In this research, we have explored the structural, morphological, electrical, and magnetic properties of Mn2+ substituted polycrystalline Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (where x = 0.0, 0.05, 0.10, 0.15, and 0.20) ceramic samples prepared by the standard solid-state reaction system. The x-ray diffraction pattern of the 20% Mn-doped BCSMTO sample has confirmed a cubic to tetragonal structural phase transition. The lattice parameter is found to increase at 15% and 20% Mn concentrations, which is caused due to the inequality in ionic radii of cations. Scanning electron microscope analysis exhibited that, with the increase in Mn contents (x = 0.0–0.15), the average grain size of the samples gets bigger but then significantly decreased at 20% Mn substitution. Dielectric constants for all the samples are higher for lower frequency regions and remain independent at a higher frequency domain. The initial permeability remains almost constant at lower frequencies and then sharply falls at the cut-off frequency, which is in agreement with the Globus model. Among all the studied samples, the 10% Mn-doped ceramic sample shows the highest relative quality factor with significantly demolished loss (tan δM). At room temperature, the M–H loop for the 10% Mn-doped sample demonstrates a domination of diamagnetic nature at a higher magnetic field. The decrement in saturation magnetization with Mn addition suggests that the double-exchange interactions in tetragonal BaTiO3 may have been weakened. The outcome of this analysis emphasizes the impact of Mn as a doping element with 10% concentration in Ba0.4Ca0.4Sr0.2MnxTi1−xO3 that provides enhanced structural and electrical properties, which are associated with homogeneous grain size, reduced porosity, and lower tangent loss.

In recent years, polycrystalline ferroelectric ceramics have received considerable attention compared to single crystals because of their ease of preparation and acquisition of various desired properties by compositional modifications.1,2 Ferroelectric ceramics or ferroelectrics refer to the group of dielectrics having the property of spontaneous polarization (i.e., they retain a dipole even after an applied voltage has been removed). Among the ferroelectric materials, BaTiO3 (BTO) based ceramics have been considered to be attractive candidates toward this aspect as compared to lead-based ferroelectric materials due to their simple crystal structure, extremely high dielectric constant, low leakage current, high stability, and anisotropic optical behavior.3 Moreover, in the electronic industry, BaTiO3 powders have been widely used for multi-layer ceramic capacitor (MLCC) production, which is relevant for various applications, such as infrared detectors, waveguide modulators, piezoelectric transducers, and ferroelectric memory gate dielectrics.4 After the revelation of ferroelectricity and piezoelectricity in BaTiO3 ceramics, in 1946, it was discovered that it had a dielectric constant 100 times larger compared to the other insulators.5,6 These findings have allowed us to manufacture smaller sized capacitors but with a larger capacity than other materials.7 However, it was observed that in the BTO earthenware production when the temperature is brought down to around 400 K, the precious stone structure of BTO changes from cubic to tetragonal due to the ionic relocation of titanium (Ti) particles, and therefore, it goes to a ferroelectric state. This transformation occurs close to room temperature, and thus, these limited stable temperature scopes of the ferroelectric tetragonal stage are hindrances for ferroelectric and piezoelectric applications.8 

In order to surpass these limitations for accomplishing the industrial requirements and to find a way of extending the working temperature, many experiments have been done by researchers, one of which is the doping of metal ions. BTO is an ABO3 type perovskite ceramic material; due to the simplicity of its crystal structure, it can easily accommodate different types of ions at the same time to form solid solutions. Dopant incorporation in ABO3 type ceramics plays an influential role in controlling the grain size and hence in the improvement of electrical characteristics of ceramic. A great deal of work has mentioned that the dielectric characteristics of polycrystalline BTO materials depend on the grain growth during the sintering process and on the concentration of the additive types.9,10 Notably, it is demonstrated that doping can transform the BTO structure into a double perovskite form, which can blend dopants on both A-site and B-site and this aspect can influence a huge change in its magnetic and dielectric properties at room temperature and it can even display multiferroicity.11 A large number of studies are available in the case of Mn doping at the B-site of BTO.12,13 It was observed that Mn doping diminishes the conductance of BaTiO3 because it constrains the electrons in a small region and therefore, the number of electrons starts to decrease with the increase in the Mn concentration.14 Moreover, former investigations reported that adding a small amount of Mn to BTO causes a ferroelectric change at Curie temperature TC.15,16 In the BaTiO3 lattice, Mn2+ substitution can occupy the Ti4+ site, but it requires a charge in compensation. Therefore, oxygen vacancies are created by this charge compensation mechanism and this might be the reason for the Mn2+ (ionic radius rMn = 0.70 Å) occupancy at the B-site (Ti4+ site, ionic radius rTi = 0.60 Å) of perovskite BaTiO3. Again, as doping elements, alkaline earth metals can improve various physical properties including ferroelectricity and conductivity in BaTiO3 and it can also induce structural transformations.17 Park et al. reported that the substitution of Ca and Sr (which are alkaline earth metals) for every cation in BaTiO3 has a noteworthy effect on the Curie temperature (TC) and hence it enhances the electrical properties.4 Furthermore, it was observed that the substitution of a Ca2+ ion at the A-site of BTO reduces the possibility of the development of a non-ferroelectric undesirable hexagonal phase.18 

Thus, in our present investigation, we have aimed to observe what change Mn doping (at the B-site) brings about when we keep Ca and Sr in a fixed ratio at the A-site (in a molar ratio of 0.4 and 0.2). However, to the best of our knowledge, a well-organized survey on the synthesis and characterization of Ca, Sr, and Mn substituted BaTiO3 ceramics still seems uncharted. Therefore, in our present study, Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (where x = 0.0, 0.05, 0.10, 0.15, and 0.20) ceramics are formulated using the conventional solid-state reaction method with the hope of having good structural, morphological, electrical, and magnetic properties through the Mn addition in the B-site.

Polycrystalline ceramic Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (where x = 0.0, 0.05, 0.10, 0.15, and 0.20) samples were formulated successfully by means of the solid-state reaction technique. Highly purified reagent grade oxides: CaO, MnCO3, BaCO3, TiO2, and SrO were used as raw materials. At first, all these raw materials were weighed accurately, and then allowed to form a homogeneous mixture, all the powders were blended in an agate mortar. After that, using a programmable furnace, this powder was calcined at 800 °C with a 3 h holding time. In order to enhance the level of homogeneity, the calcined mixtures were re-ground for 2 h. Then, these ground powders were pressed into disk-shaped pellets (the diameter is 12 mm and the thickness is 1 mm) with a uniaxial hydraulic pressing system at hand pressure. Finally, setting a 4 h holding time, the obtained pellets were sintered at 1200 °C and then again at 1250 °C. These final pellets and the ground sintered powders were then used to yield the prescribed structural, morphological, electrical, and magnetic properties of Mn-doped BCSMTO (Ba0.4Ca0.4Sr0.2MnxTi1−xO3) ceramic samples.

To explore the crystal structure of the synthesized composites, x-ray diffraction (XRD) patterns were obtained using x-ray powder diffractometer (Philips X’Pert Pro, PW 3040), which provides Cu-Kα radiation at a wavelength of 0.154 nm in the 2θ range of 20°–70°. The lattice parameter (a), bulk density (ρB), x-ray density (ρth), and porosity (P) were calculated from the XRD data. The morphological analysis of the surface of all the ceramic samples was carried out using a Scanning Electron Microscope (SEM). A computer software (ImageJ) was used to attain the grain size and distribution profile of the synthesized samples. Ceramic pellets, pasted with silver on both sides, were used to examine the room temperature dielectric behavior with the variation of frequency ranging from 1 kHz to 100 MHz with the help of a Wayne Kerr Impedance Analyzer (6500B). The Wayne Kerr Impedance Analyzer (6500B) was again used to measure the real part of permeability μi with the frequency variation using the equation μi = Ls/L0, where Ls is the self-inductance of the composites in the presence of the magnetic core and L0 is the inductance of the winding coil in the absence of the sample core, respectively. L0 = µ0N2S/Πd¯, where µ0 is the permeability of vacuum, S is the cross sectional area, N (=5) is the number of coil turns, and d¯ is the mean diameter of the specimen. If d1, d2, and h are the inner diameter, outer diameter, and thickness of the sample, then S = (d1 − d2) × h and d¯ = (d1 + d2)/2. The M–H hysteresis curves (magnetization vs applied magnetic field) of all the compositions were accomplished using a Vibrating Sample Magnetometer (VSM, EV9 Micro sense Lakeshore VSM) at RT.

XRD is one of the most valuable tools for the observation of the atomic crystal structure as well as for the determination of unit cell parameters. The crystallographic structures of the synthesized bulk ceramic samples of Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (x = 0.00, 0.05, 0.10, 0.15, and 0.20), sintered at temperatures of 1200 °C and 1250 °C, were examined by the XRD analysis and are shown in Figs. 1(a) and 1(b). As observed from Fig. 1(a), the XRD patterns of all the Mn substituted BCSMTO ceramic samples sintered at a temperature of 1200 °C clearly illustrate that the diffracted planes reflect a very lower intensity for x = 0.00 and x = 0.15 compositions. This demonstrates that the sintering densification may not have finished for undoped and 15% Mn-doped samples at a temperature of 1200 °C. On the other hand, the XRD planes for all the samples sintered at a temperature of 1250 °C show a comparatively better-crystallized perovskite structure with a negligible amount of impurity phases than the samples sintered at a sintering temperature of 1200 °C. For this reason, in our research work, we have carried out all our characterizations for the compositions sintered at a temperature of 1250 °C. The XRD planes of Ca2+, Sr2+, and Mn2+ substituted BCSMTO ceramics are in accordance with the previously reported works.19,20

FIG. 1.

XRD patterns of Mn-doped Ba0.4Ca0.4Sr0.2MnxTi1−xO3 specimens sintered at temperatures of (a) 1200 °C and (b) 1250 °C.

FIG. 1.

XRD patterns of Mn-doped Ba0.4Ca0.4Sr0.2MnxTi1−xO3 specimens sintered at temperatures of (a) 1200 °C and (b) 1250 °C.

Close modal

The distinguishable diffracted peaks indexed in Fig. 1(b) indicate that the crystal structures for undoped, 5%, 10%, and 15% Mn-doped specimens are of cubic perovskites without any identified secondary phase. The creation of single-phase perovskite structures of these compositions (x = 0.00, 0.05, 0.10, and 0.15) established the fact that A-site doping with Ca2+ and Sr2+ ions, and B-site doping with the Mn2+ ion may have almost integrated into the perovskite lattice of BTO. The XRD patterns reflect that all the samples except the 20% Mn-doped composition show single peaks around 32° and 46°. Interestingly for the 20% Mn-doped sample, a gradual emergence of splitting of (101)/(110) and (002)/(200) is distinctly noticeable, which implies the existence of tetragonal phase deformation similar to that of BaTiO3.20 However, according to Fig. 1(b), some extra peaks marked by “*” corresponding to SrTiO3 (JCPDS 01-079-0175) are distinctly noticeable in the x = 0.20 composition. These developments of secondary peaks during the solid-state synthesis are expected in cation substituted BTO materials due to its chemical kinetics of formation. To closely examine the peak positions with different Mn concentrations, the magnified XRD patterns located at 2θ = 31° to 34°, representing the (101) plane, are shown in the inset of Fig. 1(b). This figure displays that the peak positions for 5%, 15%, and 20% Mn-doped samples move toward the higher diffraction angle, which suggests that the lattice parameter may have decreased with the Mn substitution due to the shrinkage of unit cell volumes.

Taking into account the most intense peak (101), the lattice parameter “a” of each composition is calculated using the formula: a=dhklh2+k2+l2 [where dhkl is the d-spacing between the crystal planes and (hkl) are the Miller indices] and the values are shown in Table I. In order to seek the evidence of such findings, the accurate lattice constant values for individual samples were also calculated using the Nelson–Riley extrapolation method.21 The calculated values are shown in Table I as well, and as a representative among all the compositions here only for x = 0.05, the “a” parameters related to all the diffracted planes are plotted in contrast to the Nelson–Riley function, F(θ), and displayed in Fig. 2(a). The lattice constant “a” values were determined from the interception of the straight lines, which were obtained by the fit linear process for x = 0.0–0.20. There exists a correlation between the lattice parameter and the ionic radius; the lattice constant decreases proportionally with the amplified ionic radius.21 Therefore, it is expected that the lattice parameter “a” of the ceramic samples would decrease with the addition of a divalent Mn2+ ion in the cubic perovskite structure. After analyzing the calculated lattice parameter values from XRD data and from the Nelson–Riley function, it can be summarized that in both cases lattice parameters are found to show a decrement with the addition of Mn contents up to 15% (as observed in Table I). As the ionic radii of Sr2+ (1.44 Å) and Ca2+ (1.00 Å) are smaller than that of Ba2+ (1.61 Å),22 this reduced lattice constant denotes that Sr2+ and Ca2+ have entered the A-site of the perovskite structure by replacing Ba2+. Again, the B-site substitution of the Mn2+ ion with a little bit higher ionic radius of 0.645 Å than the Ti4+ ion (ionic radius of 0.605 Å)23 may have a great impact on the decrement of this lattice parameter. Moreover, it was mentioned in a previous study that when a larger Mn2+ ion enters the B-site of the crystal, the unit cell starts expanding, whereas the overall cubic symmetry is preserved.24 

TABLE I.

Structural entities of the samples, i.e., the lattice parameter, density, crystalline size, and porosity of Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (x = 0.00–0.20).

LatticeLattice
parameterparameter from
Contentfrom XRDNelson–RileyX-ray densityBulk densityCrystallinePorosity
Xa (Å)function a (Å)ρx (g/cm3)ρB (g/cm3)size D (nm)P (%)
0.00 3.94 3.97 5.36 2.72 0.305 46.24 
0.05 3.92 3.96 5.45 2.93 0.458 49.25 
0.10 3.87 3.89 5.66 3.74 0.458 33.92 
0.15 3.88 3.90 5.79 3.45 0.524 37.06 
0.20 3.90 3.92 5.89 3.19 0.612 45.84 
LatticeLattice
parameterparameter from
Contentfrom XRDNelson–RileyX-ray densityBulk densityCrystallinePorosity
Xa (Å)function a (Å)ρx (g/cm3)ρB (g/cm3)size D (nm)P (%)
0.00 3.94 3.97 5.36 2.72 0.305 46.24 
0.05 3.92 3.96 5.45 2.93 0.458 49.25 
0.10 3.87 3.89 5.66 3.74 0.458 33.92 
0.15 3.88 3.90 5.79 3.45 0.524 37.06 
0.20 3.90 3.92 5.89 3.19 0.612 45.84 
FIG. 2.

(a) Variation of the lattice parameter with the Nelson–Riley function; variation of (b) x-ray density and bulk density, and (c) bulk density and porosity with the Mn concentration.

FIG. 2.

(a) Variation of the lattice parameter with the Nelson–Riley function; variation of (b) x-ray density and bulk density, and (c) bulk density and porosity with the Mn concentration.

Close modal

Bulk densities and x-ray densities of samples with different doping concentrations are listed in Table I and also displayed in Fig. 2(b). The density of ceramic compositions acts as an important part in controlling the magnetic properties. It has been proposed that higher permeability could be obtained by increasing the density of ceramics. Here, the bulk density (ρB) was evaluated from the ratio of mass and volume, whereas the x-ray density (ρx) was measured using the molecular weight and lattice constant for all the compositions. The results of bulk density measurements indicate that the increase in the Mn concentration up to 10% can effectively promote densification, while excessive Mn hinders the densification and reduces the bulk density. The difference in atomic weights of the initial and substituted cations [the atomic weights of Mn2+ (54.938 amu) > Ti4+ (47.867 amu)]25 may have increased ρB of 5% and 10% Mn-doped samples. On the other hand, a significant enhancement in the x-ray density is observed with the increasing substitution of Mn2+ contents. According to the sum rule, the density of the mixture is the weighted sum of the individual density.26 The increase in ρx from x = 0.00 to x = 0.20 indicates that the addition of a lightly densified spinal phase might have increased the composite weights. From Fig. 2(b), it is also obvious that the bulk densities are lower in magnitude than the x-ray densities. Sometimes, during the sintering procedure, the intergranular/intragranular porosity can be formed and developed, which may guide to a distinction in bonding among the elements.27 Thus, it starts to change the inter-atomic distances and therefore makes a contribution to enlarging the unit cell volume and diminishing the density.

Figure 2(c) shows the effect of Mn addition on the bulk density and porosity of the samples. It is clear that the bulk density increases for x = 0.0–0.10 and slightly decreases for contents x = 0.15 and 0.20. The bulk density is governed by the porosities of the sintered samples. Here, the porosity values have displayed a significant change with respect to the increment of Mn contents. The composition with a concentration of x = 0.10 has the highest density of 3.74 g/cm3 and the lowest porosity of 33.92%. During the sintering densification, discontinuous grain growth starts stimulating and thus, the inter-granular porosity is diminished. When the rate of the grain growth is very high, the very fast-moving grain boundaries help to suppress the pores and therefore, no pores can be captured inside the grains.28 This contributes to the increase in bulk density, which causes a decrease in porosity. Previous investigations revealed that the porosity starts to decrease with the inflation of particle size.29,30 Therefore, the decreasing porosity for the 10% Mn-doped sample also implicates that the size of crystallites may have enhanced due to the minimization of the grain boundaries of the crystallite. Here, the crystalline size “D” for all the specimens was calculated using Scherrer’s formula and the variation of crystalline size as a function of Mn contents is shown in Table I. Hence, our present work illustrates that the porosity changes inversely with the bulk density by means of the addition of Mn contents, which is analogous to the other reported result.31 The lattice parameter, crystalline size, density, and porosity of different samples calculated using XRD data are presented in Table I.

The microstructure, surface morphology, and hence the range of average grain size of all the undoped and Mn-doped samples were measured by SEM imaging. The SEM micrographs and their respective histograms are shown in Figs. 3(a)3(e). It is clear from these images that the substitution with 5%, 10%, and 15% Mn within BCSTO ceramics degraded its surface morphology with increased grain size ranging from 0.2 µm to 1.0 µm. The SEM images also expose that the microstructures of these compounds are not dense enough and they contain more disorderly and unsystematic shaped grains, implicating the polycrystalline nature of these compositions. Since all the synthesized Ba0.4Ca0.4Sr0.2MnxTi1−xO3 samples are subjected to the 4 h long sintering process, the inhibition of the movement of the grain boundary due to the lack of activation energy may have hindered the possibility of densification without grain growth while keeping the diffusion of the grain boundary operative.32,33 Comparing Fig. 3(e) with the other images, it can be clearly seen that the amount of small grains increases with 20% Mn substitution and the range of grain size of this specific composition lies within 0.2–0.6 µm. As reported previously, the grain growth basically depends on the sintering temperature, oxygen vacancy concentration, and ion diffusion rate.34 It is suggested that a proper sintering temperature causes the pores to minimize and hence reduces the grain size, which helps to suppress the grain boundaries.27 However, the bond enthalpy of Ti–O (672.4 ± 9.2 kJ/mol) is stronger than that of Mn–O (362 ± 25 kJ/mol),35 which demonstrates that the development of Mn–O bonding requires higher energy. Moreover, it is suggested that the rate of nucleation N and the rate of grain growth G strongly influence the grain size, whereas both rates depend heavily on the composition and temperature.36 If G is less than N, the grain size would be small. Therefore, the suppressed grain size of the 20% Mn substituted sample might be interpreted due to the reduction of oxygen vacancies, which lessens the oxygen ion movements and consequently hampers the grain growth rate.

FIG. 3.

SEM micrographs of Ba0.4Ca0.4Sr0.2MnxTi1−xO3 [(a) x = 0.00, (b) x = 0.05, (c) x = 0.10, (d) x = 0.15, and (e) x = 0.20] ceramics sintered at a temperature of 1250 °C. Inset: respective histograms of images (a)–(e).

FIG. 3.

SEM micrographs of Ba0.4Ca0.4Sr0.2MnxTi1−xO3 [(a) x = 0.00, (b) x = 0.05, (c) x = 0.10, (d) x = 0.15, and (e) x = 0.20] ceramics sintered at a temperature of 1250 °C. Inset: respective histograms of images (a)–(e).

Close modal

Figures 4(a) and 4(b) illustrate the variation of the real part (ε′) and imaginary part (ε″) of complex dielectric constant (ε* = ε′ − jε″) with the increase in the frequency for Ba0.4Ca0.4Sr0.2MnxTi1−xO3 ceramics in the order of 20 Hz–1 MHz. Figure 4(a) shows that the real part of dielectric constant (ε′) of the 5% Mn substituted BCSMTO ceramic sample is maximum at lower frequencies, which decreases sharply with the raising frequency up to about 10 kHz and then becomes almost constant at higher frequencies. All the other Mn undoped and Mn-doped ceramic samples also demonstrate slight dispersion at lower frequencies and a frequency independent attitude on a broad scale of high frequencies. This observed lower frequency dispersion of ε′ is common in dielectric and ferroelectric materials and can be described in accordance with space charge polarization, which emerges because of the inhomogeneous dielectric configuration of the material as discussed in the two-layer model of Maxwell37 and Wagner.38 However, the decrease in dielectric constants with growing frequency can be interpreted by the phenomenon of dipole relaxation, according to which the dipoles are able to pursue the frequency of the activated electric field at lower frequencies.39 Again, Fig. 4(b) shows that the imaginary part of ε″ also decreases with the spreading frequency, and following a certain frequency range, it remains unchanged, which points to the decrement in the amount of hopping electrons. The Verway–de-Boer hopping mechanism40 explains that the bouncing of electrons among the ions of the same material that are dispersed arbitrarily over crystallographic counterpart lattice sites enhances the electronic conduction in perovskites. Formation of Mn2+ ions during the sintering of BCSMTO perovskites allows an enhancement of electron hopping between Mn ions in +2 and +3 valence states. The electronic exchange between Mn2+ and Mn3+ ions results in local displacement of charges according to the applied field, which is responsible for the polarization in perovskites41 and also enhances the dielectric loss. In general, the polycrystalline ceramics contain perfectly conducting grains that are isolated by insulating grain boundaries. So, the displacement of charge carriers occurs when a field is applied and if the resistances are large enough the charge carriers tend to align themselves at grain boundaries. As a result, space charge polarization is developed, which is governed by the free charge carriers and this behavior might result in the large values of dielectric constant at lower frequencies. On the other hand, at higher frequencies, they might not get sufficient time to build up and undergo the relaxation process. As a consequence, with the expanding frequency of the applied field, the electrons converse their path more frequently. So, the possibility of electrons joining the grain boundary gets reduced, which causes polarizations to cease and therefore, ε′ becomes almost static.

FIG. 4.

Variation of dielectric constant (a) real part ε′ and (b) imaginary part ε″ as a function of frequency for Ba0.4Ca0.4Sr0.2MnxTi1−xO3 ceramic samples. Insets: the behavior of ε′ and ε″ as a function of frequency in the higher frequency range.

FIG. 4.

Variation of dielectric constant (a) real part ε′ and (b) imaginary part ε″ as a function of frequency for Ba0.4Ca0.4Sr0.2MnxTi1−xO3 ceramic samples. Insets: the behavior of ε′ and ε″ as a function of frequency in the higher frequency range.

Close modal

The insets of Figs. 4(a) and 4(b) display the very interesting behavior of ε′ and ε″ for Ba0.4Ca0.4Sr0.2MnxTi1−xO3 ceramic samples in the higher frequency range. As it is mentioned in some early studies,42,43 above the infrared region (i.e., in the higher frequency range), just the electronic polarizability survives. This is because in the higher frequency range, the light electrons are still capable of following the field. At even higher frequencies, the electronic contribution dies out soon since the electrons are too weighty to follow the very rapidly oscillating field. For 20 Hz to 1 MHz frequencies, the dielectric constant ε′ behaves as the familiar static dielectric function. However, at the opposite end of the spectrum >1 MHz, it is clear from these figures that all the samples show damping. It is because in that range of frequency, the oscillation of the field is too high for the ions to follow. According to the insets of Figs. 4(a) and 4(b), dielectric constants for all the samples are negative in the range of 2.4 × 107 Hz–5.5 × 107 Hz. The wave in this frequency range does not propagate through the crystal and is defined as the forbidden gap.42 

Figure 5 shows the fluctuation of dielectric loss, tan δ as a function of frequency. At the low-frequency domain, the loss factor tan δ for all the samples is higher, but then starts to decrease rapidly at higher frequencies in accordance with Koop’s phenomenological theory.44 It is observed that the 10% Mn-doped ceramic sample exhibits a minimum dielectric loss. The large values of tan δ at lower frequencies indicate that the frequency of the applied ac field is much lower compared to the hopping frequency of electrons between Mn2+ and Mn3+ ions, and therefore, the electrons follow the field, which maximizes the loss. On the contrary, as the hopping frequency of the electron exchange between Mn2+ and Mn3+ ions at the higher frequency domain cannot chase the applied field above the specific critical frequency, the loss is diminished. Interestingly, in the higher frequency region, there is a tendency of increment of loss tangent for the 5% Mn substituted sample. This increment is owing to the relaxation function that makes tan δ much greater in the higher frequency range. Again, a large number of grains near the grain boundaries underneath the pressure of an electric field cause a localized aggregation of charge, which ensures the existence of space charge polarization.45 Generally, tan δ has very small values at higher frequencies, which indicate that dipoles with small effective masses (i.e., electrons and ferroelectric domains) mainly contribute to the dielectric constant instead of charge defects with large effective masses (i.e., oxygen vacancies). Because of the low (<0.5) loss tangent values at higher frequencies, the composition with x = 0.10 might have potential applications in high-frequency microwave devices.

FIG. 5.

Variation of loss tangent (tan δ) as a function of frequency for Ba0.4Ca0.4Sr0.2MnxTi1−xO3 ceramic samples at room temperature.

FIG. 5.

Variation of loss tangent (tan δ) as a function of frequency for Ba0.4Ca0.4Sr0.2MnxTi1−xO3 ceramic samples at room temperature.

Close modal

The variation of the real part of initial permeability (μi) with the increase in the frequency for different compositions is shown in Fig. 6(a). Permeability is a measure of the capability of a material to maintain the formation of the magnetic field within itself. The complex initial permeability can be represented as μi*=μiμi. Here, μi and μi represent real and imaginary parts of complex initial permeability. The real part of permeability indicates the reserved energy by expressing magnetization in aspect with the varying magnetic field. The values of μi indicate the squandering of energy by expressing B out of phase with H. It is clearly observed from Fig. 6(a) that in the frequency scale of 20 kHz–1 MHz, μi for all the samples remains almost constant and after a certain time, it starts to decrease slowly. Again, for all the compositions with the further expansion of frequency (at higher frequencies >108 Hz), a sudden increase in μi has been spotted, which indicates the beginning of resonance. When the applied magnetic field frequency is equal to the Larmor precession of the electron spins, resonance occurs and the energy from the field is transported toward the system for aligning the magnetic dipoles.46 This observation is in agreement with that of the Globus model. According to this model, the relation of the resonance frequency fr with the initial permeability can be expressed as

(μi1)12fr=constant.

In accordance with the above relation, the higher the permeability, the lower will be the resonance frequency and vice versa. According to the Globus and Duplex model,47 the static permeability and its product are micro-structurally dependent. Moreover, it was confirmed that the permeability of polycrystalline materials can be influenced by the magnetizing mechanisms known as domain wall motion as well as spin rotation. At lower frequencies, domain wall relaxation is usually predominating, while in the higher frequency region, domain walls become unable to follow the AC excitation field and therefore, spin rotation remains as the only active process for magnetization.48 It is worthwhile to mention that the domain wall motion may be influenced by various factors, such as grain size, intragranular porosity, domain wall displacement, and domain wall bulging.49 At higher frequencies, nonmagnetic impurities between intragranular pores and grains act as a catalyst and gradually reduce the spin and domain wall motion. Hence, the permeability decreases and the loss increases. However, the flat region of μi over a large frequency range gives the compositional stability and quality of the prepared sample and is known as the zone of utility.50 The variation of μi at 30 kHz frequency as a function of Mn contents is displayed in Fig. 6(b). The figure shows that μi decreases with the increase in Mn substitution for x = 0.05, but with the further addition of the Mn content from x = 0.10 to x = 0.15, it starts to rise. Then, a slight drop off in μi is observed for x = 0.20. Here, the highest value of permeability is obtained for x = 015. The improvement in μi for this specific composition is attributed to the decreased magnetic anisotropy since sometimes the sintering temperature causes a decrement in magnetic anisotropy by reducing the internal stresses and crystal anisotropy. Moreover, previous investigations demonstrated that the higher the density and grain size, the greater continuity of the grain to follow the magnetic flux, which leads to higher permeability.51 That is, the permeability is lower, when the grain size is reduced but the stability of permeability is higher, since the presence of small grain size interferes with the domain wall motion. It was observed from the SEM micrographs that the grain size increases from x = 0.0 to x = 0.15 and the enhancement in permeability for x = 0.15 has occurred because the microstructure is homogeneous with uniform size distribution.

FIG. 6.

Variation of (a) real part of initial permeability, μi with frequency, (b) μi with Mn concentration, and (c) loss factor, tan δM as a function of frequency for Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (where x = 0.00–0.20) ceramic samples.

FIG. 6.

Variation of (a) real part of initial permeability, μi with frequency, (b) μi with Mn concentration, and (c) loss factor, tan δM as a function of frequency for Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (where x = 0.00–0.20) ceramic samples.

Close modal

Figure 6(c) shows the deviation of magnetic loss tan δM as a function of frequency. The ratio of μi to μi is the loss tangent, and it represents the measure of inefficiency of the magnetic system. Higher values of μi and lower values of loss tangent are required for high-frequency magnetic applications. The loss occurs because the domain walls start lagging to follow the applied alternating field, which is ascribed to the imperfections in the lattice.50 It should be mentioned that tan δM is susceptible to many aspects, such as the mobility of charge carriers, porosity, domain defects, concentration of dipoles, hysteresis loss, eddy current loss, and residual loss.52 As observed from Fig. 6(c), the values of tan δM for all the synthesized samples initially decrease exponentially with the increase in the frequency (≤105 Hz) and then become almost constant up to 100 MHz. The decrease in tan δM with the increase in the frequency is caused since beyond some critical frequency, the applied electric field cannot be followed by the domain wall motion. The corresponding figure shows that the loss increases for 5% Mn addition and then gradually decreases with the further increase in the Mn content. The total magnetic loss contains hysteresis losses, eddy current losses, and residual losses. At lower frequencies, hysteresis losses vanish and therefore, the magnetic loss observed here in this investigation is due to the eddy current losses and residual losses.

Figure 7 reveals that Q for all the ceramic samples shows a noteworthy enlargement with the increment of frequency, which is in agreement with the equation Q = μitan δM. As it is mentioned before, the lag of domain wall movements occurs since the highly fluctuating applied magnetic field is very difficult to follow, so the loss is increased and is referred to as imperfections in the lattice.50 Here, at higher frequencies, Q values show the beginning of a resonance peak. The highest Q values are obtained for the 10% Mn-doped sample, whereas the lowest loss (tan δM) values [from Fig. 6(c)] were also observed for this specific composition. This is most likely due to the development of fewer imperfections, less defects, and higher bulk density compared to the other samples that are responsible for high Q. From the XRD analysis, we have already investigated that the 10% Mn-doped sample has the lowest porosity, but the maximum bulk density values among all the samples. The relative quality factor (RQF) is often used as an aid of measurement of performance. Here, all the compositions have a high RQF in the high-frequency range. Since the composition of low loss and high RQF is suitable for high-frequency magnetic applications, the 10% Mn-doped ceramic sample can be claimed to be the most suitable because of its comparatively large value of RQF with minimum tan δM.

FIG. 7.

The variation of relative quality factor with the increase in the frequency in the case of Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (where x = 0.00–0.20) samples.

FIG. 7.

The variation of relative quality factor with the increase in the frequency in the case of Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (where x = 0.00–0.20) samples.

Close modal

To survey the RT magnetic performance, M–H hysteresis loops for different Ba0.4Ca0.4Sr0.2 MnxTi1−xO3 (x = 0.0–0.20) compositions were obtained with the highest applied field up to 10 kOe. It was mentioned in a previous study that the bulk BaTiO3 ceramic reflects a diamagnetic signature due to the existence of the d0 structure of the 3d transition metal (Ti).53 In our investigation, as shown in Fig. 8, with the substitution of Ca and Sr in BTO, a well-saturated hysteresis loop is clearly observed, which indicates the presence of soft ferromagnetism in the Mn undoped BCSTO crystal. However, the M–H loops for Mn-doped BCSMTO samples clearly demonstrate that they are not saturated within the range of measured filed. The un-saturating tendency of the M–H loops for 5%, 15%, and 20% Mn-doped samples has probably emerged from the effect of mixed PM and AFM phases.54 Moreover, for the 10% Mn-doped sample, the domination of diamagnetic actions at upper fields may have resulted due to the inhibition of exchange interaction energy by the magnetic energy.2 Here, the magnetic parameters: coercive field, saturation magnetization, remanent magnetization, and anisotropy constant of the samples are calculated from the hysteresis loops and presented in Table II. The coercivity (Hc) is a measure of the magnetic field strength desired for reducing anisotropy to twist the magnetic moments. It is suggested that the emergence of large Hc as observed in the 10% Mn-doped sample might be reasoned by various factors, for instance, microstrain, magnetocrystallinity, magnetic domain size, and size distribution.55 The coercivity is in general linearly dependent on the magnetic anisotropy and demagnetization factor. Furthermore, increased grain boundaries can lead to a rise in pinning sites for the spins and therefore enhance the coercivity. From Table II, it is clearly noticeable that for the sample with x = 0.00, the saturation magnetization is found to be 0.077 emu/g and the remanent magnetization is 0.016 emu/g, which are highest among all the samples. The saturation magnetization afterward starts to decrease as the concentration of the Mn2+ ion increases from x = 0.00–0.10 and then again increases for contents x = 0.15–0.20. These observations indicate that the magnetization in Mn-doped samples is hard to realize due to the presence of the non-magnetic BaTiO3 perovskite, in which the interaction of magnetic poles on the magnetic particles is damped. This also demonstrates that the spontaneous magnetization of the composites originates from unbalanced antiparallel spin, which leads to net spins other than those aroused from the structural distortion. In a previous investigation, it was suggested that the double-exchange interactions between Mn2+ and Mn3+ ions in tetragonal BaTiO3 are the origin of ferromagnetism.13 The development of nonmagnetic hexagonal BaTiO3 at RT and the introduction of an antiferromagnetic Mn3+–Mn3+ pair dilute the ferromagnetic interaction, thus resulting in a decrement of Ms in these samples. Moreover, the weakening permeability for 5% and 10% concentrations as observed from Fig. 6(a) [the real part of permeability (μi) vs frequency graph] is also responsible for the reduction of saturation magnetization of these specific compositions. Notably, Ti4+ ions have a 0 µB magnetic moment and Mn2+ ions have a 5.9 µB magnetic moment. So, the replacement of non-magnetic Ti4+ ions at the B-site through the magnetic Mn2+ ions starts to increase the B-site magnetic moment. As a consequence, the net magnetization increases due to the suppressed A–B interactions, which results in the further increment of Ms in 15% and 20% Mn-doped samples.

FIG. 8.

Variation of magnetization (M) as a function of applied magnetic field (H) of Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (where x = 0.00–0.20) compositions at RT. The inset represents the magnified M–H curves.

FIG. 8.

Variation of magnetization (M) as a function of applied magnetic field (H) of Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (where x = 0.00–0.20) compositions at RT. The inset represents the magnified M–H curves.

Close modal
TABLE II.

The saturation magnetization, remanent magnetization, coercivity, and anisotropy constant of Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (where x = 0.00, 0.05, 0.10, 0.15, and 0.20) samples.

SaturationRemanent
ConcentrationmagnetizationmagnetizationCoercivityAnisotropy
xMs (emu/g)Mr (emu/g)Hc (Oe)constant Ka
0.00 0.077 0.015 208.208 7.97 
0.05 0.072 0.011 217.984 7.81 
0.10 0.045 0.012 280.139 6.37 
0.15 0.070 0.012 248.319 8.73 
0.20 0.073 0.011 270.817 9.89 
SaturationRemanent
ConcentrationmagnetizationmagnetizationCoercivityAnisotropy
xMs (emu/g)Mr (emu/g)Hc (Oe)constant Ka
0.00 0.077 0.015 208.208 7.97 
0.05 0.072 0.011 217.984 7.81 
0.10 0.045 0.012 280.139 6.37 
0.15 0.070 0.012 248.319 8.73 
0.20 0.073 0.011 270.817 9.89 

Again, from Table II, it is observed that the remanent magnetization also decreases with the increased Mn contents from x = 0.00–0.10, but after that it increases for x = 0.15 and then slightly decreases for x = 0.20. In this study, the anisotropy constants are also calculated using the formula K = HcMs2 and presented in Table II. It also decreases with the addition of the Mn content (x = 0.00–0.10) and increases with the further increment of the Mn content (x = 0.15–0.20). Magnetocrystalline anisotropy has a great influence on the industrial uses of ferromagnetic materials. Materials with high magnetic anisotropy usually have large coercivity, which makes them hard to demagnetize. These materials are called hard ferromagnetic materials and are used to make permanent magnets.

Mn2+ substituted Ba0.4Ca0.4Sr0.2MnxTi1−xO3 (x = 0.00, 0.05, 0.10, 0.15, and 0.20) ceramic compositions were formulated by the solid-state reaction procedure, and their structural configuration, surface morphology, and electrical and magnetic properties have been studied at the sintering temperature of 1250 °C. With 20% Mn addition, the patterns of XRD peaks exhibited a structural phase transformation from cubic to tetragonal. The calculated lattice constant “a” values found from both XRD data and fitting the Nelson–Riley function display noticeable reduction with an increase in Mn contents for x = 0.00–0.10 and then start to increase with further Mn addition. The decrease in porosity with an increase in the Mn content (up to 10%) reflects better crystallization. The SEM micrographs confirmed the highly pored morphology as well as nonuniform grain size distribution but with visible grain boundaries. Both the real and imaginary parts of complex dielectric constant have decreased at lower frequencies, and the dielectric loss, which is significantly decreased at higher frequency, for the 10% Mn-doped composite can be applied for high-frequency dielectric operations. The real part (μi) of initial permeability for all the samples displays the existence of resonance peaks at higher frequency. The RQF, which is found maximum for x = 0.10 substitution at higher frequency, signifies the high-frequency magnetic applications of this specific composition. The degradation of permeability up to 10% Mn concentration is the consequence of diminution in saturation magnetization, which is caused since the substitution of Mn2+ ions significantly weakened the A–B interaction. The anisotropy constant is found to be increased for 15% and 20% Mn substitution, and these two compositions also show higher coercivities. Since the materials with greater magnetic anisotropy and higher coercivity usually have great potential for industrial uses as ferromagnetic materials, 15% and 20% Mn-doped ceramics can be used to make permanent magnets.

The authors greatly acknowledge the Department of Physics, University of Dhaka, Dhaka. The authors are highly grateful to the Materials Science Division, Atomic Energy Center, Dhaka, and the Center for Advanced Research of Sciences, University of Dhaka, Dhaka, Bangladesh.

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