Microstructure and Properties of the Ferroelectric-Ferromagnetic PLZT-Ferrite Composites

The paper presents the technology of ferroelectric-ferromagnetic ceramic composites obtained from PLZT powder (the chemical formula Pb0.98La0.02(Zr0.90Ti0.10)0.995O3) and ferrite powder (Ni0.64Zn0.36Fe2O4), as well as the results of X-ray powder-diffraction data (XRD) measurement, microstructure, dielectric, ferroelectric, and magnetic properties of the composite samples. The ferroelectric-ferromagnetic composite (P-F) was obtained by mixing and the synthesis of 90% of PLZT and 10% of ferrite powders. The XRD test of the P-F composite shows a two-phase structure derived from the PLZT component (strong peaks) and the ferrite component (weak peaks). The symmetry of PLZT was identified as a rhombohedral ferroelectric phase, while the ferrite was identified as a spinel structure. Scanning electron microscope (SEM) microstructure analysis of the P-F ceramic composites showed that fine grains of the PLZT component surrounded large ferrite grains. At room temperature P-F composites exhibit both ferroelectric and ferromagnetic properties. The P-F composite samples have lower values of the maximum dielectric permittivity at the Curie temperature and a higher dielectric loss compared to the PLZT ceramics, however, the exhibit overall good multiferroic properties.

The modified PZT (PbZr 1−x Ti x O 3 ) materials with perovskite structures have been practically the most widely known and widely used ones for many years [18][19][20][21][22][23][24].The lead lanthanum zirconate titanate (PLZT) ceramic is a well-known perovskite material, too, with the general chemical formula of Pb 1−x La x (Zr y Ti 1−y ) 1−0.25x V B 0.25x O 3 .Similarly to the commonly-used modified PZT materials, compositional changes within the PLZT material can significantly alter the electrophysical properties thereof.PLZT materials have very good dielectric, ferroelectric, and piezoelectric properties which allow the usage of the PLZT-type material in a variety of applications [25][26][27][28][29].
The technological process of obtaining composite (P-F) samples consisted of several stages.The PLZT powder with the following chemical formula Pb 0.98 La 0.02 (Zr 0.90 Ti 0.10 ) 0.995 O 3 (material with a perovskite structure) and the ferrite powder with the chemical formula Ni 0.64 Zn 0.36 Fe 2 O 4 , were used to obtain the composite samples.The synthesis of the composite powder was carried out conventionally.
In the first stage, the PLZT ceramic powder was obtained from stoichiometric mixtures of PbO (99.9% Sigma-Aldrich, St. Louis, MO, USA), ZrO 2 (99% Tosoh, Amsterdam, The Netherlands), TiO 2 (99.8%Alfa Aesar, Karlsruhe, Germany) and La(OH) 3 prepared from La 2 O 3 (99.99%Alfa Aesar, Karlsruhe, Germany) by equilibrating it in a humid atmosphere.The technology of the PLZT powder was carried out using 3-steps of milling.The starting powders were homogenized in a planetary mill at 200 rpm for 2 h.Successively, after drying, the mixtures were calcined at 900 • C for 2 h, re-milled, and re-calcined.After a second calcination, the powders were mixed in an attritor mill and dried [35].
In the third stage, two powders (PLZT and ferrite) were mixed in the proportion of 90% of PLZT and 10% of ferrite, using a Fritsch planetary ball mill for 15 h (wet in ethyl alcohol medium).The mixture was synthesized by calcination under the following conditions: 1000 • C for 4 h.
The final sintering of the ceramic P-F composite samples was conducted under the following conditions: T s1 = 1200 • C (P-F1), T s2 = 1250 • C (P-F2) and T s3 = 1300 • C (P-F3) for t s = 2 h.The three ceramic composites obtained in the above sintering conditions have been compared in this work.For the electrical measurements, sintered pellets with a thickness of 1.0 mm were prepared.Stress was removed by annealing the pellets at 750 • C. The surfaces of the pellets were covered with a silver electrode.

Characterization
Granulometric analyses of the milled PLZT powders were performed on a Cilas Hr-850-B laser granulometer manufactured by Alcatel (Compagnie Industrielle Des Lasers, Marcoussis, France).The sintering behavior of the calcinated PLZT powders was followed using a heating-stage microscope (Leitz, Wetzlar, Germany) in the temperature range from 25 • C to 1400 • C, in an oxygen atmosphere.The synthesis temperature of the P-F composite powder was selected on the basis of differential thermal analysis DTA, as well as DTG and TG, with the usage of a Q-1500D derivatograph (MOM, Q-1500D, F. Paulik, J. Paulik, L. Erdey system, Budapest, Hungary) in the temperature range from 20 The X-ray powder-diffraction data (XRD) of the sintered P-F samples were collected at room temperature on a Phillips X'Pert Pro (PANalytical, Phillips X'Pert Pro, Eindhoven, The Netherlands) diffractometer (Cu-Kα radiation).The data were collected in the 2θ range from 10 • to 65 • , in steps of 0.02 degrees, with an integration time of 4 s/step.A scanning electron microscope SEM (Jeol Lid., JEOL JSM-7100 TTL LV, Tokyo, Japan), equipped with an energy-dispersive system EDS (Jeol Lid., EX-35170EDES, Tokyo, Japan) and a backscattered electron detector BSE (Jeol Lid., SCM-25D170, Tokyo, Japan) was used for the microstructural analysis.Prior to the SEM/EDS analyses, the samples were coated with gold to provide electrical conductivity and to avoid any charging effects.
Dielectric measurements were performed on a capacity bridge (QuadTech, Inc., 1920 Precision LCR meter, Maynard, MA, USA), for a cycle of heating (in temperature range from 20 • C to 420 • C).DC electrical conductivity has been measured using an electrometer (Keithley Instruments, Inc., 6517B, Cleveland, OH, USA) in the same temperature range.Ferroelectric hysteresis P-E loops were made using a Sawyer-Tower circuit and a high voltage amplifier (Matsusada Precision Inc., HEOPS-5B6, Kusatsu, Japan) while the data were stored on a computer disc using an A/D, D/A transducer card and the LabView computer program.Magnetic properties were obtained by applying the Quantum Design PPMS system (Quantum Design, PPMS 7T ACMS module, San Diego, CA, USA).Dynamic magnetic susceptibility (the real and imaginary part and magnetic loss) was measured versus the AC magnetization field in the range from 40 A/m to 1200 A/m for 50 Hz, 120 Hz, and 1000 Hz, as soon as the field frequency (ranging from 50 Hz to 2000 Hz, magnetic field 800 A/m).All magnetic measurement was carried out at room temperature.

Analysis of Ceramic Powder
The histogram of particle size distribution of the PLZT powder (Figure 1a) shows a single-modal Gaussian distribution, with the average particle size r a = 0.82 µm.It confirms the homogeneous powder of the PLZT with the fine particle size.Figure 1b shows the shrinkage-temperature curves of the calcined PLZT powder.Judging by the shape of the curves, it is clear that the sintering process is finished at 1100 The synthesis temperature of the P-F composite powder was determined with DTA/TG analysis (Figure 2).The DTA result has shown that the synthesis of the components of the analyzed material powder occurs to the temperature of 650 °C.The TG curve has shown the characteristic changes associated with the weight loss of the ceramic sample during heating.The extreme weight loss falls at a temperature of about 320 °C, and is associated with the evaporation of moisture from the sample.This confirms the occurrence of the peak on the DTA curve at 109 °C (endothermic maximum).The peak on the DTA curve (at 465 °C) is related to the phenomenon of nucleation and the formation of intermediate phases (including non-perovskite-pyrochlore phase).The right heating speed during the technological process eliminates the possibility of creating undesirable phases.The second clear mass loss on the TG curve is related to the formation of the perovskite The synthesis temperature of the P-F composite powder was determined with DTA/TG analysis (Figure 2).The DTA result has shown that the synthesis of the components of the analyzed material powder occurs to the temperature of 650 • C. The TG curve has shown the characteristic changes associated with the weight loss of the ceramic sample during heating.The extreme weight loss falls at a temperature of about 320 • C, and is associated with the evaporation of moisture from the sample.This confirms the occurrence of the peak on the DTA curve at 109 • C (endothermic maximum).The peak on the DTA curve (at 465 • C) is related to the phenomenon of nucleation and the formation of intermediate phases (including non-perovskite-pyrochlore phase).The right heating speed during the technological process eliminates the possibility of creating undesirable phases.The second clear mass loss on the TG curve is related to the formation of the perovskite phase, which confirms the occurrence of the peak at 643 • C on the DTA curve (endothermic maximum).Above the temperature of 750 • C further weight loss is not observed.In the whole range temperature the weight loss does not exceed 0.7%.The synthesis temperature of the P-F composite powder was determined with DTA/TG analysis (Figure 2).The DTA result has shown that the synthesis of the components of the analyzed material powder occurs to the temperature of 650 °C.The TG curve has shown the characteristic changes associated with the weight loss of the ceramic sample during heating.The extreme weight loss falls at a temperature of about 320 °C, and is associated with the evaporation of moisture from the sample.This confirms the occurrence of the peak on the DTA curve at 109 °C (endothermic maximum).The peak on the DTA curve (at 465 °C) is related to the phenomenon of nucleation and the formation of intermediate phases (including non-perovskite-pyrochlore phase).The right heating speed during the technological process eliminates the possibility of creating undesirable phases.The second clear mass loss on the TG curve is related to the formation of the perovskite phase, which confirms the occurrence of the peak at 643 °C on the DTA curve (endothermic maximum).Above the temperature of 750 °C further weight loss is not observed.In the whole range temperature the weight loss does not exceed 0.7%.

X-ray Diffraction Studies
X-ray powder-diffraction patterns of the PLZT, ferrite, and P-F materials are shown in Figure 3.The XRD studies of the ceramic powders have shown the absence of foreign phases.The diffraction

X-ray Diffraction Studies
X-ray powder-diffraction patterns of the PLZT, ferrite, and P-F materials are shown in Figure 3.The XRD studies of the ceramic powders have shown the absence of foreign phases.The diffraction peaks in the PLZT 2/90/10 patterns showed a perovskite structure, while the ferrite powder (Ni 0.64 Zn 0.36 Fe 2 O 4 ) has a spinel structure.According to the phase diagram of PLZT proposed by Haertling [36] for low La 3+ content (x ≤ 0.02), material PLZT 2/90/10 can be well indexed with the rhombohedral ferroelectric phase.For PLZT compositions with La 3+ content from the area of 0.025 ≤ x < 0.035, the coexistence of orthorhombic-rhombohedral phases occurs, while with high La 3+ content (x ≥ 0.035), the reflection is related to the presence of orthorhombic AFE phase [37].The symmetry of PLZT with high La 3+ content in [38] was also explained by a disordered rhombohedral, a monoclinic, or an orthorhombic phase.
The XRD test of the PLZT powder showed that the best-matched results were obtained for the pattern (no.PDF 77-1194) with a rhombohedral ferroelectric phase and R3c space group (with unit cell parameters: a 0 = 5.8410 Å, c 0 = 14.4160Å, and α = 89.9953• [26].In the case of the P-F composite material, the X-ray test shows peaks of the two components.Very strong peaks are derived from the PLZT component and the other weak peaks are derived from the ferrite materials with a spinel structure.
The XRD test of the PLZT powder showed that the best-matched results were obtained for the pattern (no.PDF 77-1194) with a rhombohedral ferroelectric phase and R3c space group (with unit cell parameters: a0 = 5.8410 Å, c0 = 14.4160Å, and α = 89.9953°[26].In the case of the P-F composite material, the X-ray test shows peaks of the two components.Very strong peaks are derived from the PLZT component and the other weak peaks are derived from the ferrite materials with a spinel structure.

Microstructural Properties
The SEM microstructure of the PLZT ceramic sample (obtained for the following sintering conditions: Ts = 1250 °C for ts = 2 h), characterized by a non-porous microstructure with densely packed and properly crystallized grains, is shown in Figure 4a.
The P-F composite samples (Figure 4b-d) have a fine-grained microstructure, as compared to the PLZT ceramics: small, with well-formed shaped grains of the PLZT component, surrounded by larger grains of the ferrite with a characteristic pyramidal shape.This is clearly visible in pictures taken by the backscattered electron (BSE) technique (Figure 4e-g).The ferrite grains are evenly distributed in the whole microstructure of the composite samples.As the sintering temperature grows, the increase of the average grain size is observed.The increase of the sintering temperature to 1250 °C (P-F2) causes a significant growth of the grain size.Additionally, disparities between the PLZT and ferrite grains size are much more visible.At a sintering temperature of 1300 °C, the grains grow significantly (P-F3), whereas ferrite grains are of an irregular shape.The grain boundaries of the PLZT/ferrite composites are irregular and the microstructure shows a high porosity.The high

Microstructural Properties
The SEM microstructure of the PLZT ceramic sample (obtained for the following sintering conditions: T s = 1250 • C for t s = 2 h), characterized by a non-porous microstructure with densely packed and properly crystallized grains, is shown in Figure 4a.
The P-F composite samples (Figure 4b-d) have a fine-grained microstructure, as compared to the PLZT ceramics: small, with well-formed shaped grains of the PLZT component, surrounded by larger grains of the ferrite with a characteristic pyramidal shape.This is clearly visible in pictures taken by the backscattered electron (BSE) technique (Figure 4e-g).The ferrite grains are evenly distributed in the whole microstructure of the composite samples.As the sintering temperature grows, the increase of the average grain size is observed.The increase of the sintering temperature to 1250 • C (P-F2) causes a significant growth of the grain size.Additionally, disparities between the PLZT and ferrite grains size are much more visible.At a sintering temperature of 1300 • C, the grains grow significantly (P-F3), whereas ferrite grains are of an irregular shape.The grain boundaries of the PLZT/ferrite composites are irregular and the microstructure shows a high porosity.The high sintering temperature adversely affects the ceramic composite microstructure-too high an increase of the grain leads to increased porosity and mechanical stresses.
The EDS analysis confirmed the presence of constituent elements in the test of the P-F composite samples.The results of the percentage of individual components of the P-F composite, summarized in Table 1, are the averaging of 10 randomly-chosen areas of the surface of the sample.In the case of the P-F1 and P-F2 composite samples, lead, lanthanum, and titanium deficiencies, as well as a small excess of iron, zirconium, zinc and nickel are observed, as compared to theoretical calculations.All deviations from the initial composition are within the acceptable range.In the case of the P-F3 sample, the disproportions of the components from the assumed composition are the highest (mainly a large loss of lead).composite samples.The results of the percentage of individual components of the P-F composite, summarized in Table 1, are the averaging of 10 randomly-chosen areas of the surface of the sample.In the case of the P-F1 and P-F2 composite samples, lead, lanthanum, and titanium deficiencies, as well as a small excess of iron, zirconium, zinc and nickel are observed, as compared to theoretical calculations.All deviations from the initial composition are within the acceptable range.In the case of the P-F3 sample, the disproportions of the components from the assumed composition are the highest (mainly a large loss of lead).

DC Electrical Conductivity Measurements
At room temperature, the ρ DC resistivity of the PLZT ceramics is 1.2 × 10 9 Ωm, while at the phase transition, the temperature ρ DC is 3.0 × 10 7 Ωm.At room temperature, the resistivity of the composite samples does not increase significantly.At higher temperatures, all the P-F composite samples have a higher electric conductivity, as compared to the PLZT ceramics (Figure 5).
The activation energy for the PLZT ceramics and the P-F composite samples was calculated according to Arrhenius' law [39]: where σ 0 is the pre-exponential factor, k B is Boltzmann's constant, T is the absolute temperature, and E Act is the activation energy appointed from the slope of lnσ DC vs. the 1/T plot.The values of the activation energy E Act at two characteristic regions are presented in Table 2.
a higher electric conductivity, as compared to the PLZT ceramics (Figure 5).The activation energy for the PLZT ceramics and the P-F composite samples was calculated according to Arrhenius' law [39]: where σ0 is the pre-exponential factor, kB is Boltzmann's constant, T is the absolute temperature, and EAct is the activation energy appointed from the slope of lnσDC vs. the 1/T plot.The values of the activation energy EAct at two characteristic regions are presented in Table 2.

Dielectric Properties
The temperature dependences of dielectric permittivity (ε) for the P-F composites are presented in Figure 6.At room temperature, the values of dielectric permittivity are the same for the analyzed samples.The combination of the PLZT powder and the ferrite powder to the composite form significantly reduces the maximum of dielectric permittivity at the Curie temperature (the solid line for 1 kHz in Figure 6) and shifts the phase transition slightly towards the higher temperatures.A sharp phase transition from the ferroelectric to the paraelectric phase occurs in the PLZT, while in the P-F composite samples the phase transition is characterized by the broader peaks.One of the reasons for the broad temperature range of the phase transition may be related to the disorder in the distribution of B-side ions in the perovskite cell.This leads to random local Curie temperatures in different regions of the composite material.However, the obtained composite does not lose good dielectric properties.Comparing the composite samples which were sintered at the different temperatures, it can be concluded that the highest values of dielectric permittivity are shown in the P-F2 sample.On the basis of this fact, it could be presumed that the optimal temperature in the technological process of the P-F composite is 1250 °C.Both the lower and the higher sintering temperatures extend the blur phase transition and decrease the values of dielectric permittivity.
The temperature dependencies of the dielectric loss tangent (tanδ) for the P-F ceramic composites are presented in Figure 7.The PLZT ceramics have very low values of dielectric loss (solid line for 1 kHz in Figure 7).At room temperature, the values of the dielectric loss for the P-F composite samples are low, as well.
The combination of the PLZT powder and the ferrite powder to the composite form contributes to increased dielectric loss, which can be clearly observed on the temperature waveforms.Additionally, the rise of the sintering temperature of composites causes an increase of the values of dielectric loss over the entire measurement area.However, the obtained composite does not lose good dielectric properties.Comparing the composite samples which were sintered at the different temperatures, it can be concluded that the highest values of dielectric permittivity are shown in the P-F2 sample.On the basis of this fact, it could be presumed that the optimal temperature in the technological process of the P-F composite is 1250 • C. Both the lower and the higher sintering temperatures extend the blur phase transition and decrease the values of dielectric permittivity.
The temperature dependencies of the dielectric loss tangent (tanδ) for the P-F ceramic composites are presented in Figure 7.The PLZT ceramics have very low values of dielectric loss (solid line for 1 kHz in Figure 7).At room temperature, the values of the dielectric loss for the P-F composite samples are low, as well.
The combination of the PLZT powder and the ferrite powder to the composite form contributes to increased dielectric loss, which can be clearly observed on the temperature waveforms.Additionally, the rise of the sintering temperature of composites causes an increase of the values of dielectric loss over the entire measurement area.The combination of the PLZT powder and the ferrite powder to the composite form contributes to increased dielectric loss, which can be clearly observed on the temperature waveforms.Additionally, the rise of the sintering temperature of composites causes an increase of the values of dielectric loss over the entire measurement area.

Ferroelectric Properties
The electric hysteresis loops tested at room temperature (at ν = 1 Hz frequency) of the PLZT ceramics and the P-F composite samples are presented in Figure 8.The PLZT ceramics have good saturated hysteresis loops (characteristic for ferroelectric materials) with a coercive field EC = 1.40 kV/mm (the value of spontaneous polarization Ps is 23.10 μC/cm 2 and the value of the remnant polarization is Pr = 21.10 μC/cm 2 ).In the case of the P-F composite samples, the coercive field

Ferroelectric Properties
The electric hysteresis loops tested at room temperature (at ν = 1 Hz frequency) of the PLZT ceramics and the P-F composite samples are presented in Figure 8.The PLZT ceramics have good saturated hysteresis loops (characteristic for ferroelectric materials) with a coercive field E C = 1.40 kV/mm (the value of spontaneous polarization P s is 23.10 µC/cm 2 and the value of the remnant polarization is P r = 21.10 µC/cm 2 ).In the case of the P-F composite samples, the coercive field increases and the loop loses its saturation, even with higher applied electrical fields (Table 2).
In the case of the P-F3 sample, these are extremely large.The P-E tests of the PF composite samples show that the FE behavior decreases with increasing sintering temperature.
Symmetry 2018, 10, x FOR PEER REVIEW 10 of 14 increases and the loop loses its saturation, even with higher applied electrical fields (Table 2).In the case of the P-F3 sample, these changes are extremely large.The P-E tests of the PF composite samples show that the FE behavior decreases with increasing sintering temperature.

Magnetic Properties
The AC magnetic measurements were carried out at room temperature.These investigations allowed the determination of the real and imaginary parts of susceptibility, as well as magnetic loss, for several fields and frequencies (0.5-15 Oe, i.e., 40-1200 A/m, 50 Hz, 120 Hz, and 1000 Hz) for the investigated P-F ceramic composites.
The dependence of the real and the imaginary parts of the susceptibility (a, c, e), as well as the magnetic loss tangent (b, d, f) versus the intensity of the magnetic field (at frequencies of 50, 120, and 1000 Hz) obtained for P-F composites, are presented on Figure 9.In the case of P-F1 and P-F2 samples, a similar dependence was observed.On the contrary, for the PF-3 sample, the magnetic loss at low frequencies was much higher.
Figure 10 shows the real and the imaginary part of the magnetic susceptibility, as well as the magnetic loss tangent versus the frequency of the magnetic field (10 Oe = 800 A/m), for the investigated samples measured at room temperature.The decrease of magnetic loss up to 500 Hz for all composite samples was determined.For higher frequencies, the increase of the magnetic loss was observed.
The changes of magnetic loss of the P-F composite samples are probably caused by the relaxation processes in the material.The highest magnetic loss occurred for the P-F3 composite sample, while for the P-F1 one, the magnetic loss was the lowest.

Magnetic Properties
The AC magnetic measurements were carried out at room temperature.These investigations allowed the determination of the real and imaginary parts of susceptibility, as well as magnetic loss, for several fields and frequencies (0.5-15 Oe, i.e., 40-1200 A/m, 50 Hz, 120 Hz, and 1000 Hz) for the investigated P-F ceramic composites.
The dependence of the real and the imaginary parts of the susceptibility (a, c, e), as well as the magnetic loss tangent (b, d, f) versus the intensity of the magnetic field (at frequencies of 50, 120, and 1000 Hz) obtained for P-F composites, are presented on Figure 9.In the case of P-F1 and P-F2 samples, a similar dependence was observed.On the contrary, for the PF-3 sample, the magnetic loss at low frequencies was much higher.
Figure 10 shows the real and the imaginary part of the magnetic susceptibility, as well as the magnetic loss tangent versus the frequency of the magnetic field (10 Oe = 800 A/m), for the investigated samples measured at room temperature.The decrease of magnetic loss up to 500 Hz for all composite samples was determined.For higher frequencies, the increase of the magnetic loss was observed.
The changes of magnetic loss of the P-F composite samples are probably caused by the relaxation processes in the material.The highest magnetic loss occurred for the P-F3 composite sample, while for the P-F1 one, the magnetic loss was the lowest.
Symmetry 2018, 10, x FOR PEER REVIEW 10 of 14 increases and the loop loses its saturation, even with higher applied electrical fields (Table 2).In the case of the P-F3 sample, these changes are extremely large.The P-E tests of the PF composite samples show that the FE behavior decreases with increasing sintering temperature.

Magnetic Properties
The AC magnetic measurements were carried out at room temperature.These investigations allowed the determination of the real and imaginary parts of susceptibility, as well as magnetic loss, for several fields and frequencies (0.5-15 Oe, i.e., 40-1200 A/m, 50 Hz, 120 Hz, and 1000 Hz) for the investigated P-F ceramic composites.
The dependence of the real and the imaginary parts of the susceptibility (a, c, e), as well as the magnetic loss tangent (b, d, f) versus the intensity of the magnetic field (at frequencies of 50, 120, and 1000 Hz) obtained for P-F composites, are presented on Figure 9.In the case of P-F1 and P-F2 samples, a similar dependence was observed.On the contrary, for the PF-3 sample, the magnetic loss at low frequencies was much higher.
Figure 10 shows the real and the imaginary part of the magnetic susceptibility, as well as the magnetic loss tangent versus the frequency of the magnetic field (10 Oe = 800 A/m), for the investigated samples measured at room temperature.The decrease of magnetic loss up to 500 Hz for all composite samples was determined.For higher frequencies, the increase of the magnetic loss was observed.
The changes of magnetic loss of the P-F composite samples are probably caused by the relaxation processes in the material.The highest magnetic loss occurred for the P-F3 composite sample, while for the P-F1 one, the magnetic loss was the lowest.

Conclusions
The paper presents the technological process of the ferroelectromagnetic composite (P-F) based on ferroelectric Pb0.98La0.02(Zr0.90Ti0.10)0.995O3material (PLZT) and the magnetic ferrite material (with the chemical formula of Ni0.64Zn0.36Fe2O4),with a PLZT/ferrite proportion equal to 90/10.The SEM microstructure analysis of the P-F composite showed that fine grains of the PLZT component surrounded larger ferrite grains.The ferrite grains had the shape of pyramid.The increase of the sintering temperature causes the increase in the average grain size of the P-F composite.

Conclusions
The paper presents the technological process of the ferroelectromagnetic composite (P-F) based on ferroelectric Pb0.98La0.02(Zr0.90Ti0.10)0.995O3material (PLZT) and the magnetic ferrite material (with the chemical formula of Ni0.64Zn0.36Fe2O4),with a PLZT/ferrite proportion equal to 90/10.The SEM microstructure analysis of the P-F composite showed that fine grains of the PLZT component surrounded larger ferrite grains.The ferrite grains had the shape of a pyramid.The increase of the sintering temperature causes the increase in the average grain size of the P-F composite.

Conclusions
The paper presents the technological process of the ferroelectromagnetic (P-F) based on ferroelectric Pb 0.98 La 0.02 (Zr 0.90 Ti 0.10 ) 0.995 O 3 material (PLZT) and the magnetic ferrite material (with the chemical formula of Ni 0.64 Zn 0.36 Fe 2 O 4 ), with a PLZT/ferrite proportion equal to 90/10.The SEM microstructure analysis of the P-F composite showed that fine grains of the PLZT component surrounded larger ferrite grains.The ferrite grains had the shape of a pyramid.The increase the sintering temperature causes the increase in the average grain size of the P-F composite.
At room temperature, the P-F ceramic composites show both ferroelectric and ferromagnetic properties.The frequency changes of the magnetic loss of the P-F composite samples are probably caused by relaxation processes in the material.The combination of the PLZT powder and the ferrite powder which form the composite of P-F significantly reduces the maximum dielectric permittivity at the Curie temperature, but the obtained P-F composite does not lose good dielectric properties, and exhibits overall good multiferroic properties.Both too low and too high sintering temperatures deteriorate the dielectric properties of the P-F composite samples.A high sintering temperature (1300 • C) also causes excessive grain growth, as well as increases in the heterogeneity of the microstructure.The loss of homogeneity of the microstructure composite samples deteriorates the electrophysical properties of the P-F composites.The optimal sintering temperature is 1250 • C, for a sintering time equal to 2 h.
The obtained P-F composite sample exhibits good useful properties, giving the possibility to use these types of materials to construct magnetoelectric transducers, as well as sensors in micromechatronics devices.

14 Figure 1 .
Figure 1.Graphs of PLZT powder particle size distribution the histogram and the summation curve (red line) (a) and the shrinkage-temperature curves for PLZT powder with a red reference line (b).

Figure 1 .
Figure 1.Graphs of PLZT powder particle size distribution the histogram and the summation curve (red line) (a) and the shrinkage-temperature curves for PLZT powder with a red reference line (b).

Figure 1 .
Figure 1.Graphs of PLZT powder particle size distribution the histogram and the summation curve (red line) (a) and the shrinkage-temperature curves for PLZT powder with a red reference line (b).

Figure 2 .
Figure 2. DTA, TG, and DTG of the P-F composite powder.

Figure 2 .
Figure 2. DTA, TG, and DTG of the P-F composite powder.

Figure 3 .
Figure 3. X-ray diffraction patterns of the PLZT, ferrite, and P-F composite materials.

Figure 3 .
Figure 3. X-ray diffraction patterns of the PLZT, ferrite, and P-F composite materials.

Figure 5 .
Figure 5.The lnσDC(1/T) relationship for the PLZT ceramics and P-F composite samples.

Figure 5 .
Figure 5.The lnσ DC (1/T) relationship for the PLZT ceramics and P-F composite samples.

Figure 9 .
Figure 9.The dependence of the real and the imaginary part of the susceptibility (a,c,e), and the magnetic loss tangent (b,d,f) versus the intensity of the magnetic field for P-F composites, obtained at room temperature.

Figure 10 .
Figure 10.The dependence of the real (inserted) and the imaginary (a) part of the susceptibility, and the magnetic loss tangent (b) versus frequency of the magnetic field for P-F composites, obtained at room temperature.

Figure 9 .Figure 9 .
Figure 9.The dependence of the real and the imaginary part of the susceptibility (a,c,e), and the magnetic loss tangent (b,d,f) versus the intensity of the magnetic field for P-F composites, obtained at room temperature.

Figure 10 .
Figure 10.The dependence of the real (inserted) and the imaginary (a) part of the susceptibility, and the magnetic loss tangent (b) versus frequency of the magnetic field for P-F composites, obtained at room temperature.

Figure 10 .
Figure 10.The dependence of the real (inserted) and the imaginary (a) part of the susceptibility, and the magnetic loss tangent (b) versus frequency of the magnetic field for P-F composites, obtained at room temperature.

Table 1 .
Theoretical and experimental percentages of elements of the P-F composites (expressed as oxides).

Table 1 .
Theoretical and experimental percentages of elements of the P-F composites (expressed as oxides).

Table 2 .
The electrophysical parameters of the obtained composite samples.

Table 2 .
The electrophysical parameters of the obtained composite samples.