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TbFe2)n heterostructures
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Acta Materialia 60 (2012) 1840–1845 www.elsevier.com/locate/actamat
Influence of the TbFe2 crystallization on the magnetic and magnetostrictive properties of (Fe3Ga/TbFe2)n heterostructures R. Ranchal ⇑, V. Gutie´rrez-Dı´ez, V. Gonza´lez Martı´n Dpto. Fı´sica de Materiales, Facultad Ciencias Fı´sicas, Universidad Complutense de Madrid, Ciudad Universitaria s/n, Madrid 28040, Spain Received 5 July 2011; received in revised form 22 November 2011; accepted 22 November 2011 Available online 7 February 2012
Abstract In this work we have investigated (Fe3Ga/TbFe2)n multilayers grown by sputtering at room temperature. These multilayers exhibit a large coercivity associated to the crystalline TbFe2 Laves phase. To reduce the coercivity it is necessary to control the crystallization of that material. In this work, we focus on the analysis of the properties of the TbFe2 layers. In the as-grown heterostructures we have found evidence of nanoaggregates in the TbFe2 layers. The Fe3Ga thickness and the thermal treatments have an influence on the volume of these nanoprecipitates. In the annealed samples, when increasing the Fe3Ga thickness we observe a decrease in the nanoaggregate volume and thus in the coercivity. The experimental results indicate that the crystallization of the TbFe2 depends on the Tb diffusion promoted by the thermal treatment and on the stiffness factor (Y/a) of the Fe3Ga layer. The magnetostrictive properties are also strongly influenced by the crystallization of the TbFe2. We have achieved a maximum magnetostriction constant of nearly 550 ppm with a coercive field close to 400 Oe. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Magnetic thin films; Sputtering; Magnetic properties; Crystallization
1. Introduction Crystalline TbFe2 is the material with the largest known magnetostriction constant at room temperature [1]. The study of this material is of great relevance due to its potential applications in a wide variety of magnetic devices and wireless technology. However, the large coercive field (HC) of the TbFe2 Laves phase [2] constitutes a difficulty in most of these applications. Heterostructures comprising TbFe2 and soft magnetic layers such as Ni, Fe, or FeB have already been studied [3–6]. Although the coercivity of these systems is low, the magnetostriction is also reduced. TbFe2/FeCo multilayers show excellent properties with a saturation magnetoelastic coupling coefficient (k.Y, Y being the Young’s modulus) of 27.5 MPa and a hysteresis of less than 25 Oe [7].
⇑ Corresponding author. Tel.: +34 91 394 50 12; fax: +34 91 394 45 47.
E-mail address: rociran@fis.ucm.es (R. Ranchal).
Galfenol (Fe1 xGax) exhibits a lower k than TbFe2, 440 ppm for quenched bulk samples with a 28% of Ga [8]. Galfenol is a very attractive material for magnetic sensors and actuators, not only because of its pretty high magnetostriction but also due to its low coercivity and high ductility. Several studies have focused on the structural and magnetostrictive properties of the bulk material [9– 17] and more recently, on thin films [18–23]. Heterostructures comprising Fe1 xGax and transition magnetic metal alloys such as NiFe and FeCoB have also been studied [24,25]. A review of Fe1 xGax alloys has recently been reported [26]. In this work we investigate the characteristics of asgrown and annealed (Fe3Ga/TbFe2)n heterostructures. In general, sputtered thin TbFe2 films deposited at room temperature do not exhibit high magnetostrictive properties, and post-growth thermal treatments are necessary to enhance them [27,28]. Moreover, as the TbFe2 crystalline phase is associated to a high coercivity, the control of the Laves phase crystallization is essential to engineer the
1359-6454/$36.00 Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2011.11.049
R. Ranchal et al. / Acta Materialia 60 (2012) 1840–1845
magnetic properties of these heterostructures. It has been reported that the substrate stiffness during the thermal treatment (defined by the Y/a factor, where Y is the Young’s modulus and a is the thermal expansion coefficient) is closely related to the crystallization process of sputtered thin TbFe2 films [27,28]. It is important to remark that the best results are achieved when the substrate or buffer layer has a high Y/a factor but not close to that of TbFe2. Mo is a suitable material for the TbFe2 crystallization process due to its high Y/a factor, 70 106 GPa K. On the other hand, the Tb diffusion promoted by the thermal treatment in (TbFe2/Fe3Ga)n heterostructures has also been studied [29]. This study shows that the annealing promotes the diffusion of Tb atoms from the Fe3Ga layers and/or the Fe3Ga/TbFe interfaces towards the TbFe layers. This Tb diffusion seems to be prevented in heterostructures with thin layers, and thus TbFeGa alloys might be present at the interfaces of these multilayers. Here, we report on the influence of the annealing process on the magnetic and magnetostrictive characteristics of (Fe3Ga/TbFe2)n multilayers. We have focused our investigation on the TbFe2 layers because low coercivity systems can be achieved by controlling the crystallization of the Laves phase. Our results indicate that both the Tb diffusion and the Y/a factor of the Fe3Ga layers have an effect on the magnetic and magnetostrictive properties of the studied (Fe3Ga/TbFe2)n heterostructures. Higher HC and magnetostriction values are achieved in (TbFe2/Fe3Ga)n multilayers because of the modification of the growth order. 2. Experimental techniques Samples were grown by the sputtering technique at room temperature on glass substrates. The deposition was carried out in oblique incidence with an angle between the vapor beam and the perpendicular to the sample of 25°. The Ar pressure during growth was 2 10 3 mbar and the power was 120 W for the deposition of the TbFe2 and 100 W for the Fe3Ga layers, respectively. Mo buffer and capping layers (20 nm) were deposited in all the samples. They were grown with a dc power of 90 W and at an Ar pressure of 2 10 3 mbar. We have deposited two different sets of (Fe3Ga/TbFe2)n heterostructures: (a) a set in which the Fe3Ga layers range from 12.5 to 50 nm and a constant TbFe2 thickness of 50 nm, and (b) multilayers where the Fe3Ga thickness ranged between 12.5 and 50 nm and a constant TbFe2 thickness of 12.5 nm. We have also modified the growth order and then (TbFe2(12.5 nm)/ Fe3Ga(12.5 nm))16 and (TbFe2(50 nm)/Fe3Ga(50 nm))4 heterostructures were also deposited for further analysis. Hysteresis loops at room temperature were carried out in a vibrating sample magnetometer (VSM). The temperature dependence of the magnetization has been characterized by means of SQUID magnetometry. Prior to measuring the temperature dependence of the magnetization, the sample was first cooled from room temperature to 5 K either under a saturation field of 2 kOe (field-cooled: FC) or at zero
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field (zero-field-cooled: ZFC). Then, FC and ZFC curves were recorded with an applied magnetic field of 100 Oe during the warming up. Hysteresis loops at 10 K were measured in the SQUID. Thermal treatments were carried out in Ar atmosphere during 1 h at 400 °C. After the annealing, samples were cooled down at a rate of 20 °C min 1 because fast cooling downs enhance the magnetostrictive properties of the Galfenol layers [8]. The optical cantilever method has been used to infer the magnetostrictive properties of the heterostructures [30]. 3. Results and discussion 3.1. Magnetic properties As-grown TbFe2 thin films deposited in oblique incidence on rigid buffers such as Mo consist of nanoaggregates as revealed by the presence of a cusp in the ZFC curve [28]. In this work, we have also performed FC– ZFC curves to analyze the structural properties of the (Fe3Ga/TbFe2)n heterostructures. In Fig. 1a and b we present these FC–ZFC curves for as-grown heterostructures with 12.5 or 50 nm Fe3Ga and 50 nm TbFe2 layers (set a). In these heterostructures the TbFe2 layers are always deposited on top of the Fe3Ga, and thus the latter acts as the buffer of the former. The higher the Fe3Ga thickness, the lower the influence of the Mo buffer on the TbFe2. When the Fe3Ga thickness is 12.5 nm (Fig. 1a), the ZFC curve shows a cusp at 152 K similar to that observed in single TbFe2 films directly deposited on Mo [28]. The increase of the Fe3Ga thickness up to 50 nm shifts the ZFC cusp to higher temperatures, 192 K, indicating the increase in size of the TbFe2 nanoprecipitates (Fig. 1b). Then, in the asgrown heterostructures the nanoaggregate volume in the TbFe2 layers increases with the Fe3Ga thickness. In Fig. 1c we show an example of the hysteresis loop at low temperature that was obtained for this type of heterostructure. Even when the layer thickness is as low as 12.5 nm we can observe two different magnetic behaviors in the loops: a low coercivity phase that is related to the Fe3Ga and a high coercivity material that corresponds to the TbFe2. In order to calculate the volume of the nanoprecipitates it would be necessary to know the anisotropy field (HK) and the saturation magnetization (Ms) of the TbFe2 layers [28], however these parameters are difficult to infer from the measured hysteresis loops. Then, in this work we will just qualitatively analyze the evolution of the TbFe2 nanoaggregates with the layer thickness and/or the thermal treatments. Fig. 2 shows how the annealing modifies the FC–ZFC curves of the heterostructures. The ZFC cusp of the multilayer with thin Fe3Ga layers, (Fe3Ga(12.5 nm)/ TbFe2(50 nm))7, shifts to 230 K, supporting the idea of an increase in the nanoaggregate volume upon thermal treatment (Fig. 1a). Nevertheless, the ZFC cusp of the sample with thick Fe3Ga layers, (Fe3Ga(50 nm)/TbFe2(50 nm))4, shifts to 146 K, pointing to a decrease of the nanoprecipitate
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Fig. 2. FC (d) and ZFC (s) curves for different annealed heterostructures. (a) (Fe3Ga(12.5 nm)/TbFe2(50 nm))7, and (b) (Fe3Ga(50 nm)/ TbFe2(50 nm))4.
Fig. 1. (a) FC (d) and ZFC (s) curves of the as-grown (Fe3Ga(12.5 nm)/ TbFe2(50 nm))7 heterostructure. (b) FC (d) and ZFC (s) curves of the asgrown (Fe3Ga(50 nm)/TbFe2(50 nm))4 heterostructure. (c) Hysteresis loop at 10 K of the as-grown (Fe3Ga(12.5 nm)/TbFe2(12.5 nm))16 heterostructure.
size (Fig. 2b). The negative magnetization values observed in some ZFC curves (Figs. 1a and 2a) are related to the diamagnetic contribution of the glass substrates. The results of Figs. 1 and 2 show that the Fe3Ga has a different effect on the as-grown and annealed heterostructures. In the as-grown samples we observe that the TbFe2 nanoprecipitate volume increases with the Fe3Ga thickness. However, in the annealed samples we observe the opposite behavior. The higher the Fe3Ga thickness, the lower the nanoprecipitate volume. The effect of the Fe3Ga layers when the samples are annealed can be an indication of the influence of this material on the crystallization of the
TbFe2. The increase of the Fe3Ga thickness seems to prevent the crystallization of the TbFe2. We have measured hysteresis loops at room temperature to analyze the effect of the TbFe2 microstructure on the coercivity of these two (Fe3Ga/TbFe2)n samples of set a (Fig. 3). The as-grown samples show pretty similar hysteresis loops except for the Ms (Fig. 3a). This difference can be explained taking into account that each sample has different layer thicknesses being the Ms of the Fe3Ga higher than that of TbFe2. Also, the low quality that we typically observe in our as-grown sputtered TbFe2 films is expected to decrease the Ms of those layers [27,28]. The HC of these two as-grown heterostructures is pretty similar, 125 Oe for the (Fe3Ga(50 nm)/TbFe2(50 nm))4 and 180 Oe for the (Fe3Ga(12.5 nm)/TbFe2(50 nm))7, respectively. Thus, in the as-grown samples, the different TbFe2 nanoaggregate volume seems not to have a noticeable effect on the coercivity of the samples. After the thermal treatment the HC of the (Fe3Ga (50 nm)/TbFe2(50 nm))4 multilayer slightly increases, whereas the annealed (Fe3Ga(12.5 nm)/TbFe2(50 nm))7 heterostructure shows a dramatic increase in coercivity up to 4 kOe (Fig. 3b). This extremely high HC reflects the crystallization of the TbFe2 in agreement with the increase of the TbFe2 nanoaggregate volume observed in Fig. 2a. On the other hand, the decrease of the nanoprecipitate size upon
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Fig. 4. Coercive field (HC) as a function of the Fe3Ga layer thickness (t) for as-grown (s) and annealed (d) (Fe3Ga(t)/TbFe2(12.5 nm))n heterostructures. t = 0 refers to the (TbFe2(12.5 nm)/Fe3Ga(12.5 nm))16 heterostructure.
Fig. 3. Hysteresis loops for the (d) (Fe3Ga(50 nm)/TbFe2(50 nm))4 and (h) (Fe3Ga(12.5 nm)/TbFe2(50 nm))7 heterostructures. (a) as-grown, and (b) annealed.
the thermal treatment in the (Fe3Ga(50 nm)/TbFe2 (50 nm))4 multilayer can explain the low HC attained in that sample. The increase of the Fe3Ga thickness seems to prevent the TbFe2 crystallization process. This can be due to the low Y/a stiffness factor of the Fe3Ga on the TbFe2 layers or to the Tb diffusion already observed in (TbFe2/Fe3Ga)n heterostructures [29]. In order to clarify the influence of the Fe3Ga thickness, we have studied another set of samples where the Fe3Ga thickness (t) ranges from 0 to 50 nm and the TbFe2 thickness, 12.5 nm, is fixed (set b). In Fig. 4 we present the coercivity of as-grown and annealed (Fe3Ga(t)/TbFe2 (12.5 nm))n heterostructures as a function of the parameter t. Actually, the sample with t = 0 is the (TbFe2(12.5 nm)/ Fe3Ga(12.5 nm))16 heterostructure, t = 0 being used to indicate that the first TbFe2 layer is deposited on top of Mo instead of on top of Fe3Ga. Thus, in the case of t = 0 we have modified the growth order. Regardless of the Fe3Ga thickness the coercivity is similar in all the as-grown samples, HC 190 Oe (Fig. 4). This coercivity is similar to the values observed in the multilayers previously discussed (set a). The thermal treatment has an influence on the HC of samples belonging to set b (Fig. 4). The annealed samples with a Fe3Ga thickness higher than that of TbFe2 (t > 12.5 nm) have a coercivity that is even lower than that
before the thermal treatment. The multilayer with t = 12.5 nm, i.e. (Fe3Ga(12.5 nm)/TbFe2(12.5 nm))16, does not present a noticeable modification of the HC, and the HC increases up to 400 Oe in the sample with t = 0, the (TbFe2(12.5 nm)/Fe3Ga(12.5 nm))16 multilayer. The deposition of the first TbFe2 layer on top of Mo because of the modification of the growth order doubles the HC. To confirm this result we have summarized in Table 1 the HC obtained in heterostructures with the same layer thicknesses but different growth order. In all cases, the coercivity is increased when the first TbFe2 layer is deposited on top of Mo, i.e. the HC is higher in the (TbFe2/Fe3Ga)n heterostructures (Table 1). In a previous work, we present experimental evidences of the accumulation of Tb atoms in the Fe3Ga/TbFe2 interfaces and/or Fe3Ga layers during the growth by sputtering of (TbFe2/Fe3Ga)n multilayers [29]. As-grown heterostructures have Tb–Fe layers with a Tb concentration lower than the nominal 33%. The thermal treatment can promote the diffusion of these Tb atoms to the Tb–Fe layers. If the Tb atoms diffuse upon the thermal treatment, the concentration of Tb in the Tb–Fe layers increases, being closer to the expected 33%. We also observed that the strong lattice mismatch in (TbFe2/Fe3Ga)n heterostructures with thin layers prevented the Tb diffusion [29]. The minimum in the FC curves is known as the compensation temperature (TComp) and it can be used to determine the composition of the Tb–Fe layers [29]. In this work, the presence of a TComp below 300 K reveals a Tb content lower than 28% in the Tb–Fe layers of as-grown heterostructures (Fig. 1a and b). The shift of the TComp towards higher temperatures Table 1 HC of different annealed heterostructures. Annealed heterostructure
HC (Oe)
[Fe3Ga (12.5 nm)/TbFe2(12.5 nm)]16 [TbFe2(12.5 nm)/Fe3Ga(12.5 nm)]16 [Fe3Ga (50 nm)/TbFe2(50 nm)]4 [TbFe2(50 nm)/Fe3Ga(50 nm)]4
180 400 140 250
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after the thermal treatment indicates the Tb enrichment of the Tb–Fe layers because of the Tb diffusion (Fig. 2). Taking into account Ref. [29], Tb diffusion is only expected in heterostructures comprising Fe3Ga layers with a thickness higher than that of TbFe2, so in this work diffusion is expected in samples with a Fe3Ga thickness above 12.5 nm (t > 12.5 nm). As the HC of the TbFeGa alloys is higher than that of the Fe3Ga, the reduction of these ternary alloys because of the Tb diffusion can explain the decrease of the coercivity upon the thermal treatment observed in Fig. 4 when t > 12.5 nm. Nevertheless, Tb diffusion cannot explain the increase of coercivity obtained when just the growth order is modified (samples with t = 0 and t = 12.5 nm in Fig. 4 and the samples summarized in Table 1). In those samples the layer thickness is the same and so the diffusion process is equivalent. The increase of the coercivity observed in (TbFe2/ Fe3Ga)n in comparison to (Fe3Ga/TbFe2)n multilayers needs to be explained by considering a different effect. In the (TbFe2/Fe3Ga)n multilayers, one TbFe2 layer is deposited on top of Mo that has a high Y/a stiffness factor (70 106 GPa K) and promotes an optimum crystallization of the TbFe2. In the (Fe3Ga/TbFe2)n heterostructures, all the TbFe2 layers are deposited on top of Fe3Ga that has a low Y/a factor (6 106 GPa K) and prevents the crystallization of the TbFe2. This effect is in agreement with the reduction of the TbFe2 nanoprecipitate volume observed in Fig. 2b. Then, to engineer the magnetic properties of heterostructures comprising Fe3Ga and TbFe2 is crucial to control the crystallization of the TbFe2 layers by means of the layer thickness and the growth order.
Fig. 5. (a) Magnetostriction as a function of the applied magnetic field of the (Fe3Ga(12.5 nm)/TbFe2(37.5 nm))7 multilayer. The direction of the magnetic field is parallel (d) and perpendicular (s) to the long side of the cantilever. (b) Magnetostriction constant (k) as a function of the Fe3Ga layer thickness (t) for as-grown (h) and annealed (j) (Fe3Ga(t)/ TbFe2(12.5 nm))n heterostructures. t = 0 refers to the (TbFe2(12.5 nm)/ Fe3Ga(12.5 nm))16 heterostructure.
3.2. Magnetostrictive properties Our goal is to achieve not just low coercivity but also high magnetostrictive systems. Typical magnetostriction for these samples can be seen in Fig. 5a. Fig. 5b shows the dependence of the magnetostriction constant with the Fe3Ga thickness. As previously indicated, t = 0 is used to include the (TbFe2(12.5 nm))/Fe3Ga(12.5 nm))16 heterostructure. The k of the as-grown samples is rather low and in general below 100 ppm. Our as-grown sputtered TbFe2 layers exhibit low magnetostrictive properties and thus the k of these as-grown samples seems to be mostly correlated with the Fe3Ga layers. Bulk quenched Galfenol samples can exhibit a k as high as 440 ppm for a Ga content of 28% but a k of 150 ppm has been reported in sputtered Fe81.6Ga18.4 thin films [20], a value closer to our results in as-grown multilayers. The thermal treatments clearly enhance the magnetostriction of the samples, with a maximum k of 550 ppm in the (TbFe2(12.5 nm)/Fe3Ga(12.5 nm))16 heterostructure being achieved. This enhancement can be due to several reasons. First of all, in the samples with thick Fe3Ga layers, the Tb atoms diffuse and the increase of Tb content in the Tb–Fe layers will produce an enhancement of the magnetostriction. When the Tb diffusion is prevented because the Fe3Ga
thickness is reduced, TbFeGa ternary alloys will be present. In bulk TbFeGa alloys, a large k constant of 1000 ppm has been reported when the Ga content is below 10% [31]. The existence of these ternary alloys can also explain the increase of k when the Fe3Ga thickness is reduced (Fig. 5b). However, when the growth order is modified it is necessary to consider more effects. The higher magnetostriction of the (TbFe2 (12.5 nm)/Fe3Ga(12.5 nm))16 sample seems to be related to the growth of one TbFe2 layer on top of Mo that enhances the crystallization process of that material. In the (Fe3Ga(12.5 nm)/TbFe2 (12.5 nm))16 multilayer all the TbFe2 layers are grown on top of Fe3Ga, which does not promote an optimum crystallization of the TbFe2 and lower k values are achieved. Therefore, the magnetostrictive properties of the studied multilayers depend on the crystallization process of the TbFe2. Although the TbFe2 crystallization can be prevented by means of the layer thickness combination and the growth order, the highest k value is achieved when the crystallization of the TbFe2 is enhanced and the growth order is (TbFe2/Fe3Ga)n. Further experiments need to be done to improve the characteristics of these heterostructures, but promising values of k 550 ppm and a HC of 400 Oe in the (TbFe2(12.5 nm)/Fe3Ga(12.5 nm))16, and
R. Ranchal et al. / Acta Materialia 60 (2012) 1840–1845
a k of 225 ppm and a HC of 190 Oe in the (Fe3Ga (12.5 nm)/TbFe2 (12.5 nm))16, have been achieved. 4. Conclusions We have studied the magnetic and magnetostrictive properties of (Fe3Ga/TbFe2)n heterostructures. In particular, we have focused our investigation on the crystallization process of the TbFe2 layers due to the high coercivity of the Laves phase. We have found that the magnetic and magnetostrictive properties depend on both the layer thickness and the growth order. On the one hand, the TbFe2 crystallization can be prevented by increasing the Fe3Ga thickness. On the other hand, enhanced magnetostrictive properties have been achieved in (TbFe2/Fe3Ga)n. In these latter multilayers the crystallization of the TbFe2 is improved as a consequence of the modification of the growth order. A pretty high magnetostriction constant of nearly 550 ppm has been obtained when the coercivity is 400 Oe. Acknowledgements This work has been financially supported by Universidad Complutense de Madrid (UCM) and Madrid Regional Government through the project CCG10-UCM/MAT4621. We thank “CAI Te´cnicas Fı´sicas” of UCM for SQUID characterization. We also acknowledge Prof. C. Aroca and Dr M. Maicas for fruitful discussions. References [1] [2] [3] [4]
Clark AE, Belson HS. Phys Rev B 1972;5:3642. Clark AE. Appl Phys Lett 1973;23:642. Lee HS, Cho C. IEEE Trans Magn 2008;44:4049. Quandt E, Ludwig A, Mencik J, Nold E. J Alloys Compd 1997;258:6232. [5] Shima T, Takanashi K, Fujimori H. J Magn Magn Mat 2002;239:573.
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[6] Prados C, Panagiotopoulus I, Hadjipanayis GC, Hernando A, Freijo JJ. IEEE Trans Magn 1997;33:3712. [7] Quandt E, Ludwig A. Sensors Actuat A 2000;81:275. [8] Clark AE, Wun-Fogle M, Restorf JB, Lograsso TA, Cullen JR. IEEE Trans Magn 2001;37:2678. [9] Clark AE, Restorff JB, Wun-Fogle M, Lograsso T, Schalgel D. IEEE Trans Magn 2000;36:3238. [10] Clark AE, Wun-Fogle M, Restorf JB, Dennis KW, Lograsso TA, McCallum RW. J Appl Phys 2005;97:10M316. [11] Ikeda O, Kainuma R, Ohnuma I, Fukamachi K, Ishida K. J Alloys Compd 2002;347:198. [12] Cao H, Bai F, Li J, Viehland DD, Lograsso TA, Gehring PM. J Alloys Compd 2008;465:244. [13] Xing Q, Lograsso TA. Appl Phys Lett 2008;93:182501. [14] Xing Q, Du Y, McQueeney RJ, Lograsso TA. Acta Mater 2008;56:4536. [15] Cao H, Gehring PM, Devreugd CP, Rodrı´guez-Rivera JA, Li J, Viehland D. Phys Rev Lett 2009;102:127201. [16] Laver M, Mudivarthi C, Cullen JR, Flatau AB, Chen WC, Watson SM, et al. Phys Rev Lett 2010;105:027202. [17] Ruffoni MP, Pascarelli S, Gro¨ssinger R, Sato Turtelli R, BormioNunes C, Pettifer RF. Phys Rev Lett 2008;101:147202. [18] Javed A, Morley NA, Gibbs MRJ. J Magn Magn Mater 2009;321:2877. [19] Morley NA, Javed A, Gibbs MRJ. J Appl Phys 2009;105:07A912. [20] Basantkumar RR, Hills Stadler BJ, Robinson WP, Summers EM. IEEE Trans Magn 2006;42:3102. [21] Weston JL, Butera A, Lograsso TA, Shamsuzzoha M, Zana I, Zangari G, et al. IEEE Trans Magn 2000;38:2832. [22] Butera A, Go´mez J, Weston JL, Barnard JA. J Appl Phys 2005;98:033901. [23] Dunlap RA, Deschamps NC, Mar RE, Farrell SP. J Phys: Condens Matter 2006;18:4907. [24] Lupu N, Chiriac H, Pascariu P. J Appl Phys 2008;103:07B511. [25] Lupu N, Pascariu P, Gherasim C, Chiriac H. IEEE Trans Magn 2008;44:3005. [26] Atulasimha J, Flatau AB. Smart Mater Struct 2011;20:043001. [27] Ranchal R, Prieto JL, Sa´nchez P, Aroca C. J Appl Phys 2010;107:113918. [28] Ranchal R, Lo´pez E, Prieto JL, Aroca C. Acta Mater 2011;59:2865. [29] Ranchal R, Gonza´lez-Martı´n V. J Appl Phys 2011;110:053901. [30] Ciudad D, Prieto JL, Lucas I, Aroca C, Sa´nchez P. J Appl Phys 2007;101:043907. [31] Tang YJ, Luo HL, Gao NF, Liu YY, Pan SM. Appl Phys Lett 1995;66:388.
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Influence of the TbFe2 crystallization on the magnetic and magnetostrictive properties of (Fe3Ga/TbFe2)n heterostructures R. Ranchal ⇑, V. Gutie´rrez-Dı´ez, V. Gonza´lez Martı´n Dpto. Fı´sica de Materiales, Facultad Ciencias Fı´sicas, Universidad Complutense de Madrid, Ciudad Universitaria s/n, Madrid 28040, Spain Received 5 July 2011; received in revised form 22 November 2011; accepted 22 November 2011 Available online 7 February 2012
Abstract In this work we have investigated (Fe3Ga/TbFe2)n multilayers grown by sputtering at room temperature. These multilayers exhibit a large coercivity associated to the crystalline TbFe2 Laves phase. To reduce the coercivity it is necessary to control the crystallization of that material. In this work, we focus on the analysis of the properties of the TbFe2 layers. In the as-grown heterostructures we have found evidence of nanoaggregates in the TbFe2 layers. The Fe3Ga thickness and the thermal treatments have an influence on the volume of these nanoprecipitates. In the annealed samples, when increasing the Fe3Ga thickness we observe a decrease in the nanoaggregate volume and thus in the coercivity. The experimental results indicate that the crystallization of the TbFe2 depends on the Tb diffusion promoted by the thermal treatment and on the stiffness factor (Y/a) of the Fe3Ga layer. The magnetostrictive properties are also strongly influenced by the crystallization of the TbFe2. We have achieved a maximum magnetostriction constant of nearly 550 ppm with a coercive field close to 400 Oe. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Magnetic thin films; Sputtering; Magnetic properties; Crystallization
1. Introduction Crystalline TbFe2 is the material with the largest known magnetostriction constant at room temperature [1]. The study of this material is of great relevance due to its potential applications in a wide variety of magnetic devices and wireless technology. However, the large coercive field (HC) of the TbFe2 Laves phase [2] constitutes a difficulty in most of these applications. Heterostructures comprising TbFe2 and soft magnetic layers such as Ni, Fe, or FeB have already been studied [3–6]. Although the coercivity of these systems is low, the magnetostriction is also reduced. TbFe2/FeCo multilayers show excellent properties with a saturation magnetoelastic coupling coefficient (k.Y, Y being the Young’s modulus) of 27.5 MPa and a hysteresis of less than 25 Oe [7].
⇑ Corresponding author. Tel.: +34 91 394 50 12; fax: +34 91 394 45 47.
E-mail address: rociran@fis.ucm.es (R. Ranchal).
Galfenol (Fe1 xGax) exhibits a lower k than TbFe2, 440 ppm for quenched bulk samples with a 28% of Ga [8]. Galfenol is a very attractive material for magnetic sensors and actuators, not only because of its pretty high magnetostriction but also due to its low coercivity and high ductility. Several studies have focused on the structural and magnetostrictive properties of the bulk material [9– 17] and more recently, on thin films [18–23]. Heterostructures comprising Fe1 xGax and transition magnetic metal alloys such as NiFe and FeCoB have also been studied [24,25]. A review of Fe1 xGax alloys has recently been reported [26]. In this work we investigate the characteristics of asgrown and annealed (Fe3Ga/TbFe2)n heterostructures. In general, sputtered thin TbFe2 films deposited at room temperature do not exhibit high magnetostrictive properties, and post-growth thermal treatments are necessary to enhance them [27,28]. Moreover, as the TbFe2 crystalline phase is associated to a high coercivity, the control of the Laves phase crystallization is essential to engineer the
1359-6454/$36.00 Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2011.11.049
R. Ranchal et al. / Acta Materialia 60 (2012) 1840–1845
magnetic properties of these heterostructures. It has been reported that the substrate stiffness during the thermal treatment (defined by the Y/a factor, where Y is the Young’s modulus and a is the thermal expansion coefficient) is closely related to the crystallization process of sputtered thin TbFe2 films [27,28]. It is important to remark that the best results are achieved when the substrate or buffer layer has a high Y/a factor but not close to that of TbFe2. Mo is a suitable material for the TbFe2 crystallization process due to its high Y/a factor, 70 106 GPa K. On the other hand, the Tb diffusion promoted by the thermal treatment in (TbFe2/Fe3Ga)n heterostructures has also been studied [29]. This study shows that the annealing promotes the diffusion of Tb atoms from the Fe3Ga layers and/or the Fe3Ga/TbFe interfaces towards the TbFe layers. This Tb diffusion seems to be prevented in heterostructures with thin layers, and thus TbFeGa alloys might be present at the interfaces of these multilayers. Here, we report on the influence of the annealing process on the magnetic and magnetostrictive characteristics of (Fe3Ga/TbFe2)n multilayers. We have focused our investigation on the TbFe2 layers because low coercivity systems can be achieved by controlling the crystallization of the Laves phase. Our results indicate that both the Tb diffusion and the Y/a factor of the Fe3Ga layers have an effect on the magnetic and magnetostrictive properties of the studied (Fe3Ga/TbFe2)n heterostructures. Higher HC and magnetostriction values are achieved in (TbFe2/Fe3Ga)n multilayers because of the modification of the growth order. 2. Experimental techniques Samples were grown by the sputtering technique at room temperature on glass substrates. The deposition was carried out in oblique incidence with an angle between the vapor beam and the perpendicular to the sample of 25°. The Ar pressure during growth was 2 10 3 mbar and the power was 120 W for the deposition of the TbFe2 and 100 W for the Fe3Ga layers, respectively. Mo buffer and capping layers (20 nm) were deposited in all the samples. They were grown with a dc power of 90 W and at an Ar pressure of 2 10 3 mbar. We have deposited two different sets of (Fe3Ga/TbFe2)n heterostructures: (a) a set in which the Fe3Ga layers range from 12.5 to 50 nm and a constant TbFe2 thickness of 50 nm, and (b) multilayers where the Fe3Ga thickness ranged between 12.5 and 50 nm and a constant TbFe2 thickness of 12.5 nm. We have also modified the growth order and then (TbFe2(12.5 nm)/ Fe3Ga(12.5 nm))16 and (TbFe2(50 nm)/Fe3Ga(50 nm))4 heterostructures were also deposited for further analysis. Hysteresis loops at room temperature were carried out in a vibrating sample magnetometer (VSM). The temperature dependence of the magnetization has been characterized by means of SQUID magnetometry. Prior to measuring the temperature dependence of the magnetization, the sample was first cooled from room temperature to 5 K either under a saturation field of 2 kOe (field-cooled: FC) or at zero
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field (zero-field-cooled: ZFC). Then, FC and ZFC curves were recorded with an applied magnetic field of 100 Oe during the warming up. Hysteresis loops at 10 K were measured in the SQUID. Thermal treatments were carried out in Ar atmosphere during 1 h at 400 °C. After the annealing, samples were cooled down at a rate of 20 °C min 1 because fast cooling downs enhance the magnetostrictive properties of the Galfenol layers [8]. The optical cantilever method has been used to infer the magnetostrictive properties of the heterostructures [30]. 3. Results and discussion 3.1. Magnetic properties As-grown TbFe2 thin films deposited in oblique incidence on rigid buffers such as Mo consist of nanoaggregates as revealed by the presence of a cusp in the ZFC curve [28]. In this work, we have also performed FC– ZFC curves to analyze the structural properties of the (Fe3Ga/TbFe2)n heterostructures. In Fig. 1a and b we present these FC–ZFC curves for as-grown heterostructures with 12.5 or 50 nm Fe3Ga and 50 nm TbFe2 layers (set a). In these heterostructures the TbFe2 layers are always deposited on top of the Fe3Ga, and thus the latter acts as the buffer of the former. The higher the Fe3Ga thickness, the lower the influence of the Mo buffer on the TbFe2. When the Fe3Ga thickness is 12.5 nm (Fig. 1a), the ZFC curve shows a cusp at 152 K similar to that observed in single TbFe2 films directly deposited on Mo [28]. The increase of the Fe3Ga thickness up to 50 nm shifts the ZFC cusp to higher temperatures, 192 K, indicating the increase in size of the TbFe2 nanoprecipitates (Fig. 1b). Then, in the asgrown heterostructures the nanoaggregate volume in the TbFe2 layers increases with the Fe3Ga thickness. In Fig. 1c we show an example of the hysteresis loop at low temperature that was obtained for this type of heterostructure. Even when the layer thickness is as low as 12.5 nm we can observe two different magnetic behaviors in the loops: a low coercivity phase that is related to the Fe3Ga and a high coercivity material that corresponds to the TbFe2. In order to calculate the volume of the nanoprecipitates it would be necessary to know the anisotropy field (HK) and the saturation magnetization (Ms) of the TbFe2 layers [28], however these parameters are difficult to infer from the measured hysteresis loops. Then, in this work we will just qualitatively analyze the evolution of the TbFe2 nanoaggregates with the layer thickness and/or the thermal treatments. Fig. 2 shows how the annealing modifies the FC–ZFC curves of the heterostructures. The ZFC cusp of the multilayer with thin Fe3Ga layers, (Fe3Ga(12.5 nm)/ TbFe2(50 nm))7, shifts to 230 K, supporting the idea of an increase in the nanoaggregate volume upon thermal treatment (Fig. 1a). Nevertheless, the ZFC cusp of the sample with thick Fe3Ga layers, (Fe3Ga(50 nm)/TbFe2(50 nm))4, shifts to 146 K, pointing to a decrease of the nanoprecipitate
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Fig. 2. FC (d) and ZFC (s) curves for different annealed heterostructures. (a) (Fe3Ga(12.5 nm)/TbFe2(50 nm))7, and (b) (Fe3Ga(50 nm)/ TbFe2(50 nm))4.
Fig. 1. (a) FC (d) and ZFC (s) curves of the as-grown (Fe3Ga(12.5 nm)/ TbFe2(50 nm))7 heterostructure. (b) FC (d) and ZFC (s) curves of the asgrown (Fe3Ga(50 nm)/TbFe2(50 nm))4 heterostructure. (c) Hysteresis loop at 10 K of the as-grown (Fe3Ga(12.5 nm)/TbFe2(12.5 nm))16 heterostructure.
size (Fig. 2b). The negative magnetization values observed in some ZFC curves (Figs. 1a and 2a) are related to the diamagnetic contribution of the glass substrates. The results of Figs. 1 and 2 show that the Fe3Ga has a different effect on the as-grown and annealed heterostructures. In the as-grown samples we observe that the TbFe2 nanoprecipitate volume increases with the Fe3Ga thickness. However, in the annealed samples we observe the opposite behavior. The higher the Fe3Ga thickness, the lower the nanoprecipitate volume. The effect of the Fe3Ga layers when the samples are annealed can be an indication of the influence of this material on the crystallization of the
TbFe2. The increase of the Fe3Ga thickness seems to prevent the crystallization of the TbFe2. We have measured hysteresis loops at room temperature to analyze the effect of the TbFe2 microstructure on the coercivity of these two (Fe3Ga/TbFe2)n samples of set a (Fig. 3). The as-grown samples show pretty similar hysteresis loops except for the Ms (Fig. 3a). This difference can be explained taking into account that each sample has different layer thicknesses being the Ms of the Fe3Ga higher than that of TbFe2. Also, the low quality that we typically observe in our as-grown sputtered TbFe2 films is expected to decrease the Ms of those layers [27,28]. The HC of these two as-grown heterostructures is pretty similar, 125 Oe for the (Fe3Ga(50 nm)/TbFe2(50 nm))4 and 180 Oe for the (Fe3Ga(12.5 nm)/TbFe2(50 nm))7, respectively. Thus, in the as-grown samples, the different TbFe2 nanoaggregate volume seems not to have a noticeable effect on the coercivity of the samples. After the thermal treatment the HC of the (Fe3Ga (50 nm)/TbFe2(50 nm))4 multilayer slightly increases, whereas the annealed (Fe3Ga(12.5 nm)/TbFe2(50 nm))7 heterostructure shows a dramatic increase in coercivity up to 4 kOe (Fig. 3b). This extremely high HC reflects the crystallization of the TbFe2 in agreement with the increase of the TbFe2 nanoaggregate volume observed in Fig. 2a. On the other hand, the decrease of the nanoprecipitate size upon
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Fig. 4. Coercive field (HC) as a function of the Fe3Ga layer thickness (t) for as-grown (s) and annealed (d) (Fe3Ga(t)/TbFe2(12.5 nm))n heterostructures. t = 0 refers to the (TbFe2(12.5 nm)/Fe3Ga(12.5 nm))16 heterostructure.
Fig. 3. Hysteresis loops for the (d) (Fe3Ga(50 nm)/TbFe2(50 nm))4 and (h) (Fe3Ga(12.5 nm)/TbFe2(50 nm))7 heterostructures. (a) as-grown, and (b) annealed.
the thermal treatment in the (Fe3Ga(50 nm)/TbFe2 (50 nm))4 multilayer can explain the low HC attained in that sample. The increase of the Fe3Ga thickness seems to prevent the TbFe2 crystallization process. This can be due to the low Y/a stiffness factor of the Fe3Ga on the TbFe2 layers or to the Tb diffusion already observed in (TbFe2/Fe3Ga)n heterostructures [29]. In order to clarify the influence of the Fe3Ga thickness, we have studied another set of samples where the Fe3Ga thickness (t) ranges from 0 to 50 nm and the TbFe2 thickness, 12.5 nm, is fixed (set b). In Fig. 4 we present the coercivity of as-grown and annealed (Fe3Ga(t)/TbFe2 (12.5 nm))n heterostructures as a function of the parameter t. Actually, the sample with t = 0 is the (TbFe2(12.5 nm)/ Fe3Ga(12.5 nm))16 heterostructure, t = 0 being used to indicate that the first TbFe2 layer is deposited on top of Mo instead of on top of Fe3Ga. Thus, in the case of t = 0 we have modified the growth order. Regardless of the Fe3Ga thickness the coercivity is similar in all the as-grown samples, HC 190 Oe (Fig. 4). This coercivity is similar to the values observed in the multilayers previously discussed (set a). The thermal treatment has an influence on the HC of samples belonging to set b (Fig. 4). The annealed samples with a Fe3Ga thickness higher than that of TbFe2 (t > 12.5 nm) have a coercivity that is even lower than that
before the thermal treatment. The multilayer with t = 12.5 nm, i.e. (Fe3Ga(12.5 nm)/TbFe2(12.5 nm))16, does not present a noticeable modification of the HC, and the HC increases up to 400 Oe in the sample with t = 0, the (TbFe2(12.5 nm)/Fe3Ga(12.5 nm))16 multilayer. The deposition of the first TbFe2 layer on top of Mo because of the modification of the growth order doubles the HC. To confirm this result we have summarized in Table 1 the HC obtained in heterostructures with the same layer thicknesses but different growth order. In all cases, the coercivity is increased when the first TbFe2 layer is deposited on top of Mo, i.e. the HC is higher in the (TbFe2/Fe3Ga)n heterostructures (Table 1). In a previous work, we present experimental evidences of the accumulation of Tb atoms in the Fe3Ga/TbFe2 interfaces and/or Fe3Ga layers during the growth by sputtering of (TbFe2/Fe3Ga)n multilayers [29]. As-grown heterostructures have Tb–Fe layers with a Tb concentration lower than the nominal 33%. The thermal treatment can promote the diffusion of these Tb atoms to the Tb–Fe layers. If the Tb atoms diffuse upon the thermal treatment, the concentration of Tb in the Tb–Fe layers increases, being closer to the expected 33%. We also observed that the strong lattice mismatch in (TbFe2/Fe3Ga)n heterostructures with thin layers prevented the Tb diffusion [29]. The minimum in the FC curves is known as the compensation temperature (TComp) and it can be used to determine the composition of the Tb–Fe layers [29]. In this work, the presence of a TComp below 300 K reveals a Tb content lower than 28% in the Tb–Fe layers of as-grown heterostructures (Fig. 1a and b). The shift of the TComp towards higher temperatures Table 1 HC of different annealed heterostructures. Annealed heterostructure
HC (Oe)
[Fe3Ga (12.5 nm)/TbFe2(12.5 nm)]16 [TbFe2(12.5 nm)/Fe3Ga(12.5 nm)]16 [Fe3Ga (50 nm)/TbFe2(50 nm)]4 [TbFe2(50 nm)/Fe3Ga(50 nm)]4
180 400 140 250
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after the thermal treatment indicates the Tb enrichment of the Tb–Fe layers because of the Tb diffusion (Fig. 2). Taking into account Ref. [29], Tb diffusion is only expected in heterostructures comprising Fe3Ga layers with a thickness higher than that of TbFe2, so in this work diffusion is expected in samples with a Fe3Ga thickness above 12.5 nm (t > 12.5 nm). As the HC of the TbFeGa alloys is higher than that of the Fe3Ga, the reduction of these ternary alloys because of the Tb diffusion can explain the decrease of the coercivity upon the thermal treatment observed in Fig. 4 when t > 12.5 nm. Nevertheless, Tb diffusion cannot explain the increase of coercivity obtained when just the growth order is modified (samples with t = 0 and t = 12.5 nm in Fig. 4 and the samples summarized in Table 1). In those samples the layer thickness is the same and so the diffusion process is equivalent. The increase of the coercivity observed in (TbFe2/ Fe3Ga)n in comparison to (Fe3Ga/TbFe2)n multilayers needs to be explained by considering a different effect. In the (TbFe2/Fe3Ga)n multilayers, one TbFe2 layer is deposited on top of Mo that has a high Y/a stiffness factor (70 106 GPa K) and promotes an optimum crystallization of the TbFe2. In the (Fe3Ga/TbFe2)n heterostructures, all the TbFe2 layers are deposited on top of Fe3Ga that has a low Y/a factor (6 106 GPa K) and prevents the crystallization of the TbFe2. This effect is in agreement with the reduction of the TbFe2 nanoprecipitate volume observed in Fig. 2b. Then, to engineer the magnetic properties of heterostructures comprising Fe3Ga and TbFe2 is crucial to control the crystallization of the TbFe2 layers by means of the layer thickness and the growth order.
Fig. 5. (a) Magnetostriction as a function of the applied magnetic field of the (Fe3Ga(12.5 nm)/TbFe2(37.5 nm))7 multilayer. The direction of the magnetic field is parallel (d) and perpendicular (s) to the long side of the cantilever. (b) Magnetostriction constant (k) as a function of the Fe3Ga layer thickness (t) for as-grown (h) and annealed (j) (Fe3Ga(t)/ TbFe2(12.5 nm))n heterostructures. t = 0 refers to the (TbFe2(12.5 nm)/ Fe3Ga(12.5 nm))16 heterostructure.
3.2. Magnetostrictive properties Our goal is to achieve not just low coercivity but also high magnetostrictive systems. Typical magnetostriction for these samples can be seen in Fig. 5a. Fig. 5b shows the dependence of the magnetostriction constant with the Fe3Ga thickness. As previously indicated, t = 0 is used to include the (TbFe2(12.5 nm))/Fe3Ga(12.5 nm))16 heterostructure. The k of the as-grown samples is rather low and in general below 100 ppm. Our as-grown sputtered TbFe2 layers exhibit low magnetostrictive properties and thus the k of these as-grown samples seems to be mostly correlated with the Fe3Ga layers. Bulk quenched Galfenol samples can exhibit a k as high as 440 ppm for a Ga content of 28% but a k of 150 ppm has been reported in sputtered Fe81.6Ga18.4 thin films [20], a value closer to our results in as-grown multilayers. The thermal treatments clearly enhance the magnetostriction of the samples, with a maximum k of 550 ppm in the (TbFe2(12.5 nm)/Fe3Ga(12.5 nm))16 heterostructure being achieved. This enhancement can be due to several reasons. First of all, in the samples with thick Fe3Ga layers, the Tb atoms diffuse and the increase of Tb content in the Tb–Fe layers will produce an enhancement of the magnetostriction. When the Tb diffusion is prevented because the Fe3Ga
thickness is reduced, TbFeGa ternary alloys will be present. In bulk TbFeGa alloys, a large k constant of 1000 ppm has been reported when the Ga content is below 10% [31]. The existence of these ternary alloys can also explain the increase of k when the Fe3Ga thickness is reduced (Fig. 5b). However, when the growth order is modified it is necessary to consider more effects. The higher magnetostriction of the (TbFe2 (12.5 nm)/Fe3Ga(12.5 nm))16 sample seems to be related to the growth of one TbFe2 layer on top of Mo that enhances the crystallization process of that material. In the (Fe3Ga(12.5 nm)/TbFe2 (12.5 nm))16 multilayer all the TbFe2 layers are grown on top of Fe3Ga, which does not promote an optimum crystallization of the TbFe2 and lower k values are achieved. Therefore, the magnetostrictive properties of the studied multilayers depend on the crystallization process of the TbFe2. Although the TbFe2 crystallization can be prevented by means of the layer thickness combination and the growth order, the highest k value is achieved when the crystallization of the TbFe2 is enhanced and the growth order is (TbFe2/Fe3Ga)n. Further experiments need to be done to improve the characteristics of these heterostructures, but promising values of k 550 ppm and a HC of 400 Oe in the (TbFe2(12.5 nm)/Fe3Ga(12.5 nm))16, and
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a k of 225 ppm and a HC of 190 Oe in the (Fe3Ga (12.5 nm)/TbFe2 (12.5 nm))16, have been achieved. 4. Conclusions We have studied the magnetic and magnetostrictive properties of (Fe3Ga/TbFe2)n heterostructures. In particular, we have focused our investigation on the crystallization process of the TbFe2 layers due to the high coercivity of the Laves phase. We have found that the magnetic and magnetostrictive properties depend on both the layer thickness and the growth order. On the one hand, the TbFe2 crystallization can be prevented by increasing the Fe3Ga thickness. On the other hand, enhanced magnetostrictive properties have been achieved in (TbFe2/Fe3Ga)n. In these latter multilayers the crystallization of the TbFe2 is improved as a consequence of the modification of the growth order. A pretty high magnetostriction constant of nearly 550 ppm has been obtained when the coercivity is 400 Oe. Acknowledgements This work has been financially supported by Universidad Complutense de Madrid (UCM) and Madrid Regional Government through the project CCG10-UCM/MAT4621. We thank “CAI Te´cnicas Fı´sicas” of UCM for SQUID characterization. We also acknowledge Prof. C. Aroca and Dr M. Maicas for fruitful discussions. References [1] [2] [3] [4]
Clark AE, Belson HS. Phys Rev B 1972;5:3642. Clark AE. Appl Phys Lett 1973;23:642. Lee HS, Cho C. IEEE Trans Magn 2008;44:4049. Quandt E, Ludwig A, Mencik J, Nold E. J Alloys Compd 1997;258:6232. [5] Shima T, Takanashi K, Fujimori H. J Magn Magn Mat 2002;239:573.
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[6] Prados C, Panagiotopoulus I, Hadjipanayis GC, Hernando A, Freijo JJ. IEEE Trans Magn 1997;33:3712. [7] Quandt E, Ludwig A. Sensors Actuat A 2000;81:275. [8] Clark AE, Wun-Fogle M, Restorf JB, Lograsso TA, Cullen JR. IEEE Trans Magn 2001;37:2678. [9] Clark AE, Restorff JB, Wun-Fogle M, Lograsso T, Schalgel D. IEEE Trans Magn 2000;36:3238. [10] Clark AE, Wun-Fogle M, Restorf JB, Dennis KW, Lograsso TA, McCallum RW. J Appl Phys 2005;97:10M316. [11] Ikeda O, Kainuma R, Ohnuma I, Fukamachi K, Ishida K. J Alloys Compd 2002;347:198. [12] Cao H, Bai F, Li J, Viehland DD, Lograsso TA, Gehring PM. J Alloys Compd 2008;465:244. [13] Xing Q, Lograsso TA. Appl Phys Lett 2008;93:182501. [14] Xing Q, Du Y, McQueeney RJ, Lograsso TA. Acta Mater 2008;56:4536. [15] Cao H, Gehring PM, Devreugd CP, Rodrı´guez-Rivera JA, Li J, Viehland D. Phys Rev Lett 2009;102:127201. [16] Laver M, Mudivarthi C, Cullen JR, Flatau AB, Chen WC, Watson SM, et al. Phys Rev Lett 2010;105:027202. [17] Ruffoni MP, Pascarelli S, Gro¨ssinger R, Sato Turtelli R, BormioNunes C, Pettifer RF. Phys Rev Lett 2008;101:147202. [18] Javed A, Morley NA, Gibbs MRJ. J Magn Magn Mater 2009;321:2877. [19] Morley NA, Javed A, Gibbs MRJ. J Appl Phys 2009;105:07A912. [20] Basantkumar RR, Hills Stadler BJ, Robinson WP, Summers EM. IEEE Trans Magn 2006;42:3102. [21] Weston JL, Butera A, Lograsso TA, Shamsuzzoha M, Zana I, Zangari G, et al. IEEE Trans Magn 2000;38:2832. [22] Butera A, Go´mez J, Weston JL, Barnard JA. J Appl Phys 2005;98:033901. [23] Dunlap RA, Deschamps NC, Mar RE, Farrell SP. J Phys: Condens Matter 2006;18:4907. [24] Lupu N, Chiriac H, Pascariu P. J Appl Phys 2008;103:07B511. [25] Lupu N, Pascariu P, Gherasim C, Chiriac H. IEEE Trans Magn 2008;44:3005. [26] Atulasimha J, Flatau AB. Smart Mater Struct 2011;20:043001. [27] Ranchal R, Prieto JL, Sa´nchez P, Aroca C. J Appl Phys 2010;107:113918. [28] Ranchal R, Lo´pez E, Prieto JL, Aroca C. Acta Mater 2011;59:2865. [29] Ranchal R, Gonza´lez-Martı´n V. J Appl Phys 2011;110:053901. [30] Ciudad D, Prieto JL, Lucas I, Aroca C, Sa´nchez P. J Appl Phys 2007;101:043907. [31] Tang YJ, Luo HL, Gao NF, Liu YY, Pan SM. Appl Phys Lett 1995;66:388.