Pb(Fe0.5Nb0.5)O3–BiFeO3-based multicalorics with room-temperature ferroic anomalies

The search for new single-phase multicaloric materials, combining electrocaloric and magnetocaloric effects, is just at its beginning.


I. Introduction
Multiferroics are a class of materials, where in general more than one ferroic (including relaxor), antiferroic or ferric orders coexist. 1,2 A lot of attention is paid to the so-called singlephase magnetoelectric multiferroics, combining ferroelectric and ferromagnetic properties in the same material, which are attractive due to the different origins and opposite mechanisms of both ferroic orders. [2][3][4][5] Magnetoelectric multiferroics are promising candidates for future environmental friendly solid-state cooling based on caloric effect, [6][7][8][9] which is referred to reversible adiabatic temperature changes of the material under the influence of external stimuli, such as electric or magnetic fields. [9][10][11] achieved by substitution of A-site of the crystal lattice with rare earth elements, such as Gd ions. [36][37][38][39] Rare earth elements possess strong localized magnetic moments that arise from partially filled 4f inner electron shell, making them promising candidates for enhancing the magnetic properties of the materials. 40 In previous Gd-doped BFO studies, [36][37][38][39] it was shown that the addition of 12 mol% of Gd ions enhance the ferromagnetic properties and at the same time does not increase electrical conductivity and does not lead to the formation of non-polar phase. 37,41 However, no studies on doping of PFN-BFO have been reported, therefore in this work a similar approach as reported for pure PFN and BFO, namely doping with 0.5 mol% of Mn and 12 mol% of Gd ions, will be employed.
In this work, seven different multiferroic (1-x)Pb(Fe 0.5 Nb 0.5 )O 3 -xBiFeO 3 (x = 0-0.5) solid solutions were prepared. Among them, PFN-20BFO exhibits both dielectric permittivity and magnetic susceptibility anomalies close to RT and is therefore a promising candidate for multicalorics. Further doping of PFN-20BFO with 0.5 mol% of Mn and 12 mol% of Gd ions suppress the electrical conductivity and enhance the magnetic ordering leading to improvement of the EC and MC properties. The results of this study show the importance of compositioncaloric-properties inter-correlations in PFN-BFO-based ceramics.
The composition with 0.5 mol% of Mn (substitution for B-site cations, Fe 3+ and Nb 5+ ) and 12 mol% of Gd (substitution for Bi 3+ ) was prepared according to the nominal formula 0.8Pb(Fe 0.5 Nb 0.5 ) 0.995 Mn 0.005 O 3 -0.2Bi 0.88 Gd 0.12 Fe 0.995 Mn 0.005 O 3 . Due to the problems with aggregation and dispersion of Gd 2 O 3 , for the preparation of Gd and Mn co-doped PFN-20BFO, in all starting oxides, surfactant ammonium polyacrylate (PAA) was used. The PAA, used as a steric stabilizer, was formed by mixing poly(acrylic acid) (50 wt% solution in H 2 O, Sigma Aldrich) and 1-butlyamine (99%, Alfa Aesar) in absolute ethanol. 43,44 Gd 2 O 3 (99.99%, Metall) was pre-milled in a planetary ball mill (Retsch PM400, Retsch, Haan, Germany) at 150 min −1 for 4 h in ethanol/PAA mixture using YSZ milling balls (ball diameters 3 mm). PAA was added also to all other pre-milled starting oxides and the mixtures were homogenized by milling in planetary ball mill at 150 min −1 for 2 h using YSZ milling balls (ball diameters 3 mm). A stoichiometric mixture of all starting oxides (40 g) were homogenized in a planetary ball mill at 150 min −1 for 4 h in absolute ethanol using YSZ milling balls (ball diameters 3 mm). The homogenized mixture was dried and heat-treated at 350 °C for 30 min (heating and cooling rates of 2 °C·min −1 ) in order to remove PAA from the powder. The powder mixture was mechanochemically activated in the same planetary ball mill for 14 h at 400 min −1 in a tungsten carbide milling vial (volume 80 cm 3 ), filled with 7 tungsten carbide milling balls (ball diameters 15 mm). After activation, the synthesized powder, i.e., mixtures of crystalline perovskite and amorphous phases, was re-milled in a planetary ball mill in absolute ethanol, at 200 min −1 , for 4 h, using YSZ milling balls (ball diameters 3 mm).
All prepared powders were uniaxially pressed (50 MPa) into pellets of a diameter of 8 mm followed by isostatic pressing at 300 MPa. The powder compacts were reactively sintered in double alumina crucibles in the presence of PFN packing powder. Sintering temperatures for each composition were estimated from sintering curves (details in Ref. [42]). To maintain similar grain size, PFN-100xBFO samples with lower content of BFO (x ≤ 0.1) were sintered at 950 °C for 2 h, while samples with higher content of BFO (x ≥ 0.2) were sintered at 900 °C for 2 h. Similar thermal treatment was used also for Mn-doped PFN-20BFO (900 °C for 2 h), while in Mn,Gd co-doped sample, due to the different sintering behavior, higher sintering temperature (1000 °C for 2 h) was used (shown in Supplementary material, A). Sintering runs were performed in oxygen atmosphere as suggested in Ref. [45] and used in our previous work. 23,42 The use of saturated oxygen atmosphere during the sintering prevents (or minimizes) the reduction of Fe 3+ to Fe 2+ and reduces the amount of oxygen vacancies, which consequently resulted in a lower electrical conductivity and dielectric losses of the material. 45 For all samples, heating and cooling rates of 2 °C·min −1 were used.

Characterization
The densities of the sintered pellets were determined using a gas-displacement density analyzer (Micromeritics, AcuuPyc II 1340 Pycnometer) at RT. For the calculation of the relative density, the crystallographic densities were obtained from Rietveld refinement analysis (details in Supplementary material, B). X-ray powder diffraction (XRD) patterns of crushed ceramic pellets were recorded on a X'Pert Pro (PANalytical B.V., Netherlands) highresolution powder diffractometer equipped with Cu Kα 1 radiation (45 kV, 40 mA), in reflection mode. Diffraction patterns were recorded in the 10-70° 2θ-range with a step size of 0.017° and an integration time of 100 s per step at RT. The phase identification was done with PDF-4+ database.
The microstructure of sintered ceramics was examined using a field-emission scanning electron microscope (FE-SEM, JSM-7600F, Jeol Ltd., Japan) equipped with an energy dispersive X-ray spectrometer (EDXS, Inca Oxford 350 EDS SSD, Oxford Instruments, Abingdon, U.K.). Prior to microstructural analysis, the samples were fractured for fracture-surface examination, ground and fine-polished using a colloidal silica suspension, followed by thermal etching at 775 °C for 15 minutes in the presence of PFN packing powder. The average grain size and grain-size distribution were evaluated from the digitized images of etched surfaces, processed using the Image Tool software (UTHSCSA Image Tool Version 3.00. 2002) 46 by measuring more than 500 grains. The grain size is expressed as the Feret's diameter. 47 For characterization of electrical properties, the ceramic pellets were cut, thinned and polished to a thickness of ~200 µm. In order to release the internal stresses produced during the mechanical processing, the thinned ceramic pellets were heated up to 600 °C for 1 h (with 5 °C·min −1 heating rate) and then slowly cooled down (1 °C·min −1 ) to RT. The surfaces of the pellets were coated with Au electrodes (diameter of 5 mm) using RF-magnetron sputtering machine (5Pascal, Italy).
The relative dielectric permittivity (ε) and dielectric loss (tan δ) as a function of temperature were analyzed during cooling from 200 °C to −50 °C using an Aixacct TF analyzer 2000 (aixACCT Systems GmbH, Aachen, Germany) coupled with a HP 4284A Precision LCR impedance meter (Hewlett-Packard, California, USA) in a frequency range of 0.1-100 kHz. DC current-density-electric-field (j-E) response was measured using a Keithley 237 (Keithley Instruments, Ohio, USA) high voltage source measure unit. For the j-E measurements, the samples were exposed to step-like voltages within the range of ±0.5 kV·cm −1 . In this range, from negative to positive polarity, 13 equal field steps were applied. Upon each of these steps, the voltage was maintained constant for 1.5 h for the current to stabilize, before being recorded. The specific DC electrical conductivity (σ DC ) was obtained from the slope assuming Ohm's law. AC current density-electric field (j-E) loops were measured at RT using an Aixacct TF analyzer 2000 by applying a single sinusoidal waveform with the frequency of 1 Hz and electric fields up to 80 kV·cm −1 . Polarizationelectric field (P-E) loops were automatically obtained by software using numerical integration of current signal.
The EC effect was measured using the direct method. The measurements were performed on ~80 µm-thick ceramic samples with sputtered Au electrodes. The direct EC measurements were performed in a modified differential scanning calorimeter (Netzsch DSC 204 F1, Selb, Germany), which allowed a precise temperature stabilization. The temperature changes were measured with a small-bead thermistor glued to the sample; the temperature was measured using a digital multimeter. The EC temperature change (ΔT EC ) values presented in this work are the measured maximum temperature changes multiplied by a correction factor of 1.60-1.75, which accounts for the passive part of the sample and also the thermal losses in the gold wires, the silver paste and the thermistor. The used correction factor agrees well with the results of the numerical simulation. 48 The details of the method can be found in Ref. [49].
Field-cooled (FC) DC magnetic moment vs. temperature was measured between −270 °C and 127 °C at different magnetic fields (0.02-90 kOe) using a physical property measurement system vibrating sample magnetometer (PPMS-VSM, Quantum Design, California, USA). Furthermore, the magnetization-magnetic field (M-H) hysteresis loops were obtained in magnetic fields up to 90 kOe at temperatures of −270 °C and RT. For the measurements rod-like-shaped ceramic samples were used. The M values were obtained by normalizing the magnetic moment values by the mass of the samples. The MC effect was determined by the indirect method [50], where the MC temperature change (ΔT MC ) was calculated from the magnetization versus temperature measurements at different magnetic fields (0.5-90 kOe).

A. (1-x)Pb(Fe 0.5 Nb 0.5 )O 3 -xBiFeO 3
In the first part, multiferroic properties of PFN-100xBFO (x = 0-0.5) solid solutions were investigated. All prepared ceramics were single-phase, homogeneous and highly dense (> 97.6% of the crystallographic density), as previously reported in Ref. [42]. The temperature dependences of the real component of the relative permittivity ε′ for all compositions measured at 1 kHz are shown in Fig. 1a. The peak-permittivity of PFN-100xBFO is shifting towards RT with increasing amount of BFO (x ≤ 0.2) and it reaches the lowest value in PFN-20BFO, i.e., 33 °C (Fig. 1a). This is accompanied with increasing relaxor behavior, as seen in Fig. 1b and reported in Ref. [42]. By further increasing the BFO content (x ≥ 0.3), the peak-permittivity temperature seemingly shifts to temperatures above the one of PFN (i.e., 113 °C), as visible from ε′ measurements at 100 kHz (inset of Fig. 1a). A reason for progressively increased relaxor behavior by the addition of BFO is most probably related to increased compositional disorder on the A-site of a perovskite structure, which arises when different multivalence ions occupy a crystalographically equivalent place in the crystal lattice. 42,51,52 Heterovalent substitution of the differently sized Pb 2+ (119 pm) 53 and Bi 3+ (103 pm) 53 ions appears to disturb a long-range ferroelectric order in PFN-rich PFN-BFO solid solutions as previously suggested for Pb-based relaxors. 54 In the next step, the influence of BFO addition on magnetic properties of PFN-100xBFO were studied. The M(T) curves of all compositions, measured at magnetic field of 1 kOe, are shown in Fig. 2a. In all the samples M increases with decreasing temperature reaching a maximum value of ~0.032 emu·g −1 at −270 °C in PFN. PFN exhibits the paramagnetic-toantiferromagnetic phase transition (T N ) at −120 °C. With increasing the amount of BFO, T N is almost linearly shifted towards higher temperatures reaching RT in PFN-20BFO (Fig. 2b). In PFN-50BFO T N is probably shifted beyond the measured temperature range (i.e., above 130 °C), which is in agreement with previously reported studies. 55 As demonstrated in Fig. 2b, increasing T N values with increasing amount of BFO in PFN-BFO solid solutions are directly related to the higher amount of Fe 3+ on the B-site perovskite lattice. These results agree well with the previously reported studies, where it was shown that in pure PFN 56 as well as in PFN-50BFO 57,58 the magnetic properties and thus T N depend on the number of Fe-O-Fe linkages in the structure. Therefore, in our case, substitution of similarly large Nb 5+ with Fe 3+ ions (~64 pm) 53 increase the Fe/Nb atomic ratio on the B-site sublattice, leading to higher temperature of paramagnetic-to-antiferromagnetic phase transition. On the other hand, a decrease in Nb 5+ concentration reduces the tendency for the formation of short-range spin-glass-like phase at low temperatures, which is reflected in lower M values in BFO-rich compared to PFN-rich PFN-BFO solid solutions (Fig. 2a). 25,59 In all compositions slim antiferromagnetic M-H hysteresis loops with almost zero remanent magnetization were obtained at −270 °C (Supplementary material, C), which indicates that in the PFN-BFO system increasing BFO content does not significantly affect the long-range antiferromagnetic ordering. Among all the compositions, PFN-20BFO exhibits the dielectric permittivity anomaly as well as paramagnetic-to-antiferromagnetic phase transition at RT, which makes this material a promising candidate for room-temperature multicaloric cooling. However, further evaluation of EC and MC properties revealed some issues, namely excessive electrical conductivity and poor magnetic response, that limit its cooling capability (Supplementary material, D). At higher temperatures and applied electric fields the increased electrical conductivity leads to the appearance of Joule heating (JH), which significantly reduces material's EC response reaching the maximum EC effective cooling (ΔT eff , defined as a difference between ΔT EC and ΔT JH ) of 0.50 °C at RT and 90 kV·cm while the ferromagnetic response can be enhanced by the addition of Gd ions. [36][37][38][39] Based on that, in the second part of this study the influence of Mn and Gd doping on multiferroic and multicaloric properties of PFN-20BFO is discussed.

B. Mn-doped and Gd,Mn co-doped PFN-20BFO
In the second part of this study, the effect of Mn and Gd addition on structural, microstructural and functional properties of PFO-20BFO was investigated. Firstly, the phase composition, crystal symmetry and microstructure of pure, Mn-doped and Gd,Mn co-doped PFN-20BFO were studied. The X-ray powder diffraction patterns of crushed ceramic pellets are shown in Fig. 3a. In all three samples only the peaks corresponding to the perovskite phase are present and no additional peaks corresponding to secondary phases are observed. Rietveld refinement analysis (Supplementary material, B) revealed that all three compositions exhibit pseudocubic symmetry at RT (Pm m space group), which is consistent with previous studies 3 of pure PFN-20BFO. 42,60,61 SEM images of polished sample surfaces are shown in Fig. 3b. All three samples are highly dense, i.e., equal or more than 97.6% of the crystallographic density, exhibiting homogeneous and uniform microstructures. In the microstructure of pure and Mn-doped PFN-20BFO no visible secondary phase are present, while in Gd,Mn co-doped sample small dark inclusions are observed. A qualitative EDXS analysis revealed that these dark phase regions correspond to Fe-rich and Gd,Nb-rich secondary phases (details in Supplementary material, F). The average grain size of pure and Mn-doped PFN-20BFO are in the range of 1 μm, while the grain size of Gd,Mn co-doped sample is larger (~4.8 μm) due to the higher sintering temperature. In all three samples monomodal grain size distribution, indicating the normal grain growth during the sintering, were observed (Supplementary material, G). The temperature and frequency (ν) dependence of the real component of relative permittivity ε′ and dielectric loss tan δ of all three PFN-20BFO-based compositions are shown in Fig. 4. Pure PFN-20BFO exhibits broad and diffusive peak-permittivity with the maximal value of ~7700 at 0.1 kHz and 27 °C indicating relaxor-like behavior (Fig. 4a). 42 The sample shows a large frequency dependence of peak-permittivity temperature (T ε,max ) that is spanning from 27 °C at 0.1 kHz to 53 °C at 100 kHz (i.e., ΔT ε,max of 26 °C). The addition of 0.5 mol% Mn resulted in slightly lower peak-permittivity values with T ε,max shifted to higher temperatures, reaching the maximum value of ~6200 at 0.1 kHz and 65 °C (Fig. 4b). While further addition of 12 mol% Gd shifts the T ε,max to slightly lower temperatures compared to pure PFN-20BFO with the maximum value of ~5800 at 0.1 kHz and 18 °C (Fig. 4c). Both modifications of PFN-20BFO lead to a weakening of the relaxor behavior, which is reflected in smaller ΔT ε,max , namely ~5 °C in both samples. In all three samples high room-temperature ε′ values are observed, ranging from ~5200 to ~7700 at 0.1 kHz.
The tan δ of all three compositions are frequency-dependent with relatively low values at RT, namely from 0.05 to 0.01 at 1 kHz (Fig. 4). Due to enhanced electrical conductivity in pure PFN-20BFO, above 120 °C the tan δ significantly increases. The addition of Mn ions successfully reduces the tan δ throughout entire temperature range showing very low values even at high temperatures, i.e., ~0.05 at 200 °C and 0.1 kHz (Fig. 4b). Further modification of material by Gd ions leads back to an increase in tan δ, but the values remain comparable or even slightly lower than in pure PFN-20BFO. To further investigate the influence of Mn and Gd doping on electrical conductivity, the DC j-E responses of samples were measured at RT (Fig. 5a). As expected, the highest σ DC was obtained in pure PFN-20BFO, namely ~3.0·10 −10 Ω −1 m −1 . The addition of 0.5 mol% Mn successfully reduce the electrical conductivity of the material, leading to ~20 times lower roomtemperature σ DC value, i.e., ~1.7·10 −11 Ω −1 m −1 . On the other hand, further addition of Gd does not significantly increase the electrical conductivity at RT, which still remains low and comparable to Mn-doped one (i.e., ~2.0·10 −11 Ω −1 m −1 ).
The influence of Mn and Gd doping on room-temperature P-E hysteresis loops measurements were investigated (Fig. 5b). Pure PFN-20BFO exhibits relatively slim and tilted relaxor-like P-E hysteresis loop with the coercive electric field (E C ) and saturated polarization (P S ) of ~2.8 kV·cm −1 and ~38 µC·cm −2 , respectively. The results are in agreement with the previous study. 42 The addition of 0.5 mol% Mn does not significantly affect the P-E response, where E C remains comparable to the one in pure PFN-20BFO, while P S is slightly lower, i.e., ~34 µC·cm −2 . Further addition of Gd resulted in even slimmer room-temperature hysteresis loop with lower E C and P S values, namely ~1.5 kV·cm −1 and ~32 µC·cm −2 , respectively. Note that this is most probably related to the lower peak-permittivity temperature and more pronounced cubic symmetry in this composition at RT (Supplementary material, B). In addition, both modified compositions exhibit sharp P-E loop ends, indicating low electrical conductivity of the materials even at higher applied electric fields.
As previously demonstrated in PFN 23,62 as well as BFO-based 33,34 ceramics, a reason that such a small amount of Mn successfully suppresses the enhanced electrical conductivity of the material could be related to the nature of the Mn ions, which are known to undergo different oxidation states (for example 4+, 3+ or 2+ in this case) and thus easily compensate possible charge imbalances in the system. Consequently, Mn ions most likely lead to lower concentration of charge point defects, leading to suppressed electrical conductivity at high temperatures and high electric fields. On the other hand, further addition of Gd does not significantly affect the ferroelectric and dielectric properties of the material, where dilution of ferroelectrically active Bi 3+ ions by magnetic Gd 3+ ions causes only a slight shift of T ε,max towards lower temperatures.  Fig. 6a. In all three cases an increasing trend of M with decreasing temperature is observed. Even the addition of only 0.5 mol% Mn ions slightly affect the magnetic properties of the material, where compared to pure PFN-20BFO, ~10% higher magnetization values are obtained over the entire measuring temperature range. As already mentioned, Fe-O-Fe linkages define magnetic properties in this system. [56][57][58] The substitution of larger Fe 3+ ions (64.5 pm, 3d 5 ) 53 with smaller Mn 4+ ions (53 pm, 3d 3 ) 53 changes the local structure of FeO 6 octahedra, 63,64 leading to contraction of the lattice parameters (Supplementary material, B). Consequently, shorter Mn-O bond and thus different angle of Fe-O-Mn linkage in comparison to the initial Fe-O-Fe leads to slightly different magnetic arrangement in the structure. 64,65 According to the previous studies, it was found that Fe-O-Mn can slightly distort the long-range antiferromagnetic order, inducing more ferrimagnetic-like spin configuration and thus slightly larger macroscopic magnetization. 66 The same trend, namely slightly increased magnetization in the entire temperature range, was also observed in Mn-doped PFN ceramics. 23 Further doping with Gd ions results in a significant enhancement of magnetic properties, especially at low temperatures. Compared to pure PFN-20BFO, Gd,Mn co-doped sample exhibits ~30% larger room-temperature magnetization (inset of Fig. 6a and Fig. 6b). While at lower temperatures more than 11 times higher maximum value of M is obtained, namely ~0.144 emu·g −1 at −270 °C and 1 kOe. This significantly higher magnetic response of Gd,Mn co-doped sample is also reflected in M-H hysteresis loops measurements at −270 °C (Fig. 6c). Similarly, large enhancement of M at low temperatures was also reported for Gd-doped BFO, 36,38,39,67 where such enhanced magnetic response is described by the alignment of Gd 3+ and Fe-O-Fe magnetic moments in the same direction. At low temperatures, paramagnetic Gd 3+ ions (partially filled seven 5f-orbitals) that possess high magnetic moments align with the antiferromagnetic Fe-O-Fe subsystem (f-d exchange orbital interaction), leading to significant enhancement of magnetization. 36,67,68 This contribution is less pronounced at higher temperatures, which is reflected in lower magnetization values. As seen above, the addition of Mn and Gd enhances the magnetic properties of PFN-20BFO but at the same time does not influence on its relaxor behavior that remains comparable to pure PFN-20BFO. Even more, both modified compositions exhibit relaxor-like hysteresis loops with sharp ends indicating low electrical conductivity even at higher applied electric fields, which make them promising candidates for multicalorics. Therefore, in the next step the EC and MC effects of Mn and Gd,Mn modified compositions were studied.
The temperature dependence of ΔT EC at different applied electric fields for Mn-doped PFN-20BFO is shown in Fig. 7a. The addition of 0.5 mol% Mn reduces the electrical conductivity, which enables the measurements of EC effects at higher temperatures and applied electric fields compared to pure PFN-20BFO (see Fig. S5a and compare with Fig. 7a). Mndoped PFN-20BFO exhibits high room-temperature ΔT EC value of ~1.5 °C at 150 kV·cm −1 , which is comparable with previously studied Mn-doped PFN. 23 By increasing the temperature, ΔT EC increases and reaches the maximum value of ~1.8 °C at 76 °C and 150 kV·cm −1 (Fig. 7a). Additional doping with Gd resulted in slightly lower ΔT EC , namely ~1.13 °C at RT and 140 kV·cm −1 . As seen in Fig. 4c, the peak-permittivity temperature of Gd,Mn co-doped sample is below RT. As the temperature increases, no additional anomalies are observed, resulting in a nearly constant ΔT EC of 1.15 °C ± 0.02 °C.
In Mn-doped PFN-20BFO at temperatures below 70 °C almost negligible JH (< 3%) is present, resulting in a good match of ΔT eff with the ΔT EC values (Fig. 7c). Above 80 °C, the JH starts to increase, causing a gradual lowering of effective EC cooling effect. Negligible JH (< 5 %) was also observed in Gd,Mn co-doped sample, where ΔT eff stays more or less identical to ΔT EC until 75 °C (Fig. 7d).
The temperature dependences of ΔT MC (calculated from M(T) curves shown in Supplementary material, C) at different applied magnetic fields for Mn-doped and Gd,Mn codoped PFN-20BFO are shown in Figs. 7e and 7f, respectively. A slight improvement in the magnetization of the Mn-doped sample in comparison to pure PFN-20BFO (see Fig. 6a) is also reflected in the MC effect, where ~20% higher room-temperature ΔT MC is observed. While at low temperatures ~2 times higher ΔT MC of ~0.15 °C at −270 °C and 50 kOe is obtained (insets of Figs. 7e). In agreement with enhanced magnetic properties, further addition of Gd ions causes a significant improvement of the MC effect over the entire temperature range, resulting in ~65% higher ΔT MC at RT and more than 30 times higher ΔT MC at −270 °C (~2.2 °C at 50 kOe) compared to pure PFN-20BFO (Fig. 7f). Furthermore, in this material room-temperature ΔT MC ~8·10 −3 °C is obtained at 90 kOe, which is so far the highest value among the PFN-based multicalorics. 22,23,69 On the other hand, the sample exhibits extremely high ΔT MC value of 3.6 °C at −270 °C and 90 kOe. Considering that this is an indirect measurement and that temperature of −273.15 °C (0 K) cannot be reached, we estimate that the ΔT MC value is below 3 °C.

IV. Summary and conclusions
In this work, a series of single-phase PFN-100xBFO solid solutions (x = 0-0.5) were prepared by mechanochemical synthesis followed by sintering. The influence of BFO addition on the temperatures of magnetic susceptibility and dielectric permittivity anomalies was studied. With increasing the amount of BFO, the paramagnetic-antiferromagnetic phase transition of PFN-BFO shifts toward higher temperatures, while the peak-permittivity temperature first decreases (to RT when x = 0.2) and then increases to higher temperatures (in x > 0.2 compositions). The dielectric permittivity as well as magnetic susceptibility anomalies meet at RT in PFN-20BFO composition, making it a promising candidate for multicaloric applications. However, due to the excessive Joule heating and its antiferromagnetic nature, PFN-20BFO exhibits low EC and MC effects in the entire measuring temperature range with the maximum values of 0.50 °C (at RT and 90 kV·cm −1 ) and 0.07 °C (at −270 °C and 50 kOe), respectively (Table 1). Therefore, EC and MC properties of PFN-20BFO were further improved by doping with Mn and Gd ions.
The addition of 0.5 mol% Mn successfully suppresses the electrical conductivity and consequently the Joule heating effect in the material, which resulted in high room-temperature EC effect of 1.5 °C, and even higher maximum value of 1.8 °C at 76 °C (Table 1a). At the same time, the addition of Mn also slightly improve the MC properties of the material, resulting in ~20% higher room-temperature and ~2 times higher maximum ΔT MC compared to pure PFN-20BFO (Table 1b). Further addition of 12 mol% Gd leads to significant enhancement of the MC effect in the entire measurement temperature range, exhibiting ~65% higher roomtemperature and more than 30 times higher maximum ΔT MC compared to pure PFN-20BFO (Table 1b). At the same time, the room-temperature electrical conductivity remains comparable to Mn-doped PFN-20BFO, which resulted in negligible Joule heating and EC effects of more than one degree over a wide temperature range. Despite the fact that in prepared PFN-20BFO-based materials both ferroic anomalies are located at RT, the maximum MC and EC effects are located in two different temperature regions. The use of such multicaloric material could be in applications where the operating temperature is significantly changing, for example in deep space and extreme measurement at liquid-helium temperature conditions. Further research in this area is needed to prepare better room-temperature single-phase multicalorics that can be used as cooling elements in future solid-state cooling devices for everyday use.

Supplementary material
Supporting materials are available in the online Supplementary material.