Tuning the magnetocaloric properties of La0.7Ca0.3MnO3 manganites through Ni-doping

Accepted Manuscript Tuning the magnetocaloric properties of La0.7 Ca0.3 MnO3 manganites through Ni-doping

A. Gómez, E. Chavarriaga, I. Supelano, C.A. Parra, O. Morán

PII: DOI: Reference:

S0375-9601(18)30102-6 https://doi.org/10.1016/j.physleta.2018.01.030 PLA 24938

To appear in:

Physics Letters A

Received date: Revised date: Accepted date:

5 June 2017 22 January 2018 23 January 2018

Please cite this article in press as: A. Gómez et al., Tuning the magnetocaloric properties of La0.7 Ca0.3 MnO3 manganites through Ni-doping, Phys. Lett. A (2018), https://doi.org/10.1016/j.physleta.2018.01.030

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Highlights • • • • •

La0.7 Ca0.3 Mn1-x Nix O3 (x =0, 0.02, 0.07, and 0.1) are synthesized via auto-combustion. Magnetic and magnetocaloric properties are studied. Systematic decrease in the Curie temperature is observed upon Ni2+ doping. Large magnetic entropy change is achieved for La0.7 Ca0.3 Mn1-x Nix O3 manganites. Ni2+ doping increases the relative cooling power of the parent compound.

Tuning the magnetocaloric properties of La0.7Ca0.3MnO3 manganites through Ni-doping A. Gómez1, E. Chavarriaga1, I. Supelano2, C. A. Parra2, and O. Morán1 1

Universidad Nacional de Colombia, Campus Medellín, Facultad de Ciencias, Departamento de Física, Laboratorio de Cerámicos y Vítreos, A. A. 568, Medellín, Colombia. 2 Universidad Pedagógica y Tecnológica de Colombia, Departamento de Física, Avenida Central del Norte 39-115, Tunja, Colombia Abstract. The effect of Ni2+ doping on the magnetic and magnetocaloric properties of La0.7Ca0.3MnO3 manganites synthesized via the auto-combustion method is reported. The aim of studying Ni2+-substituted La0.7Ca0.3Mn1-xNixO3 (x = 0, 0.02, 0.07, and 0.1) manganites was to explore the possibility of increasing the operating temperature range for the magnetocaloric effect through tuning of the magnetic transition temperature. X-ray diffraction analysis confirmed the phase purity of the synthesized samples. The substitution of Mn3+ ions by Ni2+ ions in the La0.7Ca0.3MnO3 lattice was also corroborated through this technique. The dependence of the magnetization on the temperature reveals that all the compositions exhibit a well-defined ferromagnetic to paramagnetic transition near the Curie temperature. A systematic decrease in the values of the Curie temperature is clearly observed upon Ni2+ doping. Probably the replacement of Mn3+ by Ni2+ ions in the La0.7Ca0.3MnO3 lattice weakens the Mn3+-OMn4+ double exchange interaction, which leads to a decrease in the transition temperature and the magnetic moment in the samples. By using Arrott plots, it was found that the phase transition from ferromagnetic to paramagnetic is second order. The maximum magnetic entropy changes observed for the x = 0, 0.02, 0.07, and 0.1 composites was 0.85, 0.77, 0.63, and 0.59 J/kg K, respectively, under a magnetic field of 1.5 T. In general, it was verified that the magnetic entropy change achieved for

1

La0.7Ca0.3Mn1-xNixO3 manganites synthesized via the auto-combustion method is higher than those reported for other manganites with comparable Ni2+-doping levels synthesized via standard solid state reaction The addition of Ni2+ increases the value of the relative cooling power as compared to that of the parent compound. The highest value of this parameter (~60 J/kg) is found for a Ni-doping level of 2 % around 230 K in a field of 1.5 T.

Corresponding author: E-mail: [email protected] Phone: 57-4-4309327 Fax: 57-4- 2604489

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1. Introduction

ABO3 compounds are oxides with a perovskite-like structure that feature a rich variety of physical properties [1]. The unique properties of these compounds have made them a lively area of research over the last few decades [2]. The parent LaMnO3 compounds behave

like

paramagnetic

(PM)

insulators

at

higher

temperatures

and

antiferromagnetic (AFM) insulators at low temperatures [3]. When the trivalent La is replaced by divalent Ca, Sr or Ba (hole doped) in the range 0.2 d x d 0.4, the material becomes a metallic ferromagnet below the Curie transition temperature (TC) [4, 5]. At T
3

undoped Gd metal [12]. Gd shows a large MCE at temperatures close to room temperature [13], but it is an expensive element and tends to be subject to oxidation. Thus it is evident that high-performance MCE materials with fewer disadvantages are necessary for potential technical applications. For this purpose, it is necessary to investigate various kinds of magnetic materials engineered under varying conditions. In the case of perovskite-type ABO3 manganites, the physicochemical properties of these compounds can be tuned through doping with metallic ions on the A site or the B site [14]. Doping at the Mn site with other transition metals is of great importance in modifying the double exchange (DE) interaction strength between Mn3+ and Mn4+ via oxygen atoms. This alters the magnetic and magnetocaloric behavior of these doped oxides [6]. In some manganites, the doping can also provoke structural transition [15] and suppression of the charge ordering [16]. The partial substitution of Mn3+ by Ni2+ ions in manganites has been previously studied in connection with their magnetoresistive behavior [17, 18]. Magnetic properties and the MCE were recently reported for the La0.7Sr0.3Mn1-xNixO3 system in Ref. [19]. The samples reported in this reference were synthesized via conventional solid-state reaction, which produces grain sizes as large as a1 Pm. In this regard, it is widely known that the dimensionality strongly impacts the magnetic and magnetocaloric response of LCMO [4]. In this regard, previous reports have shown that the properties of perovskite manganites at nanoscale are quite different from those at microscale [20, 21]. Indeed, as the particle size reduces to nanoscale, the manganites feature particular physical phenomena such as super-paramagnetism, surface spin-glass, low saturation magnetization and the low-field magnetoresistance effect [22, 23], among others. It has also been indicated that the downsizing of the particles plays an important role in the determination of the magnetic entropy change [24]. In this context, recent advances in

4

the field of permanent magnets have demonstrated that manipulation of the microstructure of known materials can lead to considerable gains in performance [25]. However, the effect of the morphology on the performance of the materials under consideration for magnetic cooling applications is less well understood. In the present investigation, the results of a systematic study of the MCE on LCMO manganites (synthesized via the auto-combustion method) substituted at the Mn site by Ni2+ are presented. Auto-combustion synthesis is characterized by fast heating rates, achieving high temperatures and short reaction times [26]. It is a straightforward preparation process that is able to produce homogeneous, very fine, crystalline, and unagglomerated multicomponent oxide powders without intermediate decomposition steps [27]. In the solution combustion synthesis, an aqueous solution of the desired metal salts is heated together with a suitable organic fuel until the mixture ignites and a fast combustion reaction starts [28]. The Ni-doped manganite system is of particular interest because the MCE, the metal-insulator (M-I) transition, and the MR effects can be tuned via Ni doping at Mn-sites over a wide range of operating temperatures [29]. In this context, the Ni-doped manganite system allows one to investigate the relation between the MCE and the transport properties, which is of central importance because the resistivity shows an M-I transition near the TC [30]. Moreover, Ni-doped manganite oxides are considered to be very good negative temperature coefficient (NTC) thermistors [31]. The Ni-doped manganite system is also of particular interest for investigating the change in the spin-wave stiffness constant (D) due to weakening of the effective ferromagnetism [32]. Indeed, the strength of the microscopic magnetic coupling of the spins associated with the ferromagnetic clusters can be estimated from the measurement of D [33]. 2. Experiment.

5

La0.7Ca0.3Mn1-xNixO3 (x = 0, 0.02, 0.07, and 0.1) powders were synthesized via the combustion

method

using

glycine

(C2H5NO2)

as

fuel.

(La(NO3)3.6H2O),

(Ca(NO3)2.4H2O) Ni (Ni(NO3)2.6H2O), and (Mn(C2H3O2)2.4H2O) were used as starting materials. All the reactants were of analytical grade. After complete dilution of the starting materials and the fuel in distilled water, based on the stoichiometric quantities, the solution was heated (under magnetic stirring) on a hot plate at around 100° C until water evaporation and gel formation took place. After that, the mixture was continuously heated at around 450° C until ignition of the glycine started. Once ignited, the gel underwent a combustion process and yielded voluminous La0.7Ca0.3Mn1-xNixO3 powders. The obtained powders were then calcined in air at 700° C for 3 h in order to remove the unreacted carbon compounds and the organic material. The combustion reaction can be expressed as: 0.7La(NO3)3+0.3Ca(NO3)2+(1-x)Mn(C2H3O2)2+xNi(NO3)2+yC2H5NO2oLa0.7Ca0.3Mn1-xNixO3

+aCO2+bN2+cH2O In order to get a general idea about the effect of the dimensionality on the physical properties of the studied system, LCMO in polycrystalline and thin film forms was prepared via standard solid state reaction and DC magnetron sputtering, respectively. Details of the preparation of these samples have been given elsewhere [34]. The final products were examined for their crystalline quality via X-ray diffraction (XRD) with a Cu KD radiation source (O=1.5406 Å) and fixed scanning steps of 0.01°. The crystalline structure and the lattice parameters were refined using the Rietveld method with the GSAS program. The surface morphologies of the samples were examined using fieldemission scanning electron microscopy (FE-SEM). Magnetization (M) measurements versus temperature (T) and magnetic field (H) were performed using a vibrating sample magnetometer (Quantum Design). The MCE was estimated in terms of

6

isothermal magnetic entropy change (-'SM), using the magnetization data and ு

డெ

employing the Maxwell relation ȟܵெ ൌ ߤ଴ ‫׬‬଴ ೘ೌೣ ቀ

డ்

ቁ ݀‫ܪ‬ǡ where μ0H is the external

magnetic field [19]. The magnetization isotherms, recorded in temperature intervals of 5 K, were used for the calculations. 3. Results and discussion 3.1. Structure

The XRD patterns in Fig. 1 show that all the samples are single phase. No additional diffraction peaks stemming from impurities or secondary phases can be observed. Additional diffraction peaks have been reported for other Ni2+-doped manganites such as La0.7Sr0.3Mn1-xNixO3 (x = 0, 0.025, 0.050, 0.075, and 0.1) prepared with the conventional solid-state reaction method [35]. The additional peaks disappeared with doping, which was attributed to enhancement of the reaction between elements due to the larger ionic size of Ni2+ compared with Mn3+ [36]. Thus it is apparent that the auto-combustion method is an efficient, time-saving, and inexpensive method for preparing high-quality LCMO powders with Ni2+-doping levels as high as 10%. The element Ni2+ is an interesting transition metal for Mn-site substitution, where the combination of Ni2+ with Mn4+ is favorable [37]. Rietveld refinement of the XRD diffraction patterns was carried out using the GSAS program. A perfect convergence of the experimental and the calculated data was obtained, as can be seen in Fig. 1. As a result, the crystalline structure was indexed in the orthorhombic system with the Pbnm space group. The results of the Rietveld refinement are summarized in Table 1. The variations of the lattice parameters a, b/—2, c, and the unit-cell volume V with Ni2+ doping are shown in Fig. 2. As expected from the slightly larger ionic size of Ni2+ (0.69 Å) compared with those of Mn3+ (0.645 Å) and Mn4+ (0.53 Å) [36], all the cell parameters and the volume exhibit a tendency to grow as the Ni2+ content increases.

7

This is a consistent and important result, which suggests that the Ni ions are mainly in the Ni2+ state and that they replace the Mn3+ ions. Interestingly, some reports have shown that the dependence of the lattice parameters of nonstoichiometric samples on Ni additions is different from that shown in Fig. 2(b) [35]. Indeed, the substitution of Ni for Mn led to a slight decrease in the lattice parameters [36]. It was suggested that a change in the Mn3+/Mn4+ ratio could be responsible for the decrease in the lattice parameters in the oxygen-deficient samples. Here, a small increase in Mn4+ with an ionic radius of 0.530 Å is expected to reduce the unit cell. The average crystallite sizes of the samples, calculated from the Rietveld refinement, are indicated in Table 1. A systematic decrease in the grain size upon Ni2+ doping is clearly in evidence. It is clear that the issue of substitution of Ni for Mn in samples synthetized by auto-combustion (in the present paper) should be further investigated using other powerful analytical techniques such as X-ray photoelectron spectroscopy or X-ray absorption near-edge structure (XANES). 3.2. Morphology

It is well known that the structural, magnetotransport, and magnetic properties of manganites are strongly governed by the nature of the grain morphology and the grain boundary [38]. Figure 3 shows the microstructure of the parent LCMO sample obtained from FE-SEM analysis. The images show that that the morphology of the compound consists of homogenous particles (rounded polyhedrons) with the grain boundaries clearly visible. The particles are abundant and almost uniform in size. The average particle size of the sample observed in the FE-SEM micrograph varies between 20 and 30 nm and is in reasonable agreement with the crystallite size obtained from the XRD pattern. The variation in the grain size as well as the grain boundaries will result in variation of the bond angle, which affects the DE interaction. As a consequence, the

8

orbital overlap and the hopping of electrons between Mn ions of different valences is reduced [39]. 3. 3. Magnetic properties

In this section, the change in the magnetic properties of the La0.7Ca0.3Mn1-xNixO3 system is discussed, taking into account intrinsic and extrinsic factors. The intrinsic factors are represented by a change in magnetic interactions such as the ferromagnetic DE interactions (Mn3+-O-Mn4+) or a change in the internal structure [40], while the extrinsic factors can arise, e.g., from a change in the grain size [41]. Figure 4(a) shows the field-cooled (FC) temperature dependence of the magnetization, M(T), taken at an applied field of 1000 Oe for La0.7Ca0.3Mn1-xNixO3 (x = 0, 0.02, 0.07, and 0.1) samples. It can be seen in Fig. 4(a) that all the samples undergo a PM–FM transition near the Curie temperature (TC), which is defined as the peak of the dM/dT vs. the T curve [upper inset of Fig. 4(a)]. Although no appreciable transition broadening is observed, the value of the TC is systematically reduced with increasing Ni2+ doping levels. Furthermore, the sample with x=0.07 features a reduced magnetization saturation. On the basis of the DE mechanism, the partial substitution of Mn 3+ ions by Ni2+ disturbs the Mn3+–O–Mn4+ chains, which leads to a change in the Mn3+/Mn4+ ratio [42]. As a consequence of this variation, the DE is weakened, and therefore the TC and the magnetic moment will decrease. In addition, the competition between FM and AFM exchange interactions is reinforced [43]. In particular, the AFM interaction between different Mn-Mn, Mn-Ni, and Ni-Ni pairs is increased due to the increase in Mn4+ caused by the appearance of Ni2+ ions in the LCMO lattice [44]. More precisely, the substitution of Mn3+ by Ni2+ ions can reduce the number of available hopping sites and create cuts in the conduction path, which weakens the DE interaction. Here, primarily it should be taken into account that the FM Ni2+-O-Mn4+ exchange interaction 9

may increase the value of the TC [45]. In turn, due to a parallel arrangement between ଶȀଷ

the Ni and Mn ion spins, the AFM Ni2+(݁௚ )–O–Mn3+(݁௚ଵ ) superexchange interaction ௚ଵ

may enhance the exchange energy of the ݁௚ spin, which could impede an eg electron from hopping to the Mn4+ sites surrounding the Ni2+ ion [46]. Thus the number of available hopping sites is reduced, and the Mn3+-O-Mn4+ DE interaction is suppressed. So although the issue of determining the charge distribution and exchange interactions in the Ni-doped manganites is far from being trivial, it is apparent that Ni2+ doping suppresses DE because this ion does not participate in the DE mechanism. As stated above, new AFM bonds such as Mn3+–O–Ni2+, Mn4+–O–Mn4+, and Ni2+–O–Ni2+, generated upon Ni2+ doping, are non-DE interactions and therefore promote AFM coupling [19]. The promotion of AF coupling then weakens the DE and lowers the TC [47, 48]. The AFM bonds may increase in number, with higher Ni2+ doping levels leading to a systematic decrease in the value of the TC, as can be seen in Fig. 4(a). The role of the internal structure in the change of the magnetic properties of Ni-doped manganites can be visualized through the gradual increase in the orthorombic distortion due to the Jahn-Teller effect, which stabilizes charge ordering in competition with ferromagnetism [49]. Indeed, doping at the Mn-site can induce a change in the angle (T) of the Mn ions [50]. In short, deviation from 180° of the Mn–O–Mn angle increases the distortion and decreases the transfer integral t = t0cos(T/2) (where t0 is the maximum value of t), which in turn decreases the DE between Mn ions. Nevertheless, it should be pointed out that at low Ni-doping levels, the structural changes are small, as suggested by the structural parameters summarized in Table 1. As previously mentioned, perovskite manganites have been the object of much research interest due to their tunable TC and saturation magnetization using

10

doping/substitution [51]. The lower inset of Fig. 4(a) shows the temperature dependence of the magnetization for an LCMO sample synthesized via standard solidstate reaction (bulk), nanocrystalline powder, synthesized via auto-combustion, and an epitaxial LCMO thin film (~100 nm thin) grown on SrTiO3 substrate via DC magnetron sputtering. From this plot, it can readily be seen that the sharp PM-FM transition occurring in the bulk LCMO is broadened in the samples with lower dimensionality. It can also be seen that the TC, estimated once again from the minima in the derivative of magnetization with temperature, shifts from 277 K in the bulk to 263 K in the nanometric powder and 261 K in the thin film. Reduced values of TC and Ms have been explained in connection with size reduction of manganite nanoparticles [52, 53]. Experimental data agree well with models predicting a core/shell morphology of the nanoparticles [54]. Such models assume a shell consisting of a magnetically dead or spin glass-like surface layer, which reduces the magnetization. From the point of view of changes in the internal structure, it is possible that nanometric-sized particles can decrease the Mn-O-Mn bond angle and increase its length, leading to a decrease in the transfer integral and consequently to a reduction in the TC [55]. In turn, finite size effects (cutoff of the correlation length) in the magnetic core of the particle contribute to the reduction of the TC. In the present paper, the auto-combustion method was utilized as a fast and energy-saving method for producing nanocrystalline LCMO material. The broadening phenomenon has been attributed to the distribution of the TC owing to polydispersity in the nanoparticle system [56]. In this regard, it has been established that in a percolative system, the distribution in the transition temperature is Gaussian, with a width that can be qualitatively linked to the disorder present [57]. From the Gaussian fit to the dM/dT curves in the lower inset of Fig. 4(a), distribution widths (*) of 12 K, 48 K, and 25 K were obtained for the bulk LCMO, the

11

nanocrystalline powder, and the thin film, respectively. This finding suggests that the local variation in the physical environment in the film, induced by particle size distribution in a sub-100 nm ensemble, is considerably increased in a sub-100 nm ensemble [4].The reduction in the average value of the TC has been previously observed in both the nanocrystalline and the thin film forms of LCMO and is usually attributed to finite size effects and disorder [52]. Figure 4(b) shows the dependence of the inverse DC magnetic susceptibility, F-1(T), on the temperature under H=1000 Oe for the LCMO samples discussed in Fig. 1(a). The

F-1(T) dependence of the all the samples varies linearly with T in the high-

temperature range, indicating Curie–Weiss behavior [F-1(T) = (T-Tp)/C]. No definite downturn in the magnetic susceptibility after the TC can be observed, which would be an indication of the presence of a Griffiths phase [58]. A Griffiths phase implies the presence of short-range FM correlation well above the TC. This phenomena is due to the increase in the magnetic moment arising from the growth of FM spin clusters in the PM region in the temperature range TC
decrease in the value of the TC. The experimental effective moment Ɋୣ୤୤ can be evaluated using the following relation:  ൌ paramagnetic

moment

୒ఽ ଷ୩ా

Ɋ௧௛௘௢ ୣ୤୤

is

12

௘௫௣ ଶ

൫ߤ௘௙௙ ൯ . In turn, the theoretical effective expressed

as

௧௛௘௢ ߤ௘௙௙ ൌ

ୣ୶୮

ଶ ሺ‫݊ܯ‬ଷା ሻ ଶ ሺ‫ ܯ‬ସା ሻ ൅ ሺͲǤ͵ െ ‫ݔ‬ሻߤ௘௙௙ [60]. The results for Ɋୣ୤୤ and Ɋ௧௛௘௢ ටͲǤ͹ߤ௘௙௙ ୣ୤୤ are listed in

ୣ୶୮

Table 2. It can be seen from the results summarized in Table 2 that Ɋୣ୤୤ increases upon Ni2+ doping. This unusual behavior is not in a manner consistent with the behavior of the TC. Nevertheless, this behavior can be explained by invoking the difference in the limits of strength of the FM-DE interactions between Mn3+ and Mn4+ [61]. In this regard, it has been demonstrated that low Ni-doping levels at Mn sites can induce drastic changes in the physical properties of manganites, due not only to the effect on the ratio of Mn3+/Mn4+ but to the capability of Ni ions to form ferromagnetic regions in a non-FM insulating matrix [36, 62]. The unclear evidence of the presence of a Griffiths phase in the F-1(T) curves would suggest that this scenario is not suitable for interpreting the increased effective paramagnetic moment upon Ni2+ doping. On the other hand, the presence of a mixture of Zener pairs (Mn3+–Mn4+) with S = 7/2 and non-paired Mn3+ ions with S = 2 in the paramagnetic state has been discussed in ୣ୶୮

several papers [63, 64]. As a consequence of this, the magnetic moment Ɋୣ୤୤ and the Curie constant of the mixture will be increased. Without doubt, further work is necessary in order to obtain a deeper insight into this complex issue. Note, however, that the value of the experimental moment of the pristine LCMO sample is in good agreement with the theoretical one. Hence the effect of Ni2+ doping on the magnetic properties of LCMO becomes evident. In order to evaluate the magnetic entropy change in the studied system, isothermal magnetization versus field, M(H), curves were recorded around the transition temperature in each sample. Figure 5 (left panels) shows the dependence of the magnetization on the applied magnetic field, recorded at different temperatures, from 0 to 1.5 T, for LCMO samples with x = 0, 0.2, 0.7, and 0.1. Below the TC, the

13

magnetization M increases sharply with the applied magnetic field up to 0.25 T and then grows smoothly towards saturation. Above the TC, the magnetization M increases more smoothly, as is typical in paramagnetic materials. This decrease is mainly due to the thermal agitation, which tends to disorder the magnetic moments. This variation indicates that there is a large magnetic entropy change associated with the FM–PM transition temperature occurring at the TC. The type of magnetic phase transition in La0.7Ca0.3Mn1-xNixO3 can be determined from the M2 vs. H/M curves (converted from the isothermal M–H data) using the Banerjee criterion [65]. The results of this analysis are presented in Fig. 5 (right panels). According to the Banerjee criterion, the magnetic transition is second order if all the M2 vs. H/M curves have a positive slope [66]. On the other hand, if some of the M2 vs. H/M curves show a negative slope at some point, the transition is first order [67]. It can readily be seen in Fig. 5 that the M2 vs. H/M curves of all the samples exhibit a positive slope at T>TC, indicating that the magnetic transition is second order for the analyzed samples. Here, it should be pointed out that most mixed-valence manganites such as La0.7Ca0.3MnO3 are associated with the firstorder character of the magnetic transition [68]. Nevertheless, changes of the magnetic phase transition order can be induced when the particle diameter is brought down from bulk to the few tens of nanometers range [69, 70]. In this context, the change of the magnetic phase transition from first-order may be attributed to surface pressure effects. Certainly it has been confirmed that the surface pressure increases in the nanoparticles [71]. The increased surface pressure as the particle size diminishes leads to a reduction in the cell volume of nanometric LCMO. By considering a spherical shape of the particles, the surface pressure is given by the relation P=2S/d, where S and d correspond to the surface tension and diameter of the particle, respectively. By taking a bulk modulus B|150 GPa (typical for many oxides) [72] and surface tension

14

S|90 N/m, the pressure necessary to produce a change in the cell volume of a4% amounts to a6 GPa. It should also be pointed out that there are a large number of spins on the surface of the nanoparticles that are generally expected to be disordered and that will lead to the destruction of any spin order [53]. This core/shell morphology model (previously mentioned) predicts that the disordered outer layer is more likely to undergo a second-order transition, from the disordered state to the paramagnetic one [73]. Therefore, it is very probable that the observed second-order magnetic phase transition in the parent LCMO sample is linked to the effects of the particle size reduction. Magnetic refrigeration technique depends on the magnetic entropy change (-'SM) of magnetic material upon magnetic field application/removal. Concretely, magnetic refrigeration is based on the MCE [74-77]. This effect is based on the decrease of the spin entropy of a magnetic material upon application of an external magnetic field. The reduction in the magnetic entropy is then compensated by an increase in the lattice entropy, resulting in an increase in the temperature of the sample. In turn, when the magnetic field is removed adiabatically, the magnetic spins tend to randomize, which leads to an increase in the magnetic entropy and a decrease in the lattice entropy. Hence the temperature of the sample diminishes. In order to evaluate -'SM, it is necessary to make a numerical approximation for the integral written in section 2. From the respective isothermal magnetization data, 'SM is approximately given by 'SM(T,'H) = σ

ெ೔ ିெ೔శభ ்೔ ି்೔శభ

ο‫[ ܪ‬78]. In the latter expression, Mi1 and Mi are the

magnetization values measured at Ti1 and Ti temperatures at a magnetic field change 'H. The calculated -'SM values for x = 0, 0.2, 0.07, and 0.1 compositions at three different magnetic fields are plotted in Fig. 6. The magnetocaloric properties of LCMO

15

samples with x = 0, 0.2, 0.7, and 0.1 under applied magnetic fields of 0.5, 1, and 1.5 T are listed in Table 3. It is evident that -'SM increases with an increasing magnetic field and the -'SM,max is shifted towards higher temperature for all the compositions. It is also clear that the largest -'SM occurs near the TC, which is a property of simple ferromagnets, due to the efficient ordering of magnetic moments induced by the magnetic field at the ordering temperature [79]. The maximum magnetic entropy changes observed for the x = 0, 0.02, 0.07, and 0.1 compositions are 0.85, 0.77, 0.63, and 0.59 J/kg K, respectively, under a magnetic field of 1.5 T. The value of -'SM for the undoped sample is in good agreement with that reported in Ref. [4], in which LCMO particles with a mean diameter of ~15 nm were obtained by the authors using the solgel method. Nevertheless, it should be mentioned that in nanocrystalline samples, the sensitive dependence of the magnetocaloric effect on synthesis details becomes even more evident. Thus differences in the -'SM values of more than an order of magnitude can be found in the literature for similarly-sized particles subject to a comparable field change [4, 80]. It can also be noted from Fig. 6 that Ni2+ doping not only shifts the TC value towards lower temperatures but also reduces the 'SM,max by a value as low as a30% for x=0.1. Probably the change in the -'SM,max values upon Ni2+ doping is due to the relatively low strength of the magnetic fields, which is insufficient to achieve a complete saturation of the magnetization in the studied samples. Similar changes in -

'SM,max have been reported for Y-doped manganites and were attributed to the low strength of the magnetic field used to measure the MCE [81]. On the other hand, -

'SM,max in manganites remained constant when the samples were measured in magnetic fields t 3 T) [19, 82]. Thus the results presented in Fig. 6 concerning the 'SM,max variation upon Ni2+ doping agree with the results reported in Ref [81]. A small variation of the-'SM,max upon doping can be considered as a plus for magnetic 16

refrigeration applications, because the same material can be used within a wide range of temperatures. In general, the magnetic entropy changes achieved for Ni2+-doped LCMO samples synthesized via auto-combustion (as in the present paper) are higher than those reported for other manganites, e.g. La0.7Sr0.3Mn1-xNixO3, with comparable Ni2+-doping levels but synthesized via standard solid-state reaction [19]. Generally, an important criterion for selecting magnetic refrigerants is the cooling power per unit volume, namely, the relative cooling power RCP [83], which corresponds to the amount of heat transferred between the cold and the hot sinks in the ideal refrigeration cycle. The RCP parameter can be calculated through the ்

relationെ ‫ ்׬‬మ οܵெ ሺܶሻ ݀ܶ, where T1 and T2 are the temperatures defining the full width భ

at half maximum (FWHM). For practical purposes, the RCP value can be evaluated using the equivalent relation RCP=~'SM,max~xGTFWHM. From the RCP calculations, summarized in Table 3, it can be seen that the addition of Ni2+ increases the RCP value compared to that of the parent LCMO. This is certainly desirable for potential applications of these materials in magnetic refrigeration. In this regard, the sample with a Ni-doping level as low as 2% features the highest RCP value (~60 J/kg) around 230 K at 1.5 T applied magnetic field. It has been observed on a number of occasions that the reduced maximum value of -'SM, which often accompanies broad magnetic entropy change peaks, can be compensated for by the increased width, resulting in an enhanced RCP over sharper transitions [84]. Table 3 shows that this scenario holds true for all the La0.7Ca0.3Mn1xNixO3

(x = 0, 0.02, 0.07, 01) samples. In the case of thin films, a strong drop in the -

'SM value has been reported [4]. Nevertheless, the strong drop in the -'SM value was

17

compensated by the increased breadth of the transition. Thus large RCP values are often encountered in materials in thin film form. The results achieved in the present study and those reported in several publications show that the MCE in perovskite manganites is certainly large. Probably the strong spin–lattice coupling, related to the magnetic ordering process, is the cause of the large MCE found in these challenging oxides [67, 85]. Indeed, it is known that the magnetic transitions in these manganites are concomitant with lattice changes, which is taken as evidence of strong spin-lattice coupling in these materials. Generally speaking, the influence of structural changes on the magnetism and the MCE in colossal magnetoresistance (CMR) manganites is primarily related to the effective one-electron bandwidth W [86], which is controlled by the A-site doping [87]. Since the La0.7Ca0.3MnO3 manganites feature a narrow W, other effects, such as collective Jahn–Teller (JT) distortions and antiferromagnetic interactions, can coexist and strongly compete with the ferromagnetic phase [88]. This finding is particularly relevant for the explanation of features of the metal-insulator and CMR observed in these manganites. Indeed, the properties of the metallic ferromagnetic state in doped manganites with relatively large W, such as La0.7Sr0.3MnO3 [89], can be satisfactorily described by the DE theory. This is not the case for the La0.7Ca0.3MnO3 counterpart. Here, it should be noted that cooperative JT distortions are present in the orthorhombic structure of La0.7Ca0.3MnO3 but not in the rhombohedral structure of La0.7Sr0.3MnO3, due to the higher symmetry of the MnO6 octahedra in this phase. Thus, the Mn-O-Mn bond angle decreases (or W decreases) with Ni-doping (Table 1), reducing the DE interaction [79]. By using the tight-binding approximation, W for ABO3-type perovskites ଷǤହ , where Θ = (π/2-Mn–O–Mn!) [90]. Hence has been determined as WvcosΘ/݀ெ௡ିை

the decrease in bandwidth W reduces the overlap between the O2p and the Mn3d

18

orbitals, which in turn decreases the exchange coupling of Mn3+–Mn4+, resulting in a decrease in the magnetic ordering [91]. 4. Summary and conclusions

La0.7Ca0.3Mn1-xNixO3 (x = 0, 00.2, 0.07, and 0.1) manganites were synthesized via the auto-combustion method, and their structural, magnetic, and magnetocaloric properties were carefully studied. It was demonstrated that auto-combustion is a fast, effective and low-cost synthesis route for obtaining high-quality powders of novel oxide materials. All the cell parameters and the unit cell volumes of the LCMO samples increased with an increase in the Ni2+ content, due to the larger ionic size of these ions compared with those of Mn3+ and Mn4+. It was verified that both the magnetic and magnetocaloric properties of the parent LCMO can be controlled through substitution of Mn3+ with Ni2+ ions. The TC of LCMO shifted to lower temperatures, probably due to the decrease in the ferromagnetic DE interactions. The observed second-order magnetic phase transition in the parent LCMO sample could be attributed to the effects of the particle size reduction. The inverse DC magnetic susceptibility data fit well to the Curie–Weiss law, indicating typical paramagnetic behavior above the TC. The effective

paramagnetic

moments

were

evaluated

both

experimentally

and

theoretically. The experimental moment increased upon Ni2+ addition, which is not in a manner consistent with the behavior of the TC. The -'Smax shifted towards lower temperatures with a small change in the amplitude upon Ni2+ addition. Interestingly, it was verified that the RCP values of the Ni2+-doped samples increased compared with that of the parent LCMO. The highest RCP value (~60 J/kg) was observed for the sample with a Ni2+-doping level of 2% around 230 K at 1.5 T applied magnetic field. The achieved results suggest that the studied Ni2+-doped LCMO manganites are

19

materials with potential to be used in magnetic refrigeration applications over a wide range of operation temperatures. Acknowledgments

This paper was supported by the Universidad Nacional de Colombia, Sede Medellín. Fruitful discussions with Dr. P.T. Phong at Khanh Hoa University, and Dr. A. Hussein at Sohag University are warmly acknowledged. One of the authors (AZ) acknowledges the financial support of the Departamento Administrativo de Ciencia, Tecnología e Innovación (Colciencias). References [1] Y. Tokura, N. Nagaosa, Science 462 (2000) 288. [2] E. Dagotto, T. Hotta, A. Moreo, Phys. Rep. 344 (2001) 1. [3] J. Hemberger, M. Brando, R. Wehn, V. Yu. Ivanov, A.A. Mukhin, A.M. Balbashov, A. Loidl, Phys. Rev. B 69 (2004) 064418. [4] P. Lampen, N. S. Bingham, M. H. Phan, H. Kim, M. Osofsky, A. Piqué, T. L. Phan, S. C. Yu, H. Srikanth, App. Phys. Lett. 102 (2013) 062414. [5] I. Hussain, M.S. Anwar, E. Kim, B.H. Koo, C.G. Lee, Korean J. Mater. Res. 26 (2016) 623. [6] L.W. Zhang, G. Feng, H. Liang, B.S. Cao, Z. Meihong, Y.G. Zhao, J. Magn. Magn. Mater. 219 (2000) 236. [7] T-Long Phan, P. Zhang, T. D. Thanh, S. C. Yu, J. Appl. Phys. 115 (2014) 17A912. [8] B. Chen, C. Uher, D. T. Morelli, J. V. Mantese, A. M. Mance, A. L. Micheli, Phys. Rev. B 53 (1996) 5094. [9] S. Raj, H.C. Padhi, P. Raychaudhury, A.K. Nigam, R. Pinto, M. Polasik, F. Pawlowski, D.K. Basa, Nucl. Instr. And Meth. In Phys. Res. B 174 (2001) 344. [10] D.N.H. Nam, N.V. Dai, L.V. Hong, N.X. Phuc, S.C. Yu, M. Tachibana, E. Takayama-Muromachi, J. Appl. Phys. 103 (2008) 043905. [11] N. A. de Olivira, P. J. Von Ranke, Phys. Rep. 489 (2010) 89. [12] M.-H. Phan, S.-C. Yu, J. Magn. Magn. Mater. 308 (2007) 325. [13] A.M. Tishin, Y.I. Spichkin, IOP Publishing Ltd, Bristol, 2003. [14] A. Marzouki-Ajmi, M. Mansouri, W. Cheikhrouhou-Koubaa, M. Koubaa, A. Cheikhrouhou, J. Magn. Magn. Mater. 433 (2017) 209.

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24

Figure captions Figure 1. X-ray diffraction patterns recorded at 300 K for La0.7Ca0.3Mn1-xNixO3 (x = 0.1, 0.02, 0.07, and 0.1) samples synthesized via the auto-combustion method. The refined profiles as well as the difference profile are shown at the bottom. Figure 2. Variation of the cell parameters of La0.7Ca0.3Mn1-xNixO3 samples with the value of x. Figure. 3. FE-SEM image of the pristine LCMO sample. Figure 4. (a) Temperature dependence of the field-cooled magnetization recorded in a 1000 Oe magnetic field for La0.7Ca0.3Mn1-xNixO3 samples. Upper inset: dM/dT versus T curves for the samples plotted in the main panel. Lower inset: M(T) dependence for an LCMO sample synthesized via standard solid-state reaction (bulk), nanocrystalline powder, synthesized via auto-combustion, and an epitaxial LCMO thin film (~100 nm thin) grown on SrTiO3 substrate via DC magnetron sputtering. (b) F -1(T) dependence for La0.7Ca0.3Mn1-xNixO3 samples at H=1000 0Oe. The straight lines correspond to the Curie–Weiss law fit to the susceptibility data above TC. Figure 5. Left panels: Magnetization versus magnetic field M(H) recorded around TC for La0.7Ca0.3Mn1-xNixO3 samples. Right panels: Arrott plots isotherms of M2 versus H/M at different temperatures for La0.7Ca0.3Mn1-xNixO3 samples. Figure 6. Temperature dependence of the magnetic entropy change for La0.7Ca0.3Mn1xNixO3

(x = 0.1, 0.02, 0.07, and 0.1) samples under different magnetic fields.

25

Table 1. Lattice parameters, atomic positions and agreement parameters of La0.7Ca0.3Mn1-xNixO3 obtained by Rietveld refinement of XRD data at room temperature. Space group corresponds to Pnma (#62), La, Ca and O1 occupy 4c Wyckoff position (x¼ z), Mn/Ni occupy 4a (0 0 0) and O2 occupy 8d (x y z). La0.7Ca0.3Mn1-xNixO3 x

0.00

0.02

0.07

0.10

a

5.4459(18) 5.4550(19) 5.4556(30)

5.4555(32)

b

7.6971(26) 7.7099(28) 7.7103(41)

7.7227(55)

5.4724(24) 5.4850(24) 5.4906(37)

5.4880(46)

V (Å )

229.4(2)

230.7(3)

230.9(3)

231.2(3)

RF (%)

7.24

5.76

8.80

8.78

F2

1.038

1.103

1.399

1.182

0.4853(8)

0.4903(12)

c 3

La/Ca

X

0.4913(14) 0.4930(9)

Z

0.0014(11) 0.0009(22) 0.0001(43)

0.0001(37)

X

0.4770(52) 0.4781(12) 0.4981(64)

0.4757(43)

Z

0.5469(47) 0.5487(63) 0.5722(83)

0.5780(89)

X

0.2220(40) 0.2146(35) 0.2285(91)

0.2330(69)

Y

0.0201(38) 0.0171(44) 0.0290(40)

0.0236(61)

Z

0.2867(51) 0.2806(65) 0.2686(57)

0.2659(93)

Mn-O1-Mn angle (°)

163.121(8) 162.718(9) 156.700(26)

153.853(27)

Mn-O2-Mn angle (°)

162.673(9) 163.083(4) 163.964(5)

162.280(15)

crystallite size (nm)

24.60

20.35

O1

O2

22.93

26

21.19

Table 2. Magnetic data for La0.7Ca0.3Mn1-xNixO3 samples with x varying between 0 and ௧௛௘௢ refer to the Curie temperature, the Curie constant, the 0.1. TC, C, Tp, ߤ௘௙௙ and ߤ௘௙௙ ௘௫௣

Weiss temperature, the experimental and theoretical effective paramagnetic moment, respectively.

0.00 0.02 x TC [K] 264 225 C [emuK/Oe.g] 0.0122 0.0164 276 250 Tp [K] ௘௫௣ 4.51 5.23 ߤ௘௙௙ [PB] ௧௛௘௢ 4.61 4.58 ߤ௘௙௙ [PB]

0.07 182 0.0218 205 6.03 4.50

0.10 174 0.0216 198 6.01 4.45

27

Table 3. Summary of magnetocaloric properties of La0.7Ca0.3Mn1-xNixO3 samples at three different magnetic field strengths.

Ni-doping level

H [T]

-∆SM [J/kg*K]

δTFWHM [K]

RCP [J/kg]

0.5

0.329

37.51

12.34

1

0.612

37.82

23.15

1.5

0.858

37.70

32.34

0.5

0.289

77.28

22.34

1

0.544

77.27

42.04

1.5

0.771

77.31

59.60

0.5

0.253

83.28

21.07

1

0.458

83.27

38.14

1.5

0.633

83.27

52.71

0.5

0.236

84.51

19.94

1

0.428

84.51

36.17

1.5

0.596

84.51

50.37

(x)

0

0.02

0.07

0.1

28

(402)

(242)

(042)

(202)

(220)

(200)

(101)

x=0.00

Intensity [a. u.]

x=0.02

x=0.05

x=0.07

x=0.10

20 Fig. 1. Morán

Fig. 1

40

60 2T [°]

80

Fig. 2

c

231

3

5.48

V [Å ]

Lattice parameter [Å ] Fig. 2. Morán

232

5.50

b/—2

230

5.46 a

5.44 0.00

0.05 Doping [x]

0.10

229

Fig. 3. Morán

Fig. 3

0

Bulk Nanometric Thin film

100 200 300 T [K] H=1000 Oe

100

Fig. 4. Morán

200 T [K]

4

x=0 x=0.02 x=0.07 x=0.1

(b)

2

-1

H=500 Oe

3

100 200 300 T [K]

1

F [x10 Oeg/emu]

25

0

Fig. 4

(a)

dM/dT

x=0 x=0.02 x=0.07 x=0.1

M/M (50 K)

M [emu/g]

50

300

0 100

200 T [K]

300

x=0.02

M [emu/g]

40

20

0 0,0

M [emu/g]

60

1,0 H [T]

1,5

x=0.07

40 20 0 0,0

Fig. 5

0,5

0,5

1,0 H [T]

1,5

2 3 2

0

1,5

150 K 160 K 170 K 180 K 190 K 200 K 205 K 210 K 215 K 220 K 225 K 230 K 235 K 240 K 245 K 250 K 255 K 260 K 265 K 270 K 275 K

50 K 100 K 110 K 120 K 130 K 140 K 150 K 155 K 160 K 165 K 170 K 175 K 180 K 185 K 190 K 195 K 200 K 205 K 210 K 215 K 220 K 230 K 240 K 300 K

2 3 M x10 [emu/g]2

0,5 1,0 H [T]

1

0

1 2 H/M [x10 Oe.g/emu]

2

1

0

0

150 K 160 K 170 K 180 K 190 K 200 K 205 K 210 K 215 K 220 K 225 K 230 K 235 K 240 K 245 K 250 K 255 K 260 K 265 K 270 K

1 3 2 3 H/M [x10 Oe.g/emu]

x=0.07

3 2 1 0

200 K 205 K 210 K 220 K 225 K 230 K 235 K 240 K 245 K 250 K 255 K 260 K 265 K 270 K 275 K 280 K

3

x=0.02

2 M x10 [emu/g]

0 0,0

x=0

3

20

2

2

M [emu/g]

40 x=0

M x10 [ emu/g]

200 K 205K 210 K 220 K 225 K 230 K 235 K 240 K 245 K 250 K 255 K 260 K 265 K 270 K 275 K 280 K 290 K

0

13 2 H/M [x10 Oe.g/emu]

50 K 100 K 110 K 120 K 130 K 140 K 150 K 155 K 160 K 165 K 170 K 175 K 180 K 185 K 190 K 195 K 200 K 205 K 210 K 215 K 220 K 230 K 240 K

20 0 0,0

0,5

Fig. 5. Morán

Fig. 5

1,0 H [T]

1,5

2

x=0.1

3

2 1

2

40

3 M x10 [ emu/g]

M [emu/g]

x=0.1

50 K 100 K 110 K 120 K 130 K 140 K 145 K 150 K 155 K 160 K 165 K 170 K 175 K 180 K 185 K 190 K 195 K 200 K 205 K 210 K 215 K 220 K 225 K

0

0

1 H/M [x10 Oe.g/emu] 3

50 K 100 K 110 K 120 K 130 K 140 K 145 K 150 K 155 K 160 K 165 K 170 K 175 K 180 K 185 K 190 K 195 K 200 K 205 K 210 K

1,0

x=0

0,5

250 T [K]

1.0

x=0.02

x=0.07

0,0 150

1,0

200 T [K]

0.5 T 1T 1.5 T

250

x=0.1

- 'SM[J/kgK]

- 'SM[J/kgK]

0.5 T 1T 1.5 T

300

0.5

0,5

100

Fig. 6. Morán

Fig. 6

0.5 T 1T 1.5 T

0,5

0,0

0.0

1,0 - 'SM[J/kgK]

- 'SM[J/kgK]

0.5 T 1T 1.5 T

T [K]

200

0,0

100

T [K]

200