France



The current interest in developing solid-state refrigeration devices has encouraged the investigation of caloric effects linked to first-order phase transitions in different families of solids [1]. Among them, (Ni, Mn)-based Heusler-type magnetic shape memory alloys have deserved particular attention due to their multicaloric nature [2]. For properly chosen compositions, martensitic transformation (MT) in Ni-Fe-Ga alloys appears around room temperature (RT) [1]. This, together with their good mechanical properties in comparison with other (Ni,Mn)-based alloys (improved with the precipitation of γ phase [3]), and the facts that with the partial replacement of Fe by Mn in the Ni&lt;sub&gt;55&lt;/sub&gt;Fe&lt;sub&gt;19&lt;/sub&gt;Ga&lt;sub&gt;26&lt;/sub&gt; alloy γ phase remains in a low concentration, increases the entropy of the MT, not significantly modifying the MT temperature [3], motivated present investigation. Two bulk Ni&lt;sub&gt;55&lt;/sub&gt;Fe&lt;sub&gt;11&lt;/sub&gt;Mn&lt;sub&gt;7&lt;/sub&gt;Ga&lt;sub&gt;27 &lt;/sub&gt;samples were prepared from pure elements (≥ 99.9 %) by arc-melting and spark plasma sintering (SPS) (bulk and SPS samples, respectively), and their phase constitution, as well as magnetocaloric (MC) and elastocaloric (eC) effects studied across the MT. The SPS sample was a dense disc (ρ ≈ 8.4 g.cm&lt;sup&gt;-3&lt;/sup&gt;; ∼ 99 % in respect of theoretical), with diameter and height of 15 and 2.2 mm, sintered from a powdered arc-melted ingot at constant pressure of 30 MPa in a Labox-210 SPS system. The time evolution of vertical displacement and temperature during the process is given in Fig. 1(a). Samples received homogenization annealing at 1273 K for 24 h, followed by a chemical ordering treatment at 773 K (1 h).&lt;br/&gt;Start and finish temperatures of the structural transition were determined from the 5 and 95 %, respectively, of the DSC peak area. RT phase constitution was determined by X-ray diffraction (XRD). As Fig. 1(b) shows, the structural transition for both samples is close to RT with similar temperatures. This highlights the effectiveness of SPS technique for consolidating (Ni-Fe-Ga)-based alloys. XRD patterns, shown in Figs. 1(c) and (d), respectively, were obtained with the X-ray beam impinging on a plane perpendicular to the gradient during solidification and on the surface of the sintered disc. In both patterns, a L2&lt;sub&gt;1&lt;/sub&gt;-type austenite (AST) is the major phase, while cubic γ phase [4] is also detected (∼ 6 %vol.). However, notice that the peak intensities in the XRD pattern for the SPS sample correspond to an isotropic polycrystal, whereas for the bulk arc melted one AST grains preferentially grow with the [1 0 0] direction oriented along the thermal gradient during solidification. Partial preferential and non-preferential grain orientation in both samples was confirmed by SEM observations (not shown).&lt;br/&gt;Magnetization measurements were done by VSM in a PPMS Dynacool system. &lt;i&gt;M&lt;/i&gt;(&lt;i&gt;T&lt;/i&gt;) curves, under 5 mT and 5 T following zero-field-cooling (ZFC) and field-cooling (FC) regimens (1.0 Kmin&lt;sup&gt;-1&lt;/sup&gt;), are depicted in Fig. 2(a). The magnetic entropy change Δ&lt;i&gt;S&lt;/i&gt;&lt;sub&gt;M&lt;/sub&gt; was estimated from the Maxwell relation through the reverse MT from a set of &lt;i&gt;M&lt;/i&gt;(&lt;i&gt;T&lt;/i&gt;) curves (shown in Fig. 2(b)), measured up to an applied magnetic field μ&lt;sub&gt;0&lt;/sub&gt;&lt;i&gt;H&lt;/i&gt;&lt;sub&gt;max &lt;/sub&gt;of 5 T. As shown in Fig. 2(c), bulk sample exhibits a moderate │Δ&lt;i&gt;S&lt;/i&gt;&lt;sub&gt;M&lt;/sub&gt;│&lt;sup&gt;max&lt;/sup&gt; value of 8.2 (12.9) Jkg&lt;sup&gt;-1&lt;/sup&gt;K&lt;sup&gt;-1&lt;/sup&gt; under μ&lt;sub&gt;0&lt;/sub&gt;Δ&lt;i&gt;H&lt;/i&gt;&lt;sub&gt;max &lt;/sub&gt;= 2 (5) T (tough │Δ&lt;i&gt;S&lt;/i&gt;&lt;sub&gt;M&lt;/sub&gt;│&lt;sup&gt;max&lt;/sup&gt; values measured are the largest reported for alloy studied [3]). The eC effect adiabatic temperature change Δ&lt;i&gt;T&lt;/i&gt;&lt;sub&gt;ad&lt;/sub&gt;&lt;sup&gt;me&lt;/sup&gt; was determined in a home-made facility under a compression uniaxial load change of 120 MPa for parallelepiped-shaped samples cut from the bulk and SPS samples; uniaxial stress for the bulk sample was applied along thermal gradient during solidification. Fig. 2(d) gives the Δ&lt;i&gt;T&lt;/i&gt;&lt;sub&gt;ad&lt;/sub&gt;&lt;sup&gt;me&lt;/sup&gt;(&lt;i&gt;T&lt;/i&gt;) curves for both samples loading and unloading the uniaxial stress. The bulk sample exhibits a │Δ&lt;i&gt;T&lt;/i&gt;&lt;sub&gt;ad&lt;/sub&gt;&lt;sup&gt;me&lt;/sup&gt;│&lt;sup&gt;max&lt;/sup&gt; of 2.4 K at a load removal of 120 MPa, whereas this magnitude significantly decreases for the SPS sample. This result agrees with the around 1/3 reduction of the eC effect expected for the direction [1 1 0] [5,6]. Hence, the above-mentioned partial crystallographic texture and isotropic nature of the samples, explain the significant difference observed in the eC effect value.&lt;br/&gt;In summary, apart from probing the efficacy of SPS technique for the shaping and consolidation of Ni-Fe-Ga alloys exhibiting similar phase transformation temperatures than those obtained for the bulk sample, the MC and eC effects were assessed for bulk samples of the Heusler alloy Ni&lt;sub&gt;55&lt;/sub&gt;Fe&lt;sub&gt;11&lt;/sub&gt;Mn&lt;sub&gt;7&lt;/sub&gt;Ga&lt;sub&gt;27&lt;/sub&gt;. Whereas the MC effect is moderate in comparison with the reported for (Ni-Mn-X)-based alloys (X= Ga, In, Sn, Sb), our results underline their good mechanical strength and the relevant role of crystallographic texture on the eC effect.

**References:**

[1] Ll. Mañosa, A. Planes, Adv. Mater. 29 (2017) 1603607.&lt;br/&gt;[2] Ll. Mañosa, A. Planes, Phil. Trans. R. Soc. A 374 (2016) 20150310.&lt;br/&gt;[3] S. Kumar Sarkar, A. Biswas, P.D. Babu, S.D. Kaushik , A. Srivastava, V. Siruguri, M. Krishnan, J. Alloys Compd. 586 (2014) 515–523.&lt;br/&gt;[4] T. Omori, N. Kamiya, Y. Sutou, K. Oikawa, R. Kainuma, K. Ishida, Mater. Sci. Engin. A 378 (2004) 403–408.&lt;br/&gt;[5] F. Xiao, M. Jin, J. Liu, X. Jin, Acta Mater. 96 (2015) 292–300.&lt;br/&gt;[6] F. Hernández-Navarro, J.P. Camarillo-Garcia, C.O. Aguilar-Ortiz, H. Flores-Zúñiga, D. Ríos, J.G. González, P. Álvarez-Alonso, Appl. Phys. Lett. 112 (2018) 164101.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/01c19f0fdf5abb60c13512568d4f30f1.jpg)
Fig. 1. (a) Vertical displacement and temperature versus time along the SPS process. (b) DSC curves for bulk and SPS samples (10 K.min&lt;sup&gt;-1&lt;/sup&gt;). Starting and finishing MT temperatures for the bulk (SPS) sample were M&lt;sub&gt;S&lt;/sub&gt;= 276 (278) K, M&lt;sub&gt;f&lt;/sub&gt;= 269 (262) K, A&lt;sub&gt;S&lt;/sub&gt;= 277 (279) K, and A&lt;sub&gt;f&lt;/sub&gt;= 286 (290) K. (c) and (d) indexed XRD patterns for bulk and SPS samples, respectively (Cu&lt;i&gt;K&lt;/i&gt;&lt;i&gt;&lt;sub&gt;α&lt;/sub&gt;&lt;/i&gt;).

![](https://assets.underline.io/uploads/markdown_image/1/image/62373c7b9f1fcd916d158bd34a413bff.jpg)
Fig 2. ZFC and FC &lt;i&gt;M&lt;/i&gt;(&lt;i&gt;T&lt;/i&gt;) curves under 5 mT and 5 T across the structural transition (a), isofield &lt;i&gt;M&lt;/i&gt;(&lt;i&gt;T&lt;/i&gt;) curves under fields from 1 to 5 T (b), and D&lt;i&gt;S&lt;/i&gt;&lt;sub&gt;M&lt;/sub&gt;(&lt;i&gt;T&lt;/i&gt;) curves (c) through the reverse MT for the bulk arc melted alloy. (d) Δ&lt;i&gt;T&lt;/i&gt;&lt;sub&gt;ad&lt;/sub&gt;&lt;sup&gt;me&lt;/sup&gt;(&lt;i&gt;T&lt;/i&gt;) for both samples under a load and unload applied stress of 120 MPa. Inset: Δ&lt;i&gt;T&lt;/i&gt;&lt;sub&gt;ad&lt;/sub&gt;&lt;sup&gt;me&lt;/sup&gt; as a function of time at 297 K for the bulk sample.

INTERMAG 2021

 Elastocaloric and magnetocaloric effects through the martensitic transformation in bulk Ni55Fe11Mn7Ga27 alloys produced by arc-melting and spark plasma sintering

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The current interest in developing solid-state refrigeration devices has encouraged the investigation of caloric effects linked to first-order phase transitions in different families of solids [1]. Among them, (Ni, Mn)-based Heusler-type magnetic shape memory alloys have deserved particular attention due to their multicaloric nature [2]. For properly chosen compositions, martensitic transformation (MT) in Ni-Fe-Ga alloys appears around room temperature (RT) [1]. This, together with their good mechanical properties in comparison with other (Ni,Mn)-based alloys (improved with the precipitation of γ phase [3]), and the facts that with the partial replacement of Fe by Mn in the Ni55Fe19Ga26 alloy γ phase remains in a low concentration, increases the entropy of the MT, not significantly modifying the MT temperature [3], motivated present investigation. Two bulk Ni55Fe11Mn7Ga27 samples were prepared from pure elements (≥ 99.9 %) by arc-melting and spark plasma sintering (SPS) (bulk and SPS samples, respectively), and their phase constitution, as well as magnetocaloric (MC) and elastocaloric (eC) effects studied across the MT. The SPS sample was a dense disc (ρ ≈ 8.4 g.cm-3; ∼ 99 % in respect of theoretical), with diameter and height of 15 and 2.2 mm, sintered from a powdered arc-melted ingot at constant pressure of 30 MPa in a Labox-210 SPS system. The time evolution of vertical displacement and temperature during the process is given in Fig. 1(a). Samples received homogenization annealing at 1273 K for 24 h, followed by a chemical ordering treatment at 773 K (1 h). Start and finish temperatures of the structural transition were determined from the 5 and 95 %, respectively, of the DSC peak area. RT phase constitution was determined by X-ray diffraction (XRD). As Fig. 1(b) shows, the structural transition for both samples is close to RT with similar temperatures. This highlights the effectiveness of SPS technique for consolidating (Ni-Fe-Ga)-based alloys. XRD patterns, shown in Figs. 1(c) and (d), respectively, were obtained with the X-ray beam impinging on a plane perpendicular to the gradient during solidification and on the surface of the sintered disc. In both patterns, a L21-type austenite (AST) is the major phase, while cubic γ phase [4] is also detected (∼ 6 %vol.). However, notice that the peak intensities in the XRD pattern for the SPS sample correspond to an isotropic polycrystal, whereas for the bulk arc melted one AST grains preferentially grow with the [1 0 0] direction oriented along the thermal gradient during solidification. Partial preferential and non-preferential grain orientation in both samples was confirmed by SEM observations (not shown). Magnetization measurements were done by VSM in a PPMS Dynacool system. M(T) curves, under 5 mT and 5 T following zero-field-cooling (ZFC) and field-cooling (FC) regimens (1.0 Kmin-1), are depicted in Fig. 2(a). The magnetic entropy change ΔSM was estimated from the Maxwell relation through the reverse MT from a set of M(T) curves (shown in Fig. 2(b)), measured up to an applied magnetic field μ0Hmax of 5 T. As shown in Fig. 2(c), bulk sample exhibits a moderate │ΔSM│max value of 8.2 (12.9) Jkg-1K-1 under μ0ΔHmax = 2 (5) T (tough │ΔSM│max values measured are the largest reported for alloy studied [3]). The eC effect adiabatic temperature change ΔTadme was determined in a home-made facility under a compression uniaxial load change of 120 MPa for parallelepiped-shaped samples cut from the bulk and SPS samples; uniaxial stress for the bulk sample was applied along thermal gradient during solidification. Fig. 2(d) gives the ΔTadme(T) curves for both samples loading and unloading the uniaxial stress. The bulk sample exhibits a │ΔTadme│max of 2.4 K at a load removal of 120 MPa, whereas this magnitude significantly decreases for the SPS sample. This result agrees with the around 1/3 reduction of the eC effect expected for the direction [1 1 0] [5,6]. Hence, the above-mentioned partial crystallographic texture and isotropic nature of the samples, explain the significant difference observed in the eC effect value. In summary, apart from probing the efficacy of SPS technique for the shaping and consolidation of Ni-Fe-Ga alloys exhibiting similar phase transformation temperatures than those obtained for the bulk sample, the MC and eC effects were assessed for bulk samples of the Heusler alloy Ni55Fe11Mn7Ga27. Whereas the MC effect is moderate in comparison with the reported for (Ni-Mn-X)-based alloys (X= Ga, In, Sn, Sb), our results underline their good mechanical strength and the relevant role of crystallographic texture on the eC effect.

**References:**

[1] Ll. Mañosa, A. Planes, Adv. Mater. 29 (2017) 1603607. [2] Ll. Mañosa, A. Planes, Phil. Trans. R. Soc. A 374 (2016) 20150310. [3] S. Kumar Sarkar, A. Biswas, P.D. Babu, S.D. Kaushik , A. Srivastava, V. Siruguri, M. Krishnan, J. Alloys Compd. 586 (2014) 515–523. [4] T. Omori, N. Kamiya, Y. Sutou, K. Oikawa, R. Kainuma, K. Ishida, Mater. Sci. Engin. A 378 (2004) 403–408. [5] F. Xiao, M. Jin, J. Liu, X. Jin, Acta Mater. 96 (2015) 292–300. [6] F. Hernández-Navarro, J.P. Camarillo-Garcia, C.O. Aguilar-Ortiz, H. Flores-Zúñiga, D. Ríos, J.G. González, P. Álvarez-Alonso, Appl. Phys. Lett. 112 (2018) 164101.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/01c19f0fdf5abb60c13512568d4f30f1.jpg)
Fig. 1. (a) Vertical displacement and temperature versus time along the SPS process. (b) DSC curves for bulk and SPS samples (10 K.min-1). Starting and finishing MT temperatures for the bulk (SPS) sample were MS= 276 (278) K, Mf= 269 (262) K, AS= 277 (279) K, and Af= 286 (290) K. (c) and (d) indexed XRD patterns for bulk and SPS samples, respectively (CuKα).

![](https://assets.underline.io/uploads/markdown_image/1/image/62373c7b9f1fcd916d158bd34a413bff.jpg)
Fig 2. ZFC and FC M(T) curves under 5 mT and 5 T across the structural transition (a), isofield M(T) curves under fields from 1 to 5 T (b), and DSM(T) curves (c) through the reverse MT for the bulk arc melted alloy. (d) ΔTadme(T) for both samples under a load and unload applied stress of 120 MPa. Inset: ΔTadme as a function of time at 297 K for the bulk sample.

poster

The magnetic nanoparticles can develop superparamagnetic properties in certain conditions1 and this feature attracted the researchers that began to study this class of materials in order to exploit their magnetic properties for applications. The main magnetic peculiarities which make these samples so interesting are the high magnetization and near zero coercive field that allow to use the magnetic nanoparticles for biomedical purposes2,3. Among all the magnetic nanoparticles, the Fe3O4 (magnetite) ones are very attractive due to their biocompatibility and because they show a low toxicity suggesting their possible use in different scopes4,5. Unfortunately, the magnetite nanoparticles are very sensitive to the action of oxygen, thus some of them might undergo oxidation to ferric hydroxide (Fe(OH)3) or to maghemite (γ-Fe2O3) phases. To limit this undesired effect and also the aggregation of primary nanoparticles of low dimension (10-30 nm), specific surface treatments and techniques of encapsulation are currently used to allow their utilization6. In this work, two samples of Fe3O4 magnetic nanoparticles coated with oleic acid (Fe3O4-OA) and citric acid (FeO4-CA) have been studied by means of DC magnetization measurements as a function of temperature (T) in Zero Field Cooling (ZFC) and Field Cooling (FC) conditions at different magnetic fields. The measurements have been performed by means of a Quantum Design PPMS equipped with a VSM option. The M(T) measurements have shown a superparamagnetic behavior in both the samples although with different characteristics which have been highlighted studying the field dependence of the blocking and irreversibility temperature. These two temperatures give important information on the nanoparticles distribution and their dispersion together with information on the diameter of the nanoparticles. In particular, the blocking temperature TB can be identified by taking the maximum of the ZFC curve which, in the case of superparamagnetic samples, has a dome shape behavior. The maximum M value of the ZFC curve can be considered as the blocking temperature of the entire sample, which, of course, differs from the blocking temperatures of the various nanoparticles of different sizes present in the sample. It can be considered as a sort of average blocking temperature of the sample useful when these materials are considered in view of applications. On the other hand, to estimate the TB distribution of the sample it is possible to use the procedure reported in our previous paper7. In this work, the attention has been focused only on the TB obtained from the maximum of ZFC curves. In Fig. 1, an example of M(T) measurement at H = 700 Oe has been reported. The panel(a) and (b) show the M(T) behavior of Fe3O4-OA and Fe3O4-CA, respectively. It can be noted in the insets that both the ZFC curves show a dome shape behavior with the presence of the maximum which identifies the TB of the samples. For the sample Fe3O4-OA, TB ≈ 160 K while for the sample Fe3O4-CA TB ≈ 145 K. In this framework, it is worth to underline that, for H = 700 Oe, the samples are in a superparamagnetic state at room temperature since Troom &gt; TB in both the cases. The temperature value where the ZFC and FC curves separate identifies the irreversibility temperature Tirr. In an ideal case where the nanoparticles are of the same dimensions, the TB and the Tirr are equal. Usually, this is quite uncommon and, in the real case, the Tirr represents the temperature at which all the nanoparticles are certainly unblocked, as the reversible phase occurs after it. Specifically, Tirr represents the unblocking moment of the particles with larger dimensions, which, given the bigger volume than the particles with lower dimensions, unblock at higher temperatures. It is evident that in our case TB ≠ Tirr for both the samples which indicates that the nanoparticles have different diameters. In particular, the temperature difference ΔT between Tirr and TB for Fe3O4-OA is lower respect to the Fe3O4-CA one. In fact, Fe3O4-OA has Tirr ≈ 197 K while Fe3O4-CA has Tirr ≈ 282 K, so giving ΔT(Fe3O4-OA) = 37 K and ΔT(Fe3O4-CA) = 137 K indicating a bigger diameters difference among the nanoparticles of the Fe3O4-CA than the Fe3O4-OA. Finally, performing M(T) measurements at several magnetic fields, the magnetic field dependence of TB and Tirr has been determined and fitted with power and exponential equations extracting interesting samples peculiarities.

**References:**

1C.P. Bean and J.D. Livingston, J. Appl. Phys. 30, S120 (1959). 2O. Veiseh, J.W. Gunn, and M. Zhang, Adv. Drug Deliv. Rev. 62, 284 (2010). 3R. Hergt, S. Dutz, R. Müller, and M. Zeisberger, J. Phys. Condens. Matter 18, S2919 (2006). 4H. wang Zhang, Y. Liu, and S. heng Sun, Front. Phys. China 5, 347 (2010). 5J. Huang, L. Bu, J. Xie et al, ACS Nano 4, 7151 (2010). 6A. Ali, H. Zafar, M. Zia et al, Nanotechnol. Sci. Appl. 9, 49 (2016). 7M. Polichetti, M. Modestino, A. Galluzzi et al, Mater. Today Proc. 20, 21 (2020).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/5cc8cd2f476b6ca2d52b8abd6a717159.jpg)
Fig. 1 M(T) ZFC-FC curves for Fe3O4-OA (a) and Fe3O4-CA (b) at H = 700 Oe. The insets show an enlargement of the region near the maximum of ZFC curves and the arrows indicate the blocking temperatures of the samples.

 Field dependence of blocking and irreversibility temperature in Fe3O4 magnetic nanoparticles coated by oleic and citric acid

Recently, studies of structured systems based on magnetic fluids have become popular. These media include magnetic emulsions, magnetorheological suspensions, magnetic elastomers, ferrogels [1]. Due to significant changes in magnetic, mechanical, rheological, optical, and other properties under the action of a field, such systems are referred to as Smart Materials. One of the most typical representatives of such systems are magnetic emulsions - dispersed systems in the form of droplets of a magnetic fluid of micron and submicron size, suspended in other immiscible and non-magnetic fluid media [2,3]. In this work, we present the results of studying the possibilities of using magnetic emulsions with low interfacial tension as magneto-optical media. Low interfacial tension magnetic emulsions were produced as follows. Magnetic fluid based on kerosene with magnetite nanoparticles was mixed with AMG-10 hydraulic oil. The volume concentration of magnetite in the initial magnetic fluid was 5%. The size range of magnetite nanoparticles is 8–20 nm. Depending on the preparation conditions, the size of microdroplets in the magnetic emulsion could vary from 1000 to 10000 nm. To study the optical effects in magnetic emulsions, we used an optical setup consisting of a helium-neon laser with a radiation wavelength of 633 nm. The sample under study was placed in two types of cuvettes: in the form of a cylinder and in the form of a thin flat layer. Transmission spectra of the magnetic emulsion were studied using an ELLIPS-1891 spectral ellipsometer in the spectrophotometer mode in the wavelength range of 400-1050 nm. Exposure to a magnetic field leads to a rapid (within 0.1-0.3 s) drawing of microdroplets of the magnetic emulsion into long needle-like structures. This effect leads to a significant change in the transparency of the magnetic emulsion (Fig. 1), which opens up prospects for the use of such systems as magnetically controlled optical filters. The most pronounced effect is achieved in the longitudinal field (the relative change in transparency reaches 25-30%). In this case, the magnitude of the change in transparency can be controlled not only by the magnitude of the field strength but also by its orientation relative to the light beam. The effect of changing transparency has a pronounced saturation at a magnetic field strength of over 5 kA/m. For light with a wavelength of less than 550 nm, the magnetic emulsion is practically opaque and therefore has a characteristic yellow-brown color. The mechanism for changing the transparency of magnetic emulsions under the action of a field seems to be quite simple. The transparency of the system T is determined by the concentration of particles and the cross-section of light scattering by an individual particle. When the emulsion droplets are stretched along the direction of the field (and the light beam in this case), the attenuation cross-section significantly decreases at a practically unchanged concentration of particles, which leads to a pronounced effect of changing the transparency. When exposed to a magnetic field perpendicular to the direction of propagation of light, a pronounced effect of small-angle anisotropic scattering of light on elongated needle-shaped microdroplets appears in magnetic emulsions (Fig. 2). A qualitative interpretation of the effect of diffraction scattering in polydisperse magnetic emulsions can be based on the representation of the system in a magnetic field as an irregular diffraction grating, which is formed from microdroplets of a magnetic emulsion elongated in a magnetic field. In the absence of a field, the microdroplets have a spherical shape and the intensity of light diffracted by polydisperse drops smoothly decreases from the center to the periphery of the diffraction pattern without pronounced minima. The action of the field in the direction perpendicular to the light beam leads to a strong elongation of the emulsion droplets and the appearance of pronounced anisotropy of small-angle light scattering (light is scattered mainly in the plane perpendicular to the direction of droplet extension). In a magnetic field, the optical properties of magnetic emulsions change significantly. Deformation of drops of a magnetic emulsion in a magnetic field creates an oriented structure in the emulsion, in which pronounced effects of anisotropic attenuation and diffraction scattering of light are observed. A magnetic field directed along the light beam increases the transparency of the emulsion without changing the transmission spectrum. This allows magnetic emulsions with low interfacial tension to be used as neutralizing filters with fine-tuning of light transmission by the magnetic field. The effect of diffraction by needle-shaped microdroplets can be used in magneto-optical sensors to determine the orientation and strength of the magnetic field. The use of magnetic emulsions can significantly expand the range of applications of liquid magnetic nanostructures [4-6]. This work was financially supported by the Ministry of Science and Higher Education of Russia (project no. 0795-2020-007).

**References:**

S. Odenbach, Arch. Appl. Mech., Vol. 86, p. 269 (2016). A.R. Zakinyan and Y.I. Dikansky, J. Magn. Magn. Mater., Vol. 431, p. 103 (2017). I.M. Subbotin and A.O. Ivanov, Magnetohydrodynamics, Vol. 52, No. 1, p. 269 (2016). J. Philip and J.M. Laskar, J. Nanofluids., Vol. 1, p. 3 (2012). V. Mahendran and J. Philip, Appl. Phys. Lett., Vol. 100, 073104 (2012). P. Das, M. Colombo, and D. Prosperi, Colloids and Surfaces B: Biointerfaces.. Vol. 174. p. 42 (2019).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/7d7554bc0e8f618c6e7a7879009bc8a0.jpg)
The transmission spectrum of a magnetic emulsion under the action of a magnetic field

![](https://assets.underline.io/uploads/markdown_image/1/image/d74cdba524b23f06f55a3d6a7bbd89e3.jpg)
Diffraction light scattering in a magnetic emulsion

 Magnetic Emulsions as Prospective Magneto-Optical Media

A magnetic field pulser works by feeding a pulsed signal to a switching device which drives current through an inductor to generate magnetic field. These pulsers can be designed to output different levels of magnetic field intensity depending on its application. At the same time, a short switching time, on both the rise and fall end, is preferred to keep the outputted magnetic field clean and in synced with the initial pulsed signal. This is especially important in some use cases. In application of the magnetic field pulser is in an all-optical communication system. The pulser generates a magnetic field which causes a polarization rotation in the light signal as per the theory of Faraday rotation [1]. In this situation, the ideal magnetic field intensity is relatively high to ensure sufficient polarization rotation for the end sensors to detect and a short switching time to prevent signal deterioration or contamination. Another application of the magnetic field pulser is to perform transcranial magnetic stimulation, a medical treatment for certain mental disorders [2]. Small, controlled pulses of magnetic field are targeted at parts of the brain to stimulate it without the need to perform open surgery. The noninvasive nature of this treatment is safer and more accessible for the patients. For this application, the pulser generates a magnetic field with much lower intensity and a shorter pulse width. These are some of the many fields where a variant of magnetic field pulser is found. The mechanics behind a magnetic field pulser can be perceived as simple on paper. The pulser design can be broken into four parts: 1. Switching device 2. Inductor coil 3. Dissipation loop 4. Capacitor banks These are by no means complicated systems. However, applications such as the all-optical communication system requires a magnetic field pulser with high magnetic field intensity and fast switching. That factor coupled with the restrictions in physical dimensions mean the pulser is using a relatively large inductor coil. This introduces a fair amount of reactance into the circuit hence affecting the switching time. Hence, a balance between the magnetic field intensity and switching time is needed. The larger reactance and parasitic effects also affect the natural response of the circuit. This investigation will look to tackle several of these factors to bring improvements to an aging design. Firstly, this investigation will examine the switching device of the magnetic field pulser. In most cases, a silicon-based transistor is used as the switching device. Recent developments in transistor technologies have enabled production of transistors with higher drain current, superfast switching, and smaller on-resistance. Compared to transistors from yesteryears, these new products are better replacements to achieve the balance as mentioned above. A quick comparison between several pre-selected transistors can be seen in Figure 1 attached below. These are based on the same magnetic field pulser design by Pritchard, Weber, and Mina (2013) albeit with different transistors [3]. Alternately, the technology behind gallium nitride-based (GaN) transistors has make it another potential candidate as the switching device. Compared to silicon-based (Si) transistors, GaN transistors show higher electron mobility. This means a similarly sized GaN transistor (compared to a Si transistor) will have higher drain current, faster switching time, and smaller on-resistance [4]. By comparing the different parameters in the transistor technologies, this investigation looks to present a better pulser design that can output a higher magnetic field intensity with faster switching time compared to prior published designs. Next, this investigation also analyzes the natural response of the magnetic field pulser. As previously mentioned, the natural response of the circuit is affected by the presence of the relatively big inductor coil. As a result, it is common to see either underdamped or overdamped response on either or both rise and fall ends. This investigation will consider ideal compensation components to mitigate these effects. For better understanding, the investigation looks at the rise and fall end as separate RLC circuits. As a result, the new pulser design can produce a magnetic field pulse that is clean on the rise, fall, and steady-state sides as shown in Figure 2 attached below. In the case of the all-optical communication system, the cleaner magnetic field output prevents cross signal contaminations which results in bad light signal to the end sensors. On the other hand, in the case of the transcranial magnetic stimulation, the cleaner magnetic field output ensures no accidental over or under-stimulation of the patient’s brain. In conclusion, this investigation seeks to improve upon the magnetic field pulser design that has been around for a while. The authors will present a design that shows a significant improvement over previously published designs and a paper that shows the procedure of achieving it. The authors are confident the paper can be of tremendous help to readers who are looking into creating their own magnetic field pulser to fit into their individual requirements.

**References:**

1. J. W. Tioh, M. Mina, and R. J. Weber, “All-optical integrated switch utilizing Faraday rotation,” IEEE Trans. Magn., vol. 46, no. 6, pp. 2474–2477, Jun. 2010. 2. J. Selvaraj, P. Rastogi, N. P. Gaunkar, R. L. Hadimani, and M. Mina, “Transcranial magnetic stimulation: Design of a stimulator and a focused coil for the application of small animals,” IEEE Transactions on Magnetics, no. 99, pp. 1–5, 2018. 3. J.W. Pritchard, M. Mina, and R.J. Weber, “Magnetic Field Generator Design for Magneto-Optic Switching Applications,” IEEE Transactions on Magnetics, Vol. 49, No. 7, pp. 4242-4244, 2013. 4. L. F. Eastman and U. K. Mishra, "The toughest transistor yet [GaN transistors]," IEEE Spectrum, vol. 39, no. 5, pp. 28-33, May 2002, doi: 10.1109/6.999791.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/d576a83153d4d12099210e0d49982323.jpg)
Comparing magnetic field pulsers using three different transistors including one from previously published paper

![](https://assets.underline.io/uploads/markdown_image/1/image/25b4de1ed386a824751fab013e51892f.jpg)
Comparing magnetic field pulser before and after adding compensation components

 Design improvements for a magnetic field pulser: A look into switching device and circuit effects

Among a variety of magneto-optical effects we focus on the inverse Faraday effect (IFE) that plays a crusial role for optical magnetic switching. The IFE is the emergence of magnetization in a transparent medium illuminated by a circularly polarized high-frequency electromagnetic wave [1]. It provides a new way for efficient control of magnetization at GHz and THz rates. Quantitatively the IFE magnetic field is determined by the distribution of the electric field inside the magnetic layer [2]: HIFE∼[E×E*]. Recently it has been demonstrated that optical confinement of light in the magnetophotonic crystals (MPC) or in magnetophotonic microcavity leads to the significant increase of the IFE [2-4]. It happens due to the photonic crystal (PC) defect mode existing in such magnetophotonic crystals or microcavities. These states are spectrally located inside the band gaps and reveal themselves as a sharp peaks in transmittance spectra of a nanostructure. These defect modes also provide strong localization of light inside the magnetic layer that leads to the significant enhancement of the IFE. In this work we propose the novel design of the MPC nanostructures with gradual variation of the magnetic layer thickness. So, the spectral position of the defect mode of the MPC nanostructure depends on the spatial position of input light spot on the sample surface. Therefore, the resonant enhancement of the IFE associated with the peak in transmittance band gap also changes with the input light spot position. Besides that, the design of the MPC nanostructure with perforated magnetic layer provide a spatially localized emergence of the IFE in the spots of several microns that seems very promising for practical applications. The principle scheme of the PC nanostructure with smooth magnetic layer consist of the magnetic layer sandwiched between two non-magnetic Bragg mirrors (Fig. 1). Iron-garnet films doped with bismuth (or BIG films) provide a large Faraday effect and have a relatively low optical absorption in the long-wavelengths visible spectral range [2]. Illuminating this nanostructure by the circularly polarized light at the normal incidence one can observe a defect mode, when the resonance peak appears in the bandgap of the transmittance spectrum. The spectral position and amplitude of this peak depend on the thickness of the magnetic layer h. The spectral position of the defect mode changes from 0.6 μm for 120nm-thick magnetic layer to 0.675 μm for 180nm-thick one. So, the gradual tuning of the resonance spectral position can be realized in sample with the gradient change of the BIG layer thickness. The IFE is determined by the electromagnetic field intensity localized in magnetic layer. In Fig. 2 the spectral dependence of the total energy in the BIG layer normalized by the layer thickness is presented for several values of h. Therefore, the magnitude of the inverse Faraday effect can also be detuned by moving the input beam along the gradient direction of the sample. Besides the smooth gradient magnetic layer, the PC nanostructure with the magnetic layer with etched discs is proposed. In the BIG magnetic layer with the constant thickness the discs with diameter of dozen microns are etched. The etching depth is different for various discs. As a result, the thickness of the magnetic layer is smaller in the area of these discs. Therefore, the enhanced IFE occurs at the different frequencies for each disc. This approach allows us to achieve spatial localization of the IFE limited just by the disc diameter. To sum up, we propose two designs of the MPC nanostructures with gradient of the magnetic layer thickness for tunable inverse Faraday effect. The spectral position and amplitude of the transmittance peak and resonant IFE depend on the thickness of the magnetic layer. The first design of the MPC nanostructure has smooth magnetic layer with thickness gradient that allows to tune the spectral position of peak in transmittance spectra and the IFE gradually. The second design has perforated magnetic layer and provide spatially localized emergence of the IFE in the spots of several microns. The sample fabrication and the experimental verification of the reported results is expected. This study was supported by Russian Science Foundation (project No. 19-72-10139).

**References:**

A. Kirilyuk, A.V. Kimel, and T. Raising, Rev. Mod. Phys. Vol. 82, pp. 2731 (2010). M.A. Kozhaev, A.I. Chernov et al., Sci. Rep. Vol. 8, p. 11435 (2018). T. Goto, A.V. Baryshev, M. Inoue et al., Phys. Rev. B, Vol. 79, p. 125103 (2009). M. Inoue, et al., J. Phys. D. Appl. Phys. Vol. 39, p. R151 (2006).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/3d88bf2c81e92be02ad6b0e66654c6dd.jpg)
Fig. 1. The principle scheme of the photonic crystal nanostructure with smooth magnetic layer (the gradient of the magnetic layer thickness isn’t shown).

![](https://assets.underline.io/uploads/markdown_image/1/image/3327d1b959c15f1285f79297fab57267.jpg)
Fig. 2. Spectra of the total energy localized in the magnetic layer of the MPC for various values of h.

 Photonic Crystal Nanostructures with the Magnetic Layer of Gradient Thickness for Optical Magnetic Switching

Introduction We have been trying to develop magneto-optical spatial light modulators (MO-SLMs) driven by current-induced domain wall motion (CIDWM) because MO-SLMs offer a high-frame rate and high-resolution adequate for holographic display applications [1-3]. Each pixel consists of a magnetic nanowire that modulates the beam by MO Kerr effect (MOKE) and two hard magnets (HMs) embedded at both ends of the nanowire. When a pulse current is injected into the nanowire, a magnetic domain of the nanowire expands and shrinks according to the direction of electron flow. We fabricated a single light modulator with a dimension of 1 μm × 2 μm and successfully demonstrated magnetization reversal by pulse current injection [2]. Using the fact that the reversal magnetic field of the nanowires changes depending on the presence or absence of HMs, we fabricated an array device that displays a magnetic grating pattern with different magnetization directions for each pixel row by applying an external magnetic field and succeeded in measuring diffraction characteristics [3]. The SLM should have high diffraction efficiency because the diffracted beam’s intensity provides the brightness of reproduced images. The intensity of the diffracted beam by the magnetic grating is proportional to the MO figure of merit (FOM), which is the product of the reflectivity and the square of the complex Kerr rotation angle of the nanowire. In this study, we optimized the device structure and the incident angle and polarization direction of the incident beam to amplify the FOM and enhance the diffracted beam’s intensity. Experimental method Figure 1(a) shows the structure of the array device. The nanowire material was Gd24Fe76, which was expected to provide large MOKE and magnetization reversal at low current density. The HMs material was a Co/Pd multilayer with perpendicular magnetic anisotropy. The HMs’ magnetization directions were fixed anti-parallel using the different coercivity. The red dotted line indicates a single pixel. The light modulators were arranged at a pitch of 1 μm × 2 μm. HM1 and the nanowires were partially shared with adjacent pixels to expand the aperture ratio. The width and length of the nanowires was 0.5 μm and 3.8 μm, respectively. The widths of HM1 and HM2 were 0.1 μm, and the lengths were 600 and 0.8 μm, respectively. The coercivity of HM1 and HM2 was controlled by fabricating them at different lengths [2]. When the external magnetic field is applied to the array device, only the magnetization direction of the nanowire with HM2 is reversed. Figure 1(b) shows a two-dimensional magnetic grating pattern displayed on the array. Figure 1(c) shows a cross-sectional illustration of the nanowire. Inside the nanowire, the beam is reflected complexly at each interface. We calculated the dependence of the FOM on the thickness of the Si–N cap layer and the incident beam angle using an optical analysis simulator for multilayers (Multilayer from MM Research, Inc. [4]). The thickness of the cap layer was set at 5 nm steps from 5 to 30 nm, and the incident angle was set at 1° steps from 0° to 89° in the vertical direction of the film surface. Figure 1(d) shows the optical setup for measuring the diffracted beam. A linearly polarized laser beam with a wavelength of 532 nm was incident at a tilt in the x-axis direction. The first-order diffracted beam’s intensity was measured using a laser power meter. The magnetic grating pattern was turned on and off by applying an external magnetic field from an electromagnet placed at the back of the array device. The incident angle was set to be between 0° and 15° since the first-order diffraction angle of the array device is 15.4°. The polarization direction of the incident beam was set to p-polarization and s-polarization, and the diffracted beam’s intensity was measured, respectively. Results and discussions Figure 2(a) shows the FOM calculated by the simulator. The FOM is larger when the thickness of the cap layer is thinner, and it has a maximum at an incident angle of around 50 degrees in this device. It is worth noting that the FOM does not depend on the polarization direction. Figures 2(b) and (c) show the diffraction efficiency of the array device when p- and s-polarized beams were incident, respectively. The diffraction efficiency behaved differently depending on the polarization direction of the incident beam. This is attributed to the pixel shape, which is 1 μm × 2 μm and not square. For s-polarized beam, the diffraction efficiency showed the same behavior as that of the FOM, i.e., it was maximum when the thickness of the cap layer was thin, and the incident angle was large. When the thickness of the cap layer was 5 nm, the diffraction efficiency increased by a factor of 5.9 by optimizing the incident angle and polarization direction. These results showed that the diffracted beam’s intensity can be enhanced by amplifying the FOM of Gd–Fe film, including cap and under layers in MO-SLMs. Also, it showed that the maximum intensity of the diffracted beam can be obtained by considering the polarization direction in MO-SLMs with non-square pixel shapes.

**References:**

[1] K. Aoshima et al., Jpn. J. Appl. Phys., Vol.57, p.09TC03 (2018) [2] R. Higashida et al., Jpn. J. Appl. Phys., Vol.59, p.053001 (2020) [3] R. Higashida et al., Opt. Eng., vol. 59(6), p.064104 (2020) [4] M. Mansuripur, J. Appl. Phys., vol. 67, p.6466 (1990)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/44044a584f48cf1201e7cd6c4532bb82.jpg)
Fig. 1 (a) Domain wall motion type light modulator array (b) illustration of the two-dimensional magnetic grating (c) cross-sectional illustration of the nanowire (d) optical setup for measuring the diffracted beam

![](https://assets.underline.io/uploads/markdown_image/1/image/b8f86f6a8b7ad01b4e8abfb026097c9d.jpg)
Fig. 2 (a) Calculated FOM of the Gd–Fe film including cap and under layers, (b) diffraction efficiency of the fabricated array device in the p-polarized incident beam, and (c) s-polarized incident beam

 Enhancement of a Diffracted Beam by Optimizing an Incident Beam and Cap Layer Thickness in a Domain-wall-motion Type Light Modulator Array

Perovskite manganites RxA1-xMnO3 (R = Rare earth, A = alkaline earth: 0&lt;x&lt;1) are well-known for their correlated charge, spin and lattice degrees of freedom [1]. While studies on hole-doped manganites are almost complete, the physical properties of electron-doped manganites are still being understood. Recent studies on R0.15Ca0.85MnO3 (R = Y, Gd and Dy) revealed interesting metamagnetic transition in the antiferromagnetically ordered state across which the samples show insulator to metal-like change in the electrical resistivity and also switches from an inverse magnetocaloric effect (MCE) to normal MCE [2]. In this work, magnetic and magnetocaloric properties of electron-doped Pr0.15Ca0.85MnO3 are studied in order to understand the effect of light rare earth substitution at the Ca-site of antiferromagnetic CaMnO3. Polycrystalline Pr0.15Ca0.85MnO3 has been prepared by solid state reaction. The X-ray diffraction (XRD) data were collected at several temperatures (T) from 300 K down to 25 K (Cu Ka, Rigaku). It is found that Pr0.15Ca0.85MnO3 crystallizes in orthorhombic structure (space group Pnma, no. 62, oP20) at 300 K and orders antiferromagnetically at 147 K (TN). Paramagnetic susceptibility follows Curie-Weiss law. The effective paramagnetic moment (meff) and paramagnetic Curie temperature (qp) values are 4.47 mB/f.u and +118 K respectively. The positive qp value indicates the presence of coexisting ferromagnetic interactions in the system. The field dependent magnetization measured at 5 K in fields up to 70 kOe is almost linear suggesting the presence of dominant antiferromagnetic interactions. The magnetization value at 5 K in 70 kOe field is about 0.11 mB/f.u. only. From low temperature XRD, it is observed that the crystal structure changes to monoclinic (space group P21/m, no.11-2, mP20) across the magnetic transition. No metamagnetic transition is evidenced in Pr0.15Ca0.85MnO3 at 5 K in fields up to 70 kOe suggesting robust antiferromagnetic state. Magnetocaloric effect (MCE) has been estimated in terms of isothermal magnetic entropy change (ΔSm) using magnetization vs field (M-H) isotherms measured around TN. Positive MCE (i.e. inverse MCE) is observed below TN [Fig. 1a]. The maximum ΔSm value is about 6.4 Jkg-1K-1 at 137.5 K for 70 kOe field change. The peak in inverse magnetocaloric effect coincides with the magnetostructural transition temperature. The maximum value of isothermal magnetic entropy change, ΔSmmax, at 142.5 K is plotted as a function of magnetic field change (ΔH) and is fitted to a power law, i.e. ΔSmmax µ (ΔH)n [Fig. 1b] and the ‘n’ value is found to be 1.75. It is recalled that a value of n = 2 is expected for antiferromagnets [3] and the decrease in the ‘n’ value could be indicative of the coexistence of ferromagnetic interactions in the system. Though Pr0.15Ca0.85MnO3 shows a strong electrical resistivity value in the antiferromagnetically ordered state, the electrical resistivity value decreased by several order below TN. Non adiabatic polaron hopping conduction is observed above TN and polaron activation energy is found to be ~40 meV. Electron spin resonance studies were performed from 300 K to down to 7 K reveal Dysonian broadening as a result of the intrinsic electrical conductivity of the material. The g-value observed at 300 K is less than the free electron g-value and it increases with decreasing temperature down to TN. Thermopower measurements showed strong temperature dependence near TN and a large negative thermopower of ~40 mV/K was obtained around 100 K.

**References:**

1. Y. Tokura, Rep. Prog. Phys. 69 (2006) 797 2. Lakshman Dhal, Eesha Andharia, N. Shukla, T. Geetha Kumary, A. K. Nigam, S. K. Malik, P. N. Santhosh and R. Nirmala, J. Magn. Magn. Mater. 474 (2019) 215 3. A. Biswas, S. Chandra, T. Samanta, M. H. Phan, I. Das and H. Srikanth, J. Appl. Phys. 113 (2013) 17A902

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/7c4e653992c0eeae931c3f9dc5f0db42.jpg)
Fig. 1 (a) Isothermal magnetic entropy change (ΔSm) vs temperature for various field changes and (b) Maximum magnetic entropy change vs ΔH at TN and the fit to power law, (ΔH)n, with n = 1.75.

 Inverse Magnetocaloric Effect and the Magnetostructural Transition in Pr0.15Ca0.85MnO3 Manganite

Recently, solid-state magnetic refrigeration (MR) technique has great interest of research as a promising technology based on magnetocaloric effect to replace the conventional cooling techniques. Many literature reviewed La0.45Nd0.25Sr0.3MnO3/5%CuO composite shows convincing magnetocaloric properties. The present study investigate the detail magnetocaloric study of La0.45Nd0.25Sr0.3MnO3/ x Wt% (M)O, (x=0.0, 0.25, and 0.50, M = Cu, Co and Ni) nano-composite samples. Pure phase exchange coupled nanocomposites of magnetically soft oxides were prepared via one-pot autocombustion method and heat treated in air at 1100OC for 12 h. The influence of wt% of transition metal oxide to the composite on the structural, morphology, magnetic and magnetocaloric properties have been investigated using X-ray diffraction, SEM and SQUID. XRD pattern of the composite reveals the presence of a distinct pure and single-phase rhombohedral structure of perovskite-based material La0.45Nd0.25Sr0.3MnO3(LNSMO) with space group R-3c and monoclinic structure of metal oxide(MO) with the ICCD numbers (LNSMO: 01-082-1152). No trace of secondary phase was detected within the sensitivity limit of the experiment. Rietveld refinement was performed using fullprof software to calculate prepared sample’s mass ratio, lattice parameter and cell volume. The lattice parameters aren’t much effected due to presence of very small amount of metal oxide and values are matched with published literature. The SEM techniques were used to identify the obtained nano-structural morphologies. The NiO/CoO/CuO material were located at the grain boundaries with uniform homogeneity due to the uniaxial pressure effect. Temperature dependence of field-cooled (FC) curve exhibits second-order transition with Tc near room temperature. The temperature depending of magnetization as a function of applied field for all the samples set were measured and conformed the paramagnetic behavior above the transition temperature and ferromagnetic behavior below the transition temperature. The isothermal magnetic entropy change (ΔSm) was calculated to evaluate the magnetocaloric property of the material. The ΔSm have a maximum value around the vicinity of the transition temperature and decrease with any further increase of the temperature. The maximum value of ΔSm for 5T field were found near room temperature. All the values of ΔSm are higher than ferrites oxides. The reason could be the grain size and particle distribution. The uniform distribution of NiO and CoO particles in the grain boundaries than CuO gives the enhance value of entropy change with CoO and NiO compare to CuO. Also, the grain size of La0.45Nd0.25Sr0.3MnO3 for CoO has higher as compare to CuO. This is the effectiveness of one pot autocombustion technique in achieving effective growth of fine grains of respective phases. Cooling efficiency can be calculated from the peak position of ΔSM called relative cooling power (RCP). RCP values of our nanoparticle samples being the maximum for 5wt%CoO. These results make our compounds promising materials for magnetic refrigeration technology. Our study suggests that the La0.45Nd0.25Sr0.3MnO3/ x Wt% (A)O, (x=0.0, 0.25, and 0.50, A = Cu, Co and Ni) nano-composite materials with various morphologies could exhibit interesting magnetocaloric properties near room temperatures

**References:**

1. Tishin, A. M., & Spichkin, Y. I. (2016). The magnetocaloric effect and its applications. CRC Press. 2. GschneidnerJr, K. A., Pecharsky, V. K., & Tsokol, A. O. (2005). Recent developments in magnetocaloric materials. Reports on progress in physics, 68(6), 1479. 3. Gschneidner Jr, K. A., & Pecharsky, V. K. (2000). Magnetocaloric materials. Annual Review of Materials Science, 30(1), 387-429.. 4. Fkhar, L., Mahmoud, A., Boschini, F., Schrijnemakers, A., Hamedoun, M., Benyoussef, A., ... & Mounkachi, O. (2020). A study of magnetic and magnetocaloric properties of 0.95 (La0. 45Nd0. 25Sr0. 3MnO3)/0.05 CuO composites prepared by spray drying. Inorganic Chemistry Communications, 119, 108129. 5. Fkhar, L., Lamouri, R., Mahmoud, A., Boschini, F., Hamedoun, M., Ez-Zahraouy, H., ... & Mounkachi, O. (2020). Enhanced Magnetic and Magnetocaloric Properties of La0. 45Nd0. 25Sr0. 3MnO3/CuO Composite. Journal of Superconductivity and Novel Magnetism.. 6. Kallel, N., Kallel, S., Hagaza, A., & Oumezzine, M. (2009). Magnetocaloric properties in the Cr-doped La0. 7Sr0. 3MnO3 manganites. Physica B: Condensed Matter, 404(2), 285-288. 7. Kumaresavanji, M., Sousa, C. T., Pires, A., Pereira, A. M., Lopes, A. M. L., & Araujo, J. P. (2014). Room temperature magnetocaloric effect and refrigerant capacitance in La0. 7Sr0. 3MnO3 nanotube arrays. Applied Physics Letters, 105(8), 083110. 8. Fang, Y. K., Yeh, C. C., Chang, C. W., Chang, W. C., Zhu, M. G., & Li, W. (2007). Large low-field magnetocaloric effect in MnCo0. 95Ge1. 14 alloy. Scripta materialia, 57(6), 453-456.

**Images:**

 Magnetocaloric study of La0.45Nd0.25Sr0.3MnO3/ MO (MO=CuO, CoO and Ni) nanocomposites

The precise study of the bonding in materials is an important part of material characterization. The Maximum Entropy Method (MEM) is a powerful tool for deriving accurate electron density distribution in crystalline materials using experimental data. The MEM provides us with the least biased deduction, which is compatible with certain given information. The bonding nature and the distribution of electrons in the bonding region can be clearly visualized using this technique. MEM electron densities are always positive and even with the limited number of data, one can evaluate reliable electron densities resembling true densities. In this present work, the attention is focused on the growth, structural, and electron density distribution of Gd3Fe5-xAlxO12 (0 ≤ x ≤ 1) using X-ray powder data. Gd3Fe5-xAlxO12; x = 0.0, 0.25, 0.50 and 1.0 samples were synthesized via auto combustion method and heat-treated in air at 1100oC for 12 h. The phase identification and lattice parameter of the powders performed using X-ray diffraction data. The Rietveld analysis was performed for the present analysis, which proves the synthesized material of Gd3Fe5-xAlxO12 (0 ≤ x ≤ 1) stays in high-crystalline with single-phase garnet structure, the structural and profile parameter were compared with the reported results from JCPDS-48-0077. The Rietveld analysis was performed using JANA 2000 software package, the results show that the lattice parameters and the cell volume decrease with the increasing the Al, this is clearly evidencing the Al content is replacing the Fe. From this analysis, we observed that a gradual shift in the peaks from the lower angle to the higher angle as compared to the pure Gd3Fe5O12. The lattice parameters for Gd3Fe5-xAlxO12 decreases from 12.50A to 12.44A due to smaller ionic radii of Al3+ ion (0.535A) replacing Fe3+ ions (0.645A). The atomic site occupancy and bond distance analysis for each site were derived from Rietveld refinement for Gd3Fe5-xAlxO12, the Al is preferred to the tetrahedral site for x ≤ 0.5 and prefers both sites; tetrahedral and octahedral thereafter. The structure factors from the Rietveld analysis were further taken to determine the electron density distribution analysis by MEM method. The bonding nature and radii of atoms/ions and interaction between the atoms were clearly analyzed and represented by 3D, 2D, and 1D MEM maps. The mid bond electron densities between atoms were analyzed. The two-dimensional MEM electron density maps on the (001) and (202) planes, and the one-dimensional electron density profiles along Fe-O, Gd-O, and O-O directions were calculated. From this analysis, the substitution of Al3+ ions actually reduces the lattice parameter and volume of the Gd3Fe5-xAlxO12. The 2D map shows the depletion of charge density in the valence regions are clearly seen. Our study suggests that the Al3+ doped Gd3Fe5O12 nano-materials could exhibit interesting electronic properties for data storage applications.

**References:**

1. Lassri, H., Hlil, E. K., Prasad, S., & Krishnan, R. (2011). Magnetic and electronic properties of nanocrystalline Gd3Fe5O12 garnet. Journal of solid state chemistry, 184(12), 3216-3220. 2. Nakamoto, R., Xu, B., Xu, C., Xu, H., & Bellaiche, L. (2017). Properties of rare-earth iron garnets from first principles. Physical Review B, 95(2), 024434. 3. Ali, K. S., Saravanan, R., Israel, S., & Arda, L. (2010). Localized ferromagnetic charge ordering through charge density analysis in nano sized diluted magnetic semiconductor Co2+: ZnO. Physica B: Condensed Matter, 405(7), 1763-1769. 4. Ali, K. S., Saravanan, R., Pashchenko, A. V., & Pashchenko, V. P. (2010). Local distortion in Co-doped LSMO from entropy-maximized charge density distribution. Journal of Alloys and Compounds, 501(2), 307-312. 5. Saravanan, R., & Rani, M. P. (2011). Metal and alloy bonding-an experimental analysis: Charge density in metals and alloys. Springer Science & Business Media.

**Images:**

 Structural and Electronic properties on Gd3Fe5-xAlxO12 (x= 0.25, 0.5, 1.00) using Rietveld, Maximum Entropy Method (MEM)

The family of manganese oxides with the general formula R1-xAxMnO3 (R=lanthanide and A=divalent alkaline earth) has a non-negligible potential in magnetic cooling particularly because of the high chemical stability and interesting magnetocaloric properties at room temperature [1]. In this work, we report a detailed study of the thermodynamic performance of Pr0.6Sr0.4MnO3 in a functional magnetic cooling system up to the theoretical limit of its magnetocaloric effect. In this context, a one-dimensional numerical model of the AMR cycle is developed, while the required magnetocaloric properties of the oxide are calculated in the framework of Mean Field Theory (MFT). The obtained results are deeply analyzed and then compared with the reference Gadolinium (Gd) metal. Moreover, it is revealed that using Pr0.6Sr0.4MnO3 as refrigerant requires high magnetic fields to generate reasonable temperature spans for application. Under a magnetic field of 8 T, the studied oxide can generate an AMR temperature span of 33 K. However, such a span can be achieved by employing Gd and using only a magnetic field source of 1.45 T. The thermodynamic efficiency was evaluated by directly considering the involved magnetic forces between the Pr0.6Sr0.4MnO3 regenerator and the magnetic field source. The associated magnetic work is numerically calculated by using Altair Flux 3D software. Particularly, it was found that the coefficient of performance (COP) can be enhanced by a factor of two when balancing magnetic forces (Fig. 1) despite of the moderate magnetocaloric effect shown by Pr0.6Sr0.4MnO3 under practical fields (1 to 2 T). On the other hand, the hidden thermodynamic performances of this compound are also pointed out by investigating its magnetocaloric properties in the theoretical limit zone. In this way, our calculations unveil that is possible to reach cold source temperatures under -22°C even when using helium as heat transfer fluid (Fig. 2). Finally, a magnetic cooling device working with Pr0.6Sr0.4MnO3 as refrigerant and fully powered by solar cells is proposed.

**References:**

[1] M. Balli, S. Jandl, P. Fournier, and A. Kedous-Lebouc, Appl. Phys. Rev. 4, 021305 (2017).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/c8bd35af86f03d8a8875379d767e6ec9.jpg)
Fig. 1 The simulated magnetic forces in the case of one and two separated regenerators of Pr0.6Sr0.4MnO3.

![](https://assets.underline.io/uploads/markdown_image/1/image/1c021cfc7ac84b35b464f05f9437c523.jpg)
Fig. 2 Temperature evolution of the hot and cold ends as a function of the number of cycles for using helium as transfer fluid

 On the real potential of R1-xAxMnO3 oxides in magnetic cooling: Pr0.6Sr0.4MnO3 as a case of study

The materials that exhibit mutual control of more than one property by external stimuli such as magnetic, electric and stress fields are known as multifunctional material. They are useful to design smart sensors and can be extended as a functional refrigerant for cooling. The magnetic refrigerator is an alternative to the conventional refrigerators as they do not emit toxic greenhouse gases to the environment. However, the product cost of a magnetic refrigerator is expensive due to the raw material and magnetic components required to achieve the maximum entropy change at a given magnetic field strength (≤1.5 T). Assuming the combination of electric and magnetic properties may yield adiabatic temperature change at low (magnetic/electric) field, the piezoelectric/magnetic nanocomposites are proposed. Poly(methyl methacrylate)/Poly(vinylidenefluoride-co-hexafluoropropylene) (PMMA/PVDF-HFP) blend and PMMA/PVDF-HFP/Zn0.5Cu0.5Fe2O4 flexible nanocomposites were fabricated by solution casting method. Pristine polymer blend and PMMA/PVDF-HFP/ferrite nanocomposite are having 67 and 62% of F(b), respectively. Zn0.5Cu0.5Fe2O4 and PMMA/PVDF-HFP/Zn0.5Cu0.5Fe2O4 nanocomposite show superparamagnetic nature with the saturation magnetization of 44 and 5.3 emu/g, respectively. The composite film is thermally stable up to 753 K and later it shows 80% of weight loss. The magnetic entropy change (DSM) of Zn0.5Cu0.5Fe2O4 and PMMA/PVDF-HFP/Zn0.5Cu0.5Fe2O4 show interesting table-like flat variation with respect to magnetic field (0 to 40 kOe) and temperature (100 to 400 K). The DSM (refrigerant capacity) at 300 K is found to be -0.69 (293 J/kg) and -0.094 (40 J/kg) J/kg/K for Zn0.5Cu0.5Fe2O4 and PMMA/PVDF-HFP/Zn0.5Cu0.5Fe2O4, respectively. Ferroelectric polarization change as a function of electric field is analysed at room temperature. The composite film exhibiting both magnetism and ferroelectricity is proved and would be attractive to design different compositions to increase the temperature variation for cooling devices.

**References:**



**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/db2e5a1e5dec581473008e717fb6b89a.jpg)
Figure 1 (A) X-ray diffraction patterns, (B) and (C) M-H hysteresis loops of Zn0.5Cu0.5Fe2O4 and PP/Zn0.5Cu0.5Fe2O4 nanocomposite respectively.

Elastocaloric and magnetocaloric effects through the martensitic transformation in bulk Ni55Fe11Mn7Ga27 alloys produced by arc-melting and spark plasma sintering

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