Size dependence of martensite transformation temperature in nanostructured Ni–Mn–Sn ferromagnetic shape memory alloy thin films

Surface & Coatings Technology 204 (2010) 3773–3782

Contents lists available at ScienceDirect

Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t

Size dependence of martensite transformation temperature in nanostructured Ni–Mn–Sn ferromagnetic shape memory alloy thin films Ritu Vishnoi, Davinder Kaur ⁎ Functional Nanomaterials Research Laboratory, Department of Physics and Center of Nanotechnology, Indian Institute of Technology Roorkee, Roorkee-247667, India

a r t i c l e

i n f o

Article history: Received 13 February 2010 Accepted in revised form 22 April 2010 Available online 28 April 2010 Keywords: DC magnetron sputtering FSMA Martensitic transformation Ni–Mn–Sn thin films

a b s t r a c t In the present study, we report the influence of grain size on structural and phase transformation behaviour of nanostructured Ni–Mn–Sn ferromagnetic shape memory alloy thin films synthesized by dc magnetron sputtering. With increase in substrate temperature, the structural phase changes from austenite with L21 cubic crystal structure to martensite with monoclinic structure. In addition, field-induced martensite– austenite transformation is observed in magnetization studies using superconducting quantum interference device magnetometer. The martensitic transformation behaviour of these films depends critically on the microstructure and dimensional constraint. Both, the martensite start temperature (Ms) and austenite finish temperature (Af) of these nanostructured films decreases with decreasing grain size. The excess free volume associated with grain boundaries has been observed to increase with decrease in grain size which in turn leads to an increase in the number of grain boundaries. It has been proposed that the grain boundaries impose constraints on the growth of the martensite and confine the transformed volume fraction in nanocrystalline structure. A martensite phase nucleated within a grain will be stopped at the grain boundaries acting as obstacles for martensite growth. The investigations revealed that below a critical grain size of 10.8 nm, the austenite phase is observed to be more stable than the martensite phase which leads to the complete suppression of martensitic transformation in these films. © 2010 Elsevier B.V. All rights reserved.

1. Introduction Ferromagnetic shape memory alloys (FSMA) are of considerable interest because of their exceptional magnetoelastic properties and large magnetic field-induced strain due to the rearrangement of twin variants in the martensite phase [1,2]. In FSMA, the shape memory effect can not only be controlled by changing the temperature but also by varying the magnetic field. This property makes FSMA of scientific interest for developing new thermal or magnetically driven actuators [3]. The most extensively studied FSMA is those of the Ni–Mn–Ga system due to its large magnetic field-assisted superelasticity [4], large magnetic shape memory effect, low temperature ferromagnetic martensitic phase, large magnetic field-induced strain (MFIS) [5], giant magnetocaloric effects due to large entropy change across the martensitic transition [6,7], high response frequency and premartensitic phase transition [8]. Heusler alloy Ni2MnGa is one of the most promising candidates for magnetically actuated devices due to its low hysteresis, unique magnetoelastic properties [9] and large MFIS that is an order of magnitude higher than that found for magnetostrictive materials [1,10]. Murray et al. reported 6% MFIS at room temperature in five-layer modulated tetragonal martensite phase of Ni–Mn–Ga

⁎ Corresponding author. E-mail address: [email protected] (D. Kaur). 0257-8972/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2010.04.053

[11] whereas Sozinov et al. achieved very large MFIS of 9.5% in orthorhombic 7 M martensitic phase of this alloy [1]. Although, these Ni–Mn–Ga based FSMA showed reasonably good magnetoelastic properties, their application in devices was limited due to some practical problems, such as the high cost of gallium and low martensitic transformation temperature (TM). The potential applications of Ga based FSMA are also limited due to their high brittleness. Due to these limitations with Ni–Mn–Ga based FSMA, the search for Ga free FSMA has recently been started. Sutou et al. investigated the structural properties of Ni50Mn50 − xSnx (10 ≤ x ≤ 16.5) type alloys and observed a martensitic transformation from cubic L21-type crystal structure to orthorhombic four-layered martensite structure [12]. Khovaylo et al. also reported martensitic transition in Ni–Mn–Sn alloy [13]. The shape memory effect of Ni–Mn–X [X: Sb, Sn, In] FSMA has been well investigated in their bulk form [12–16]. But the knowledge of such phenomenon in Ni–Mn–X thin films is sparse. This martensitic transition is of immense importance especially in FSMA thin films as being potential candidates for microelectromechanical systems (MEMS) applications. Nanocrystalline materials are the materials which have nanometer-sized crystallographically ordered regions or grains. The properties of nanocrystalline materials deviate from those of single crystals or coarse grained polycrystals with the same average chemical composition. This deviation is due to the reduced size and/or dimensionality of the nanometer-sized crystallites and numerous

3774

R. Vishnoi, D. Kaur / Surface & Coatings Technology 204 (2010) 3773–3782

interfaces between adjacent crystallites [17]. Earlier it has been reported in nanocrystalline materials of Ni-Ti [18–20] and Fe based alloys [21,22] that the martensite transformation temperature is dependent not only on the chemical composition but also on the crystal grain size in the sample. It has been concluded that there is a critical size at which the martensite transformation is suppressed fully on cooling. The reported value of critical grain size for Ni-Ti and Ni-TiCu nanostructured shape memory alloys is 17 nm [20] and 16 nm [23] respectively. Kajiwara et al. studied the martensitic transformation behaviour in fine particles of Fe based alloys and reported critical grain size of 20 nm [24]. H. Rumpf et al. studied the effect of annealing temperature in NiMnGa FSMA films and found that the shape memory effect is not observed until a critical grain size is obtained that only takes place for annealing temperatures in the range of 500 °C [25]. To the best of our knowledge there is no report of grain size effect on martensite transformation behaviour in Ni–Mn–Sn based FSMA thin films. Hence a systematic study is required to understand the martensitic phase transformation behaviour of thin films of Ni–Mn– Sn for application of these films in emerging micro-devices such as magnetically driven MEMS which require a high quality of FSMA thin films. The present study explored the effect of substrate temperature on Ni–Mn–Sn FSMA thin films grown by DC magnetron sputtering on Si (100) substrate. The magnetron sputtering has important specific advantages such as low levels of impurities, easy control of deposition rate and production of thin films of various morphology and crystallographic structures. Earlier, we have reported the deposition of NiTi based shape memory alloy thin films via magnetron sputtering [26,27]. The main aim of the present study is to investigate the effect of grain size on structural and phase transformation behaviour of Ni– Mn–Sn FSMA thin films. The investigation revealed that the martensitic transformation behaviour of these films depend critically on its microstructure and dimensional constraint and below a critical grain size of 10.8 nm, the austenite phase is observed to be more stable than the martensite phase which leads to the complete suppression of martensitic transformation. 2. Experimental details The Ni–Mn–Sn FSMA thin films were deposited on (100) silicon (Si) substrate at various substrate temperatures (TS) ranging from 450 to 650 °C by dc magnetron sputtering system [20]. The films deposited at TS of 450, 500, 550, 600 and 650 °C are represented as samples A1, A2, A3, A4 and A5 respectively. The substrates were initially cleaned thoroughly in an ultrasonic bath with a mixture of distilled water and trichloroethylene in 4:1 ratio, washed with boiled acetone and etched in 5% HF solution to remove the surface native oxide layer prior to load into the chamber. Ni50Mn35Sn15 target of 1 in. diameter and 3 mm thickness was used. Before every sputtering run, the targets have been pre-sputtered for 15 min in order to ascertain the same state of the targets in every run. Substrate holder was rotated at 20 rpm in a horizontal plane to achieve a uniform film composition. The target to substrate distance was fixed at approximately 5 cm. Before sputtering deposition, the chamber was evacuated to a base pressure of the order of 10− 7 Torr and then backfilled with argon to the desired pressure. The chamber pressure was measured using a combination vacuum gauge (Pfeiffer vacuum) and maintained at a constant value of 20 mTorr. The deposition was carried out for 45 min at a fixed sputtering power of 100 W. No post-annealing was performed after deposition. The thickness of the films was ∼ 1 μm, as determined by surface profilometer and cross sectional field emission scanning electron microscope (FESEM). The orientation and crystallinity of the films were studied using a ´ Bruker AXS D-8 advanced diffractometer of Cu Kα (1.54 A˚) radiations in θ–2θ geometry at a scan speed of 1°/min. To obtain a profile fitting with good signal, polycrystalline silicon powder was used for

instrumental correction. The coherently diffracting domain size (dXRD) was calculated from the integral width of the diffraction lines using the Scherrer's equation [28], after background subtraction and correction of instrumental broadening. The surface topography and microstructure were studied using FESEM (FEI Quanta 200F) and atomic force microscopy (AFM). The temperature dependence of magnetization M–T of the films in a low external magnetic field (H = 100 Oe) was measured in the temperature range 10 K ≤ T ≤ 310 K using a superconducting quantum interference device (SQUID) magnetometer (MPMS, Quantum Design) using a scan length of 4 cm with 24 data points in each scan. To further characterize the magnetic properties, the magnetic field dependence of the magnetization M–H of the samples was measured in fields up to 1 T. 3. Results and discussions 3.1. Structural properties Fig. 1 shows the X-ray diffraction (XRD) pattern of Ni–Mn–Sn thin films deposited at various TS. All the films exhibit nearly single phase with no detectable secondary or impurity phases. It is observed that the films deposited at TS below 400 °C are amorphous. The films deposited at TS ≤ 600 °C are in L21 austenite phase with prominent (220) reflection, while the films deposited at TS of 650 °C are observed to exhibit martensite phase. In general, the peaks in XRD pattern corresponding to L21 structure should have two types of superlattice reflections in addition to fundamental reflections. One type of superlattice reflections appears, when h, k, and l are all odd (h, k, and l are Miller indices). Another type of superlattice reflection appears, when h, k, and l are all even. For the fundamental reflections, h, k, l are all even and satisfy the condition h + k + l = 4n [29]. For the films deposited at TS ≤ 600 °C, both the superlattice peaks (111) and (200) along with the fundamental peak (220) are clearly observed suggesting the existence of L21 structure in the films. However in case of film deposited at TS = 650 °C, sharp peak of (1 2 7) reflection at 44.8° and other less intense peaks of (1 1 5̄), (1 0 11), (2 0 5̄) and (2 1 11) reflections are observed which correspond to monoclinic martensitic structure. An analysis of the XRD pattern reveals that the lattice parameters of the films, deposited at TS of 450, 500, 550 and 600 °C are 5.925, 5.964, 5.978 and 5.972 respectively. These values are close to the lattice parameter of bulk Ni–Mn–Sn (a ∼ 6.009 ) [14]. Similarly the lattice parameters of the film deposited at 650 °C which

Fig. 1. XRD spectra of the Ni–Mn–Sn films deposited at substrate temperatures ranging from 450 to 650 °C.

R. Vishnoi, D. Kaur / Surface & Coatings Technology 204 (2010) 3773–3782

´ ´ ´ is in martensite phase are a = 4.296 A˚, b = 5.573 A˚ and c = 29.756 A˚ ´˚ and are in close agreement with their bulk values (a = 4.333 A, ´ ´ b = 5.570 A˚ and c = 29.971 A˚) [14]. The summary of these lattice parameters and crystal structure are given in Table 1. The XRD pattern shows that with increase in TS, the (220) reflection becomes sharp up to 600 °C. Increase in TS from 450 to 600 ºC increases grain size as a result of faster grain growth which results in reduction of grain boundaries that favours the formation of closed packed structure which lead to the development of strong (220) plane texture. The reduction in grain boundaries with increase in TS is due to the energy enhancement of atoms absorbed on the substrate which now have sufficient energy to overcome the grain boundary barrier and consequently grain size increases. When thermal energy is given to the film, adatom diffusion takes place, which leads to the growth of favourably oriented grains. This suggests that the number of grains per unit volume decreases, their size increases, and the grain boundary area per unit volume decreases. For the film deposited at TS of 600 °C, a state of higher thermodynamic stability has been reached. Further increase in TS (film A5) results in dominance of martensite phase which has been ascribed to the considerable increase in grain size that decreases the additional energy barrier originating from the grain boundaries. In order to determine the density of grain boundaries, the excess free volume associated with the grain boundaries is calculated using the following expression [30]: 2

ΔVF =

2

ðL + d =2Þ −L L2

where L is the crystallite size and d is the mean width of the grain boundaries. In most prior calculations of the excess free volume, the width of the grain boundary has been assumed to be a constant (d ≈ 1 nm), independent of the grain size [31]. Inset of Fig. 1 shows that the grain boundary free volume decreases with increase in crystallite size. The mean crystallographic particle size of the films was calculated from the XRD data, using the Scherrer's formula [28]. The average value of calculated crystallite size was 4.6, 10.8, 20.3, 21 and 23.2 nm for the films A1, A2, A3, A4 and A5 respectively which reflects an enhancement in grain size with increase in deposition temperature. Fig. 2(a)–(e) shows the FESEM micrographs of Ni–Mn–Sn thin films A1 to A5 respectively. These images clearly resolve the change of microstructure of these films with change in deposition temperature. It is clear from these images that grains are pyramid type for the films A1 to A4 which diffuse to spherical shape for the film A5. It further confirms the phase change form austenite to martensite phase above a substrate temperature of 600 °C. The grain size increases with increase in TS and their values are reported in Table 1. The film surfaces have no evident defaults such as impurity, holes and cracks as can be seen from Fig. 2(a)–(e). The surface morphology of these films was also analyzed using AFM and is shown in Fig. 3(a)–(e). The values of grain size and surface roughness calculated from AFM are given in Table 1. The grain size increases with increase in TS, which is in agreement with the XRD and

3775

FESEM results. The root-mean-square roughness (Rrms) and average roughness (Ravg) of the surface was obtained from AFM scans over substrate areas of 2 μm × 2 μm, three times at a different spot for each sample. Rrms and Ravg are defined from the following relationships [32]: P 1 N ∑ jz − zj; N i=1 i " #1 = 2 P2 1 N = ∑ jz − z j N i=1 i

Ravg = Rrms

where N is the number of surface height data and z̅ is the mean-height distance. AFM micrographs reveal that the Ravg of the film increases with phase transformation from austenite to martensite phase. The film A5 with martensite phase shows Ravg of 42.8 nm and Rrms of 54.3 nm respectively, while the films A1–A4, which exhibit austenite phase, show Ravg ≤ 25 nm. It was observed that the grain sizes, determined by AFM and FESEM, are quite similar. But, the overall particle size shown by FESEM and AFM is much bigger than that calculated by XRD, which is ascribed to the fact that FESEM and AFM shows agglomeration of the particles whereas XRD gives an average mean crystallite size. The XRD and FESEM/AFM data can be reconciled by the fact that smaller primary particles have a large surface free energy and would, therefore, tend to agglomerate faster and grow into larger grains. The composition of the films, determined by energy dispersive Xray analysis (EDAX), were found to be Ni50.3Mn35.2Sn14.5, Ni49.7Mn35.1Sn15.2, Ni50.5Mn34.8Sn14.7, Ni49.6Mn34.9Sn15.5 and Ni50.2Mn34.9Sn14.9 for samples A1, A2, A3, A4 and A5 respectively. Fig. 4(a) shows a typical EDAX spectrum for the Ni–Mn–Sn film deposited at 550 °C. All the peaks can be identified as arising from the film (Ni, Mn and Sn). Analysis of the spectra after background subtraction and deconvolution of peaks confirm the composition of the Ni–Mn–Sn film to be close to the nominal composition of the target. Fig. 4(b) reveals good composition congruency of the films with that of target. The dotted lines show the Ni, Mn and Sn concentration of the target used for the preparation of films. 3.2. Magnetic properties The magnetic and structural transition temperatures of Ni–Mn–Sn thin films of different grain sizes were determined from the thermomagnetic measurements at a low magnetic field strength of H = 100 Oe on heating and cooling in the temperature range between 5 K and 310 K using SQUID magnetometer (Fig. 5). The corresponding transformation temperatures are reported in Tables 2 and 3. Before measurements, the samples were prepared in the zero field-cooled state by cooling them from 310 to 5 K in the absence of magnetic field. Then the external field was applied and the readings were recorded during heating the samples (ZFC). Then the samples were further cooled but now in the presence of field i.e. field-cooled (FC) and measurements were done. Finally, the samples were heated in the presence of field i.e. field-heated denoted as (FH) and measurements

Table 1 Crystal structure, lattice parameter, grain size and average roughness for Ni–Mn–Sn thin films deposited at substrate temperatures ranging from 450 to 650 °C. A: austenite phase; M: martensite phase. Sample name

Substrate temperature (TS) °C

Phase

A1 A2 A3 A4 A5

450 500 550 600 650

A A A A M

Lattice parameter

Grain size (nm)

a (Å)

b (Å)

c (Å)

XRD along (220) peak

FESEM

AFM

5.925 5.964 5.978 5.972 4.296

5.925 5.964 5.978 5.972 5.573

5.925 5.964 5.978 5.972 29.756

4.6 10.8 20.3 21.0 23.2

52.9 76.7 85.8 107.8 144.1

56.7 82.6 88.5 117.6 160.7

Average roughness AFM (nm) 15.6 16.0 24.5 25.0 42.8

3776

R. Vishnoi, D. Kaur / Surface & Coatings Technology 204 (2010) 3773–3782

Fig. 2. FESEM images of Ni–Mn–Sn films deposited at (a) 450, (b) 500, (c) 550, (d) 600 and (e) 650 °C.

were taken. ZFC measurements are not shown here for convenience. The area of the substrates used for the measurements was 0.5 × 0.5 cm2. The magnetization data for all the films was corrected to account for the diamagnetic contribution of the Si (100) substrate using equation: Mfilm ðHÞ = Mtotal ðH Þ−χsubstrate ·H where χsubstrate is the susceptibility of the substrate, Mtotal is the magnetization of (film + substrate) and H is the applied magnetic

field which is parallel to the film surface. The measured value of susceptibility for Si (100) substrate comes out to be − 5.736 × 10− 8. A change in the magnetic moment has been observed at structural as well as magnetic transition points. Curie temperature (TC) was estimated using differentiation of magnetization data. The TC temperature is defined as the temperature at which the slope of magnetization versus temperature curve is the largest. In film A1 (Fig. 5(a)), with grain size of 4.6 nm, the magnetization decreases continuously with increase in temperature from 5 to 215 K above which it becomes constant. There is no signature of phase

Fig. 3. AFM images of Ni–Mn–Sn films deposited at (a) 450, (b) 500, (c) 550, (d) 600 and (e) 650 °C, respectively.

R. Vishnoi, D. Kaur / Surface & Coatings Technology 204 (2010) 3773–3782

3777

Fig. 4. (a) EDX spectrum for the Ni–Mn–Sn film deposited at 550 °C. (b) Temperature dependence of compositions of the Ni–Mn–Sn films.

transformation between martensite and austenite phase in this film. Hysteresis between FC and FH curves characteristic of first-order phase transformation is also not present as the FH curve retraces the FC curve, which further confirms the absence of phase transformation. The absence of phase transformation phenomenon could be due to the very small grain size (∼ 4.6 nm) in this film causing large number of grain boundary interfaces and associated excess free volume as shown in the inset of Fig. 1. Martensitic transformation is a first-order solid– solid phase transformation which involves distortion of the lattice, shuffling of atomic planes and is displacive in nature. Due to the displacive character, the transformation proceeds by small cooperative movements of the atoms. In film A1, the large number of grain boundaries, which act as potential barriers separating the individual grains, limits such type of atomic movements, impose constraints on the growth of the martensite structure and confine the transformed volume fraction in the nanocrystalline structure. The grain boundary width and consequently the height of the energy barrier corresponding to the grain boundaries are proportional to the width of the intergranular region, which increases with decrease in the grain

size. A martensite plate, if nucleated within a grain, is stopped at the grain boundaries. To propagate the transformation, the plate has to exert stress that should be sufficient enough to stimulate nucleation and growth of favourable martensite variants in the adjacent grains [33]. However, due to insufficient thermal activation energy required for martensitic nucleation, at 450 °C, phase transformation could not occur in this film. Further, as this film possesses grain size below critical value required for phase transformation, the austenite phase becomes more stable than the martensite phase even at low temperatures and martensitic transformation cannot take place. In Film A2 (Fig. 5(b)), with increase in temperature in the martensitic state, it is observed that the magnetization starts increasing but the increase in magnetization is very low which could be due to the smaller grain size of 10.8 nm. Then, a slight signature of phase transformation from martensite to austenite phase is observed at lower temperature of ∼56 K which is the austenite start temperature (As) and film gets fully converted to austenite at Af ∼ 111 K followed by the continuous decrease in magnetization. The FH curve does not retrace the FC curve but show hysteresis,

3778

R. Vishnoi, D. Kaur / Surface & Coatings Technology 204 (2010) 3773–3782

Fig. 5. FC and FH M (T) curves measured in an applied field of 100 Oe for Ni–Mn–Sn films deposited at (a) 450, (b) 500, (c) 550, (d) 600 and (e) 650 °C, respectively. Inset: heating and cooling curves where phase transition occurs.

R. Vishnoi, D. Kaur / Surface & Coatings Technology 204 (2010) 3773–3782

3779

Table 2 Values of saturation magnetization (SM), coercivity (Hc), remanent magnetization (Mr) (at 5, 75 and 200 K) and Curie temperatures for Ni–Mn–Sn thin films deposited at substrate temperatures ranging from 450 to 650 °C. Substrate temperature (TS) °C

TM C (K)

TC (K)

450 500 550 600 650

– – – – 161

190.0 210.4 230.0 248.5 –

75 K

200 K

5K

75 K

200 K

5K

75 K

200 K

1.33 2.10 5.68 6.37 1.14

1.16 1.79 5.34 6.12 –

0.34 0.95 2.79 4.14 –

320.7 300.0 158.5 96.8 112.5

120.0 76.0 67.1 36.5 43.5

30.0 20.2 16.2 14.3 27.2

0.77 1.01 3.52 2.92 0.53

0.55 0.70 2.70 1.52 0.03

0.04 0.11 0.73 0.95 0.01

Table 3 Details of transformation temperatures obtained from magnetization versus temperature curves of Ni–Mn–Sn thin films deposited at substrate temperatures ranging from 450 to 650 °C. Grain size (nm)

450 500 550 600 650

4.6 10.8 20.3 21.0 23.2

Remanent magnetization Mr (μB/f.u.)

5K

indicating the occurrence of first-order structural transition. TC of this film is ∼210 K. Thus with increasing temperature ferromagnetic martensite has first transformed into the ferromagnetic austenite and then into the paramagnetic austenite. FH curve lies below FC curve which indicates that hysteresis is due to the structural change and not due to magnetic changes. The thermal hysteresis between FC and FH curves is observed to be 27 K. Fig. 5(c) shows the thermomagnetic curves of the film A3 having grain size of 20.3 nm which clearly indicates the phase transformation from martensite to austenite phase and vice versa during subsequent heating and cooling cycles. With the increase in temperature from 5 K to 265 K, initially the magnetization increases slowly followed by the sudden increase in magnetization across the martensite–austenite phase transformation during which detwinning takes place as temperature increases to As. Twinned martensite phase gets completely converted to single cubic austenite phase at Af. The inset shows the phase transition region clearly along with the transformation temperatures. The behaviour of magnetization at low temperature is due to the decrease in magnetic anisotropy with increase in temperature while the sudden change in spontaneous magnetization across the martensite–austenite phase transition is due to the change in exchange interaction because of different lattice parameters in martensite and austenite phases [34]. With further increase in temperature, the magnetization decreases which is due to the ferromagnetic to paramagnetic transformation of austenite phase. Thus, the initial rise and the subsequent fall in magnetization with increasing temperature in film A3 is indication of the consequent martensitic and magnetic (TC ∼ 210 K) phase transformations. The similar behaviour is observed for the FC curve. Below Ms, the magnetization decreases with decreasing temperature which indicates that the martensitic state does not sustain long range ferromagnetism at temperatures below Ms. The start and finish temperatures of the martensite phase transformation are Ms = 92.4 K and Mf = 61 K, while those for austenite are As = 82.5 K and Af = 113 K. The thermal hysteresis of transformation is 20 K. Fig. 5(d) displays the magnetic behaviour of film A4 having bigger grain size of 21 nm. The FC and FH behaviour is similar to that observed for the film A3 except that the transformation temperatures Ms, Mf, As and Af (shown in inset of Fig. 5(d)) are shifted towards

Substrate temperature (TS) °C

Coercivity Hc (Oe)

Saturation magnetization SM (μB/f.u.)

Transition temperatures (K) Ms

Mf

As

Af

– 84.0 92.4 121.0 299.0

– 52.3 61.1 73.8 288.0

– 56.2 82.5 95.7 291.3

– 111.3 112.9 140.7 302.3

Hysteresis width (ΔK) – 27 20 19 3

higher temperatures with slight decrease in hysteresis width of phase transformation. M–T curve of film A5 is shown in Fig. 5(e). Film A5 is in martensite phase at room temperature as clear from the XRD, AFM and FESEM data. The M–T curve of this film shows that the magnetization decreases with increase in temperature from 5 to 200 K and exhibit a magnetic transition at 161 K which is the curie temperature of the martensite phase (TM C ). There is no long range magnetic ordering in the temperature range TM C ≤ T ≤ Ms. With further increase in temperature, there is a sudden increase in magnetization at ∼295 K which represents the martensitic transition (clearly shown in the inset of Fig. 5(e)) occurring in a narrow temperature range. The positions of maxima in the FC and FH curves correspond to Ms and Af respectively. Below Mf, FC and FH measurement curves retrace. The hysteresis width is significantly small (3 K) in this film. The TC cannot be observed in this film in the measured temperature range showing the TC should be above room temperature. This is advantageous for the application of FSMA films in MEMS devices. The magnetization of films A1, A2 and A5, is observed to be smaller than that of the films A3 and A4 which shows that the full magnetic ordering is not attained in these films. It is also clear from the curves in Fig. 5 that the transition temperatures Ms, Mf, As and Af of all the films increases with increase in TS. This change in transition temperatures with TS can possibly be due to the following three factors (i) change in the film composition due to the different temperatures of substrate during deposition (ii) precipitate formation and (iii) change in the grain size with TS. In our case, the effect of different composition of the films is disregarded because all the films are of nearly same composition as evident by EDAX measurement. Further, no precipitate formation occurs as evident from XRD, AFM and FESEM data so this effect is also excluded. Therefore the only possibility for the increase in Ms, Mf, As and Af with TS could be due to the formation of bigger grains with increase in TS. Grain boundaries provide transformation energy barrier for martensitic transformation. As grain size increases, grain boundaries and the width of intergranular region decreases which in turn decreases the energy barrier which was responsible to restrict the martensitic transformation. Also, the increase in grain size leads to reduction in strain energy and thus less driving force is required for transformation. This leads to rise of As and Af in the reverse transformation and Ms and Mf in the forward transformation. Thus, the increase in Ms, Mf, As and Af with increasing TS is ascribed to the increase in grain size. Fig. 6 clearly shows the variation of Ms, Mf, As and Af as a function of grain size. The TC also increases with increase in TS which is also due to the formation of bigger grains in the films. The ferromagnetic nature of the films A1, A2, A3, A4 and A5 was investigated by the isothermal magnetization versus field M–H data taken at different temperatures below TC. Each measurement was performed by cooling the sample from 310 K down to the required temperature of interest in zero field and then varying the field from 0 to 1 T. Prior to each isothermal M–H measurement, the sample was heated above TC (whenever possible) in order to bring it in demagnetized state. After cooling to the required temperature,

3780

R. Vishnoi, D. Kaur / Surface & Coatings Technology 204 (2010) 3773–3782

Fig. 6. Grain size dependence of transition temperatures Ms, Mf, As and Af.

measurements were started after some waiting times in order to ensure that thermal equilibrium was established, and this was maintained throughout succeeding measurements by using sufficiently slow (field) sweep rates. The hysteresis loops were measured with the field parallel to the film plane. In order to find out the direction of applying magnetic field the magnetization versus field measurements were performed in both the ways, i.e. by applying magnetic field parallel to the film plane as well as perpendicular to the film plane. It was observed that the magnetization gets saturated at small field for the field parallel to the film plane but do not saturate even at higher field for field perpendicular to the film plane. This suggests that the easy axis of magnetization lies in the film plane i.e. the anisotropy in case of our samples is low along the in-plane direction and the films possessed in-plane magnetization. Therefore, the rest of the measurements were performed along in-plane direction. This behaviour has also been reported previously on polycrystalline [35] and single crystalline thin films [36]. M–H plots of Ni–Mn–Sn films A1–A5, measured at 5 K (well below martensitic transition), 75 K and 200 K are shown in Fig. 7(a)–(c). In order to determine the values of coercivity (Hc) and retentivity (Mr), the M–H loops are analyzed in low field region and are shown in the inset of each figure. Table 2 lists the detailed values of the saturation magnetization (SM), Hc and Mr of the films. From Fig. 7(a) it is clear that for films A1–A4 the value of SM increases with increase in grain size, however a sharp decline in the value of SM has been observed in case of film A5, which could be due to the formation of martensite phase. Generally, it has been observed that the martensite phase exhibit lower SM than that of austenite phase [37]. The value of Hc decreases with increase in TS from 450 to 600 °C and it again increases to a value of 112.5 Oe for the film A5 deposited at TS of 650 °C. The sudden increase of Hc for the film A5 is due to its martensitic nature. The decrease of Hc with TS has also been reported earlier [38,39]. Fig. 7 (b) and (c) shows that the values of SM, Hc and Mr follow the same behaviour at 75 K and 200 K as that at 5 K. An increase in TS from 450 to 600 °C results in a reduction in Hc almost by a factor of 3 at 5 K and 2 at 200 K. The decrease in Hc is due to increase in grain size and crystallinity-degree of (220) texturing of films as TS increases from

Fig. 7. Magnetization hysteresis loops measured at 5 K (a), 75 K (b) and 200 K (c) for Ni–Mn–Sn thin films deposited at various substrate temperatures. Inset in all the three figures shows the enlarge view of the low field region.

R. Vishnoi, D. Kaur / Surface & Coatings Technology 204 (2010) 3773–3782

450 to 600 °C, whereas increase in Hc of film for TS N 600 °C is due to formation of martensite phase at room temperature. As the Hc of martensite phase is higher than austenite phase, the Hc increases for TS N 600 °C. As expected, the values of Hc and Mr of the martensite phase of these films at low temperature (5 K) are higher than that of austenite phase of films at high temperature (200 K). The Hc is of great importance as it specifies the operating magnetic field range of active devices [40]. Further, low Hc of films is especially important for application in magnetic resistive devices. Thus, it is desirable for the grown films to possess low Hc [41]. Since the Hc decreases with increasing TS and the film grown at 600 °C shows minimum Hc therefore it is preferred to grow film at 600 °C for the use of Ni50Mn35Sn15 FSMA films in devices for faster response. In Fig. 7(c), the film A5 is paramagnetic at 200 K which is expected as its TM C is 161 K and the film A1 is still showing very small ferromagnetic hysteresis even though its TC is ∼190 K. It is because of broad magnetic transition in this film as evident from the curve (a) in Fig. 5. The hysteresis loop of film A2 is linear up to about 100 Oe, after which it acquires a curvature due to the presence of ferromagnetic short range correlations, 200 K being close to TC (210 K). The hysteresis loops show nearly rectangular behaviour for films in austenite phase, deposited at high TS of 550 and 600 °C, as shown in Fig. 7(a)–(c). Fig. 8 shows the TS dependence of Hc and SM at 200 K. The measured SM at low TS of 450 and 500 °C was much smaller than the theoretically predicted value of 4.3 μB/f.u. [13,42,43]. This might indicate that at low temperatures, the full magnetic order is not attained in Ni–Mn–Sn films. The reduced SM indicates a non-magnetic intergranular volume which suppresses any intergranular exchange interaction. The SM of film A4 is 4.14 μB/f.u, which agrees well with the bulk value of 4.3 μB/f.u. The finite Hc observed in our films is the result of domain wall pinning at grain boundaries. The pinning of the domain wall becomes effective in the martensite phase as indicated by the higher Hc values of the films at 5 K in the martensite phase (Table 2) as compared to those at 200 K in the austenite phase. It is well known that Hc is associated to the crystalline anisotropy and also the grain size of the films [39]. The different Hc observed in the Ni– Mn–Sn films is attributed to the micro-structural differences such as the grain size, crystalline quality and roughness of the Ni–Mn–Sn films due to variation in TS. 4. Conclusion In summary, the Ni–Mn–Sn FSMA thin films have been grown successfully on Si (100) substrate by DC magnetron sputtering. We

Fig. 8. Effect of substrate temperature (TS) on saturation magnetization (SM) and coercivity (Hc) measured at 200 K.

3781

believe that the martensitic transformation behaviour depends critically on the microstructure and dimensional constraint of the films. Grain size is found to have a considerable effect on phase transformation behaviour of these films. XRD pattern of the films reveals austenite phase with preferential (220) reflection for the films with grain size in the range between 4.6 and 21 nm. However, only the films with grain size above 23 nm exhibit martensite phase at room temperature with monoclinic structure. Both the microstructure and grain size of the films are found to depend on the substrate temperature. It is observed that the fine nanostructure of crystalline films resulting from deposition at low temperature ≤ 450 °C does not allow the occurrence of phase transformation until a critical grain size is obtained during the coalescence process that takes place only for deposition at TS ≥ 500 °C. The magnetic behaviour of the films differs in martensite and austenite phases considerably. This difference, we believe, occurs due to the structural change as well as the formation of martensitic variants in the low temperature phase. With increase in grain size, the Hc is observed to be significantly reduced. The pronounced improvement of the magnetic properties in terms of SM and Hc with increasing grain size is probably related to the reduction in excess free volume associated with grain boundaries. Below a critical grain size of 10.8 nm, the austenite phase is observed to be more stable than the martensite phase which leads to the complete suppression of martensitic transformation in these Ni–Mn–Sn films. Acknowledgements The financial support provided by Ministry of Communications and Information Technology (MIT), India under Nanotechnology Initiative Program with Reference No. 20(11)/2007-VCND and DRDO India is highly acknowledged. References [1] A. Sozinov, A.A. Likhachev, N. Lanska, K. Ullakko, Appl. Phys. Lett. 80 (2002) 1746. [2] S.J. Murray, M. Farinelli, C. Kantner, J.K. Haung, S.M. Allen, R.C. O'Handley, J. Appl. Phys. 83 (1998) 7297. [3] M.A. Marioni, R.O.C. O'Handley, S.M. Allen, S.R. Hall, D.I. Paul, M.L. Richard, J. Feuchtwanger, B.W. Peterson, J.M. Chambers, R. Techapiesancahoroenkij, J. Magn. Magn. Mater. 35 (2005) 290. [4] V.A. Chernenko, V. L'Vov, J. Pons, E. Cesari, J. Appl. Phys. 93 (2003) 2394. [5] S.J. Murray, M.A. Marioni, A.M. Kukla, J. Robinson, R.C. O'Handley, S.M. Allen, J. Appl. Phys. 87 (2000) 5774. [6] L. Pareti, M. Solzi, F. Albertini, A. Paoluzi, Eur. Phys. J. B 32 (2003) 303. [7] M. Khan, S. Stadler, J. Craig, J. Mitchell, N. Ali, IEEE Trans. Magn. 42 (2006) 3108. [8] V.V. Kokorin, V.A. Chernenko, E. Cesari, J. Pons, C. Segui, J. Phys. Condens. Matter 8 (1996) 6457. [9] P.J. Webster, K.R.A. Ziebeck, S.L. Town, M.S. Peak, Philos. Mag. B 49 (1984) 295. [10] K. Ullakko, J.K. Huang, C. Kantner, R.C. O'Handley, V.V. Kokorin, Appl. Phys. Lett. 69 (1996) 1966. [11] S.J. Murray, M. Marioni, S.M. Allen, R.C. O'Handley, T.A. Lograsso, Appl. Phys. Lett. 77 (2000) 886. [12] Y. Sutuo, Y. Imano, N. Koeda, T. Omori, R. Kainuma, K. Ishida, K. Oikawa, Appl. Phys. Lett. 85 (2004) 4358. [13] V. Khovaylo, V. Koledov, V. Shavrov, V. Novosad, A. Korolyov, M. Ohtsuka, O. Saveel'eva, T. Takagi, Adv. Funct. Mater. 13 (2006) 474. [14] T. Krenke, M. Acet, E.F. Wassermann, X. Moya, L. Manosa, A. Planes, Phys. Rev. B 72 (2005) 014412. [15] S.B. Roy, M.K. Chattopadhyay, P. Chaddah, Phys. Rev. B 71 (2005) 174413. [16] A.K. Nayak, K.G. Suresh, A.K. Nigam, J. Phys. D:Appl. Phys. 42 (2009) 035009. [17] H. Gleiter, Acta Mater. 48 (2000) 1. [18] Y. Qin, L. Chen, Y. Zhu, L. Zhang, J. Mater. Sci. Lett. 15 (1996) 1155. [19] T. Kikuchi, K. Ogawa, S. Kajiwara, T. Matsunaga, S. Miyazaki, Y. Tomota, Philos. Mag. A 78 (1998) 467. [20] A. Kumar, D. Singh, D. Kaur, Surf. Coat. Technol. 203 (2009) 1596. [21] Q. Meng, Y. Rong, T.Y. Hsu, Phys. Rev. B 65 (2002) 174118. [22] W. Qin, Z.H. Chen, J. Alloys Compd. 322 (2001) 286. [23] A.M. Glezer, E.N. Blinova, V.A. Pozdnyakov, A.V. Shelyakov, J. Nanopart. Res. 5 (2003) 551. [24] S. Kajiwara, S. Ohno, K. Honma, Philos. Mag. A 63 (1991) 625. [25] H. Rumpf, C.M. Craciunescu, H. Modrow, Kh. Olimov, E. Quandt, M. Wuttig, J. Magn. Magn. Mater. 302 (2006) 421. [26] A. Kumar, D. Singh, R.N. Goyal, D. Kaur, Talanta 78 (2009) 964. [27] A. Kumar, D. Singh, R. Kumar, D. Kaur, J. Alloys Compd. 479 (2009) 166. [28] B.D. Cullity, Elements of X-ray Diffraction, Addison-Wesley, Reading, MA, 1970, p. 102. [29] P.J. Webster, J. Phys. Chem. Solids 32 (1971) 1221.

3782

R. Vishnoi, D. Kaur / Surface & Coatings Technology 204 (2010) 3773–3782

[30] P.P. Chattopadhyay, P.M.G. Nambissan, S.K. Pabi, I. Manna, Phys. Rev. B 63 (2001) 054107. [31] R. Banerjee, E.A. Sperling, G.B. Thompson, H.L. Fraser, S. Bose, P. Ayyub, Appl. Phys. Lett. 82 (2003) 4250; R.W. Siegel, Processing of Metals and Alloys, Vch, Weinheim, Germany, 1991, p. 583. [32] H.–.S. Zhang, J.L. Endrino, A. Anders, Appl. Surf. Sci. 255 (2008) 2551. [33] T. Waitz, V. Kazykhanov, H.P. Karnthaler, Acta Mater. 52 (2004) 137. [34] T. Krenke, E. Duman, M. Acet, E.F. Wassermann, X. Moya, L. Manosa, A. Planes, Nat. Mater. 4 (2005) 450. [35] V.A. Chernenko, M. Ohtsuka, M. Kohl, V.V. Khovailo, T. Takagi, Smart Mater. Struct. 4 (2005) 245. [36] J.W. Dong, J.Q. Xie, J. Lu, C. Adelmann, C.J. Palmstrom, J. Cui, Q. Pan, T.W. Shield, R. D. James, S. McKernan, J. Appl. Phys. 95 (2004) 2593.

[37] J.D. Santos, T. Sanchez, P. Alvarez, M.L. Sanchez, J.L. Sanchez Llamazares, B. Hernando, Ll. Escoda, J.J. Sunol, R. Varga, J. Appl. Phys. 103 (2008) 07B326. [38] C. Liu, Z.Y. Gao, X. An, M. Saunders, H. Yang, H.B. Wang, L.X. Gao, W. Cai, J. Magn. Magn. Mater. 320 (2008) 1078. [39] L.J. Singh, Z.H. Barber, Y. Miyoshi, Y. Bugoslavsky, W.R. Branford, L.F. Cohen, Appl. Phys. Lett. 84 (2004) 2367. [40] J.A. Caballero, Y.D. Park, A. Cabbibo, J.R. Childress, F. Petroff, R. Morel, J. Appl. Phys. 81 (1997) 2740. [41] J. Giapintzakis, C. Grigorescu, A. Klini, A. Manousaki, V. Zorba, J. Androulakis, Z. Viskadourakis, C. Fotakis, Appl. Phys. Lett. 80 (2002) 2716. [42] E.Z. da Silva, O. Jepsen, O.K. Andersen, Solid State Commun. 67 (1988) 13. [43] J. Kubler, A.R. Williams, C.B. Sommers, Phys. Rev. B 28 (1983) 1745.