40In10 ribbons measured at 250 K (up) and
160 K (down). The crystal structures are L21 austenite and 14M monoclinic martensite, respectively.
The texture associated to columnar grains was found. In the case of In10 sample, at 250 K the columnar grains perpendicular to the ribbon surface corresponds to the 400 reflection of the L21 phase whereas at 160 K corresponds to the 040 reflection of the 14M monoclinic martensite. In the case of Sn13 sample, at 250 K the columnar grains perpendicular to the ribbon surface corresponds to the 200 reflection of the L21 phase whereas at 150 K corresponds to the 022 reflection of the 7M orthorhombic martensite. The sample In16 does not present martensite transformation, at least in the temperature interval analyzed. It is known that transformation depends on the valence electron concentration e/a [38]. Table 1 shows the lattice parameters corresponding to XRD patterns. The values are similar to that detected in previous works [33-35].
The XRD spectra corresponding to the Ni50Mn40-xIn10+x system shows a L21 structure at room temperature. The main peak corresponds to the 220 reflection and a relative increase/decrease of the 422/400 reflection is found as increasing In content. In bulk alloys with close composition the structure at room temperature corresponds to the 10M martensite [18]. Furthermore, mixed modulated structures with texture were found at room temperature in melt-spun Ni50Mn36In14 ribbons [37]. Probably, the modification of the production conditions or small changes in the composition favors the development of different structures. It is known, in other family of alloys, that depending on melt-spinning quenching rates it is possible to develop materials with the desired structure [39]. Thus, materials with different transformation temperatures and magnetoelastic
148 Solid Phase Transformations II
behavior can be produced by controlling production conditions. Heat treatments can modify the structure but sometimes remains the same [37]. In textured ribbons, the columnar shape (1-dimension) makes difficult structural changes that must be more probable in bulk materials (3-dimensions).
20 40 60 802 Theta / deg
Intens
ity / a.u.
2 2 0
1 1 1
3 1 1
4 0 0
4 2 0
1 1 1 4 2 2
2 0 0
160 K
L2
2 0 0
2 2 0
3 1 1
4 2 04 0 0 4 2 2
250 K
L21
1
2 2 2
2 2 2
Fig. 6 XRD diffraction pattern for as-quenched Mn50Ni34In16 ribbons measured at 250 K (up) and
160 K (down). The crystal structure is L21.
Solid State Phenomena Vol. 150 149
20 40 60 802 Theta / deg
Intens
ity / a.u.
2 2 0
0 2 2
0 0 2
2 2 2 4
0 0
0 4 0
0 0 4
2 0 2
4 2 2
2 4 22 0 0
160 K
7M
2 0 0
2 2 0
3 1 1
4 2 0
4 0 0
4 2 2
250 K
L21
Fig. 7 XRD diffraction pattern for as-quenched Ni50Mn37Sn13 ribbons measured at 250 K (up) and
160 K (down). The crystal structures are L21 austenite and 7M orthorhombic martensite, respectively.
Preferential orientation is not eliminated by heat treatment. The lattice parameters of the L21
phase are compared with those of alloys with in reported by Krenke et al. [22], see Fig. 8.There is a linear behaviour and an extension of the stability of the L21 phase at room temperature. Both the structural, Tm, and the magnetic, TC, temperatures of the ribbons are lower than those of the bulk samples. Due to the ribbons small and constrained grains, a large under-cooling is necessary for the martensitic transformation. The decreased TC of the ribbon is probably associated with the increased degree of quenched-in short-range disorder around defects, as proposed by Chernenko et al. [40].
Table 1 XRD lattice parameters corresponding to ribbons Mn50Ni40In10, Mn50Ni34In16 and Ni50Mn37Sn13.
T / composition Mn50Ni40In10 Mn50Ni34In16 Ni50Mn37Sn13
250 K L21
a = 0.6004(4) nm L21
a = 0.5989(3) nm L21
a = 0.5979(4) nm
160 K
Monoclinic 14M a = 0.4265(6) nm b = 0. 5772(5) nm c = 2.979(4) nm
θ = 93.32º
L21
a = 0.5981(4) nm
Orthorhombic 7M
a = 0.6145(6) nm b = 0. 6064(4) nm c = 0.5612(7) nm
150 Solid Phase Transformations II
12 16 20 24 28at%. In
5.98
6.00
6.02
6.04
6.06
6.08
Lattice
param
eter (A)
o
Fig. 8 Lattice parameter of the L21 ordered cubic structure as a function of In at.%. Cros (this work), rhombs [22].
Mn50Ni40In10 and Ni50Mn37Sn13 ribbons are good candidates to perform the thermomagnetic
analysis. For the In10 composition, the ZFC, FC and FH thermomagnetic curves recorded at 50 Oe and 50 kOe are shown in Fig. 9. In the ZFC-curve at 50 Oe, see Fig. 9(a), the magnetization keeps low values and a small positive slope with the increase in temperature and has a sudden raise at around 220 K probably due to the structural martensite-austenite phase transformation found by XRD. A similar behavior is observed in the FH-curve at 50 Oe. The FC-curve simultaneously confirms its reversible and thermally hysteretic character. The characteristic temperatures of the respective martensite-austenite phase transformation are MS = 213 K, Mf = 173 K, AS = 222 K, and Af = 243 K. TC for the high-temperature austenite phase is ~310 K. The inset shows the low-field heating and cooling dM/dT(T) curves proving that the phase transition occurs in a broad temperature interval. The thermal hysteresis can be estimated from the maximum in dM/dT between heating and cooling in the transition region. Its value is as large as ∆T= 38 K.
In Mn50Ni40In10 ribbons, the heating and cooling thermal dependence of saturation magnetization measured at a high field value of 50 kOe, shown in Fig. 9b, exhibits two well distinct ferromagnetic phases and illustrates the reversible and abrupt change in MS as well as the field-dependence of the martensitic transformation. A significant decrease in the characteristic temperatures of the structural transformation is observed (estimated values: ∆MS = -20 K, ∆Mf = -40 K, ∆AS = -22 K, and ∆Af = -22 K). Martensite shows a lower saturation magnetization than austenite. Furthermore, complementary hysteresis loops measured at 270 K and 150 K confirm the ferromagnetic ordering exhibited by both phases [35]. Saturation magnetization is higher for austenite phase as it was previously observed in thermomagnetic curves of Fig. 9. M(H) curve for austenite approaches to saturation at a lower field value being magnetically softer than martensite, as a result of its higher symmetry cubic structure and hence lower magnetocrystalline anisotropy. Coercive field, HC, for austenite is negligible within the uncertainty of the measurement (± 2 Oe), while HC for martensite is 93 Oe [35].
Solid State Phenomena Vol. 150 151
0 50 100 150 200 250 300 3500
10
20
30
40
50
60
0 50 100 150 200 250 300 3500
20
40
60
80
100
40 80 12016020024028020
40
60
80
150 200 250 300 350
311 K
234 K
dM/dT (arb. u
nits)
Temperature (K)
196 K
Mf= 173 K
MS= 213 K
AS= 222 K
Af= 243 K
38 K
(a) H = 50 Oe
FH
FC
Mag
netiz
ation (emu/g)
Temperature (K)
(b) H = 50 kOe
Mag
netiz
ation (emu/g)
Temperature (K)
o
o
o
o
MS
Mf
Af
AS
Mag
netization (em
u/g)
Temperature (K)
Fig. 9 Temperature dependence of the magnetization measured at H = 50 Oe (up) and H = 50 kOe (down). The arrows indicate heating or cooling regime. The upper inset shows the corresponding
dM/dT(T) curve at 50 Oe.
Similar analysis was performed in the Ni50Mn37Sn13 ribbons [33,34]. Here we show the results corresponding to ribbons as quenched and as thermally treated at 800ºC during 2 hours. The annealing process was followed by a water quenching. After thermal treatment the chemical composition was shifted to Ni50.3Mn35.3Sn14.4. As an example of thermomagnetic dependences, Figs. 10a and 10b show the M(T) curves recorded at 50 Oe and 50 kOe, respectively, in ZFC, FC and FH mode of as quenched ribbons.
152 Solid Phase Transformations II
Fig. 10 Temperature dependence of magnetization with temperature measured at H = 50 Oe (up)
and H = 50 kOe (down), for as quenched Ni50Mn37Sn13 ribbons. In the ZFC-curve, M(T) stays in zero with the increasing temperature up to 100 K, afterwards a
quasi-linear monotonic increase starts. At 236 K an abrupt and sharp increase in the magnetization confirms the occurrence of the martensite-austenite transformation detected by X-ray diffraction. The same trends were found in the thermally treated ribbon. It is worthy of mention the magnetically softest behavior of austenite in annealed samples that reaches a large magnetization value at 50 Oe (one order of magnitude higher than the measured value in as-spun samples). The transition from ferromagnetic ordering to paramagnetic state for this phase occurs at 313 K in both samples. The structural transformation, in the as quenched ribbons, shows starting and finishing temperatures of MS = 218 K, Mf = 207 K, AS = 224 K, and Af = 232 K, while a thermal hysteresis of 15 K between cooling and heating curves was measured. The structural transformation, in the thermally treated ribbons, shows starting and finishing temperatures of MS = 226 K, Mf = 218 K, AS = 237 K, and Af = 244 K, while a thermal hysteresis of 21 K between cooling and heating curves was measured. In agreement with the change in composition these values are shifted with respect to
Solid State Phenomena Vol. 150 153
as-spun samples being close to those reported for the bulk Ni50Mn36Sn14 alloy [20,21]. In the thermally treated ribbons, M(T) at the FH-curve monotonously decreases until the transformation begin to start, and coincides with the magnetization value for the FC-curve in the 10 – 215 K range. At 50 kOe, see Fig. 10b, heating and cooling MS(T) shows two well distinct ferromagnetic phases illustrating a reversible and abrupt change in MS as well as the field-dependence of the martensitic transformation. A decrease in the characteristic temperatures of the structural transformation can be observed (estimated values: ∆MS = -11 K, ∆Mf = -8 K, ∆AS = -7 K, and ∆Af = -7 K).
Fig. 11 Hysteresis loops for Ni50.3Mn35.3Sn14.4 ribbons (thermally treated) measured at 150 and 270
K with the field-applied parallel to the rolling direction.
Hysteresis loops, of thermally treated ribbons, measured at 270 K and 150 K (Fig. 11), underline the ferromagnetic ordering present in both austenite and martensite. The same behavior was found in the as quenched ribbons [33]. When the magnetic field is applied parallel to the rolling direction, M(H) curve for austenite shows high susceptibility at low fields and fast approach to saturation. This magnetically soft behavior is a consequence of the low anisotropy of austenite in agreement with its cubic symmetry. Domain walls displacement would be the dominant magnetization process in this phase. Furthermore, the loop shape suggests that the easy magnetization direction lies along the rolling direction and therefore perpendicular to the major axis of columnar grains. A different situation is found for martensite, since the influence of rotation process can be observed from low fields up to 5 kOe when M(H) approaches its saturation value, indicating that anisotropy axis for this phase tilts from the rolling direction. The inset of Fig. 11 zooms into the low magnetic field region. Coercive field, HC, for austenite is ~ 28 Oe, while HC for martensite is ~ 42 Oe. These values differ from results obtained in as quenched ribbons, HC, for austenite is ~ 15 Oe, while HC for martensite is ~ 95 Oe. Hysteresis loops for both phases do not saturate up to near 10 kOe when the magnetic field is applied perpendicular to the ribbon plane. The loop is anhysteretic for austenite, but HC is ~ 88 Oe for martensite and harder magnetic properties are found, confirming that magnetization easy axes are in the ribbon plane for both phases.
The effect of magnetic field on reverse transformation was studied by measuring magnetization isotherms, under increasing and decreasing the field, in the transformation temperature range. Before recording each isotherm, the sample was cooled down to 200 K, to assure that the measurement starting in the martensitic phase, and then heated up to the measuring temperature in zero-field. After stabilizing the temperature, the M(H) curve was measured and then the measuring
154 Solid Phase Transformations II
protocol is repeated again. M(H,T) curves corresponding to as quenched ribbons are shown in Fig. 12. All of them have the signature of metamagnetism owing to the effect of the magnetic field on the transformation. Consequently, a hysteresis between field-up and field-down curves is observed. The behavior is emphasized as the temperature approaches to 240 K where the inflection point of the low-field M(T) curve was obtained [34]. At this temperature a well defined discontinuity in the first derivative of the field-up M(H) curve is observed at H = 17 kOe. Over this field value a progressive magnetic-field-induced transformation from martensite to austenite is detected. As temperature approaches to ´pure´ austenitic or martensitic phase these effects vanish.
Fig. 12 Magnetization isotherms for Ni50.3Mn35.3Sn14.4 ribbons (thermally treated) measured in the
temperature interval where the reverse martensitic transformation occurs.
Conclusions
The melt-spinning technique is proposed as an effective method to prepare single-phase Ni-Mn-(In, Sn) ribbons with single ordered structure. The composition remains close to that of the starting master alloy. From morphological analysis, columnar grains and preferential orientation have been obtained. In as-quenched conditions, austenite shows the cubic L21 crystal structure, transforming into a 7M orthorhombic (Ni50Mn37Sn13) or 14M monoclinic martensite (Mn50Ni40In10).
The thermomagnetic characterization performed in melt-spun Mn50Ni40In10 and Ni50Mn37Sn13 alloy ribbons allow us to confirm the existence of martensitic-austenitic transformation with both phases exhibiting ferromagnetic ordering. Martensite shows a lower saturation magnetization than austenite. Magnetization easy axis lies along the rolling direction for austenitic phase while the anisotropy direction for martensitic one tilts out of it. Furthermore, it has been also stated that both alloys present field-induced reverse martensite transformation. A large inverse magnetocaloric effect, comparable to that of bulk materials, originated from the abrupt change in magnetization around structural transition, together with advantages of the single stepproduction by melt spinning, points to the fact that the NiMnSn and MnNiIn Heusler ribbons could be of potential interest not only as ferromagnetic shape memory materials but as magnetocaloric ones.
Acknowledgements
FICYT is acknowledged by J.L. Sánchez Llamazares (COF07-013). This work has been supported by the Spanish MEC under projects MAT2006-13925-C02-01, MAT2006-13925-C02-02 and NAN-
Solid State Phenomena Vol. 150 155
2004-09203-C04-C03. Authors are grateful to Prof. V. Shavrov, Dr. V. Koledov and Dr. V.V. Khovaylo for helpful discussion. Serveis Tècnics de Recerca of UB is also acknowledged.
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