Effect of co-substitutions (Ca–Mn) on structural, electrical and magnetic characteristics of bismuth ferrite

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CERAMICS INTERNATIONAL

Ceramics International 41 (2015) 9078–9087 www.elsevier.com/locate/ceramint

Effect of co-substitutions (Ca–Mn) on structural, electrical and magnetic characteristics of bismuth ferrite Jyoshna Routa,n, R.N.P. Choudharya, H.B. Sharmab, S.R. Shannigrahic a

Department of Physics, Institute of Technical Education and Research, SOA University, Bhubaneswar 751030, India b Department of Physics, Manipur University, Canchipur 795003, India c Institute of Material Institute of Materials Research and Engineering, AnSTAR (Agency for Science, Technology and Research), 3 Research Link, Singapore 117602, Singapore Received 20 December 2014; received in revised form 19 March 2015; accepted 27 March 2015 Available online 7 April 2015

Abstract The polycrystalline samples of (Bi1
xCax)(Fe1
xMnx)O3 with x ¼0, 0.05, 0.10, and 0.15 mol fraction were prepared by synthesizing through initial mixing of precursors by high-energy ball milling, and subsequent thermal treatment of the milled powders. Study of basic crystal parameters using x-ray diffraction technique of the calcined samples suggests that all the materials are crystallized in the distorted-perovskite (rhombohedral) structure which was supplemented and supported by their FTIR studies. The dielectric properties of the materials were investigated as a function of temperature and frequency and the temperature dependent dielectric constant and loss tangent of the modified system vary with substitution concentration (x) and frequency. Detailed studies of impedance exhibit that the electrical properties of the prepared systems are strongly dependent on temperature and frequency; and thus show a good correlation with their microstructures (i.e. bulk, grain boundary, etc.). Further, it is interesting to see the contributions of bulk, grain boundary and electrode in the impedance plots. The decrease in value of bulk resistance on increasing temperature suggests the existence of negative temperature co-efficient of resistance (NTCR) behavior in the materials. Room temperature M–H hysteresis measurements reveal that magnetization increases due to the structural distortion, and partial destruction of spin cycloid is caused by co-doping in BiFeO3 ceramics. Unlike magnetic response, the ferroelectric properties of the samples were found be enhanced on increasing mole fraction of x in BiFeO3. & 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Keywords: High energy ball milling; X-ray diffraction; Magnetization; Ferroelectric

1. Introduction In some insulating materials, an external magnetic field can induce electric polarization and an external electric field can induce changes in magnetization; such a phenomenon is known as the magnetoelectric (ME) effect and materials exhibiting this effect are called magnetoelectrics [1,2]. The coexistence of ferroelectric and ferromagnetic properties induces coupling between the magnetic and ferroelectric domains [3], and thus those materials could be used either as ferroelectrics or as magnetic materials and both in devices like ferroelectric capacitors, magnetic data-storage media, transducers, spintronics, actuators, sensors, and optoelectronic n

Corresponding author. Tel.: þ91 9437294494. E-mail address: [email protected] (J. Rout).

http://dx.doi.org/10.1016/j.ceramint.2015.03.283 0272-8842/& 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

devices [4,5]. However, the number of materials exhibiting magnetoelectricity is limited, and there is a strong dependence on specific conditions of the synthesis process. Among those materials, BiFeO3 (BFO) occupies a center stage due to its room temperature multiferroic property. The room temperature phase of BFO is classified as rhombohedral (R3c space group) with high ferroelectric Curie point (TC E830 1C), high anti-ferromagnetic (AFM) Neel temperature (TN E370 1C) and narrow band gap (2.3 eV) [6,7]. The major drawback of BFO is its large leakage current, which arises due to the existence of small amounts of Fe þ 2 and oxygen vacancies [8], and due to difficulty in synthesizing single-phase BFO in its narrow temperature range of phase stabilization. When BFO is fabricated in the ceramic form by a solid-state reaction process, single-phase material is difficult to obtain. The kinetics of phase formation often leads to formation of

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some secondary or impurity phases, such as Bi2Fe4O9, Bi25FeO40, and Bi2O3 because of the volatilization of some reactants and phase decompositions at high temperature [9]. In order to obtain the pure phase and enhance the multiferroic properties (i.e., ferroelectric, ferromagnetic or/and ferroelastic) of BFO, many steps have already been taken such as doping/substitution with rare earth elements [10,11], fabricating composites [12,13], and or forming solid solutions with other ferroelectric materials [14,15]. The strategy of substitution at the Bi-site/Fe-site and/or both at Bi/Fe site has been investigated widely. Some of the trivalent Bi-site substituent commonly used are La þ 3, Gd þ 3, Nd þ 3, and Dy þ 3; more recently, divalent Bi-site substituent's (such as Sr þ 2, Ca þ 2, Ba þ 2, and Pb þ 2) and Fe-site substituent's (such as V þ 5, Cr þ 5, Mn þ 4, and Ti þ 4) are also investigated [16–18]. These substitutions have been found to be effective for improving the ferroelectric or/and the magnetic properties. Since the ferroelectricity and magnetism in BFO are caused by the lone pair electrons of Bi þ 3 ions at A-sites and partially filled “d” orbital electrons of Fe þ 3 ions at Fe-sites respectively. The divalent ion substitution results in charge imbalance creating oxygen vacancies and/or charge transformation of a fraction of Fe þ 3 to Fe þ 2/Fe þ 4. It has been reported that quadrivalent ions (Mn þ 4, Ti þ 4, and Zr þ 4) substitution at Fe-site reduces oxygen vacancies [19,18]. Among them Mn substitution at Fe-sites not only reduces the leakage current of BFO, but also may produce double exchange (Mn þ 3– O
2–Mn þ 4), leading to enhanced magnetism of BFO [20]. Based on the literature survey, we have synthesized new samples in anticipation of getting better properties in terms of reduced leakage current and enhanced ferromagnetic properties. In this paper we report on synthesis and characterization (structural, electrical and magnetic properties) of Bi1
xCaxFe1
xMnxO3 (x¼ 0.00, 0.05, 0.10, and 0.15) samples. 2. Experimental The co-substituted (Ca, Mn) samples of BiFeO3 (i.e. Bi1
xCaxFe1
xMnxO3 (x¼ 0.00, 0.05, 0.10, and 0.15) (BCFMO) were prepared by high-energy ball-milling method followed by thermal treatment. The stoichiometric proportion of high-purity (4 99.9%) ingredients: Bi2O3 (with 2 mol% extra), Fe2O3, CaCO3 and MnO2 (all from M/S Loba Chemie Co, India) were weighed carefully and ball-milled in a zirconium-grinding jar of 250 ml capacity with 10 mm zirconium balls in a planetary ball mill (M/s Retsch PM 100, Germany). The wet milling (in toluene medium) was carried out with ball to powder weight ratio of 10:1 and speed 400 rpm for 30 h. After ball milling, the mixture was dried at 100 1C for 24 h, and then all the samples were calcined at 700 1C. The basic crystal structure of the samples at room temperature was analyzed using x-ray diffraction (XRD) data collected by x-ray powder diffractometer (PANalytical, X'Pert-PRO) with Cu-Kα radiation in a wide range of Bragg angle θ (20o r 2θ r 60o) at room temperature. Fourier transform infrared spectra (FT-IR) were recorded at room temperature using the Perkin-Elmer (model L1280034) in the range of 400–2000 cm
1. The calcined powders were mixed thoroughly with polyvinyl alcohol, PVA (binder) solution, and then pressed into disk-

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shape pellets (size ¼ 10 mm diameter and 1–2 mm thickness) at the pressure of 5  106 kg/m2. The pellets of pure BFO were sintered at 7801C for 2 h, and other pellets with x ¼ 0.05, 0.10, and 0.15 were sintered at 850 1C for 2 h in air atmosphere. The surface morphology of the samples was obtained by a scanning electron microscope (SEM, FEI, QUANTA-250). A sintered pellet was polished with fine emery paper; silver contacts were made on the surface of pellets for all electrical measurements. The dielectric and electrical parameters were measured using a phase sensitive meter (PSM 1735, N4L) at room temperature. The electric field dependence of polarization (P–E hysteresis loop) of all the samples was recorded using P–E loop tracer (M/s Marine India Co.). Magnetic measurements (M–H loop) of the samples were carried out using the physical property measurement system (PPMS; Quantum Design) at room temperature with a maximum magnetic field of 10 kOe. 3. Results and discussion 3.1. Structural and micro-structural characteristics Fig. 1 shows room temperature XRD patterns of BCFMO powders calcined at 700 1C. All the peaks of the XRD patterns (except few peaks of impurity phase) were indexed in different crystal systems with observed peak position (i.e. Bragg angle (2θ) and intensity) using a standard computer software or program package “POWDMULT” [21]. From the best agreement between observed (obs) and calculated (cal) values of inter-planar distance d (Σ (dobs
dcal ¼ minimum), it is found that the above parameters of all the samples match well with those of rhombohedral crystal system. There is no indication of change in basic structure of BFO on the increasing value of x up to its 15%. A few un-indexed peaks observed in the XRD patterns (marked with n in Fig. 1) correspond to secondary or impurity phases [22]. The percentage of the content of the impurity phase was estimated using a standard empirical relation: Pi ¼ Pp (Ii/Ip þ Ii) where i and p represent the amount of impurity and perovskite phase P respectively [23], and was found to be 2–5%. The content of impurity phase decreases

Fig.1. (a) X-ray diffraction patterns of (Bi1
xCax)(Fe1
xMnx)O3 (x¼ 0.00, 0.05, 0.10, and 0.15) samples calcined at 700 1C; (b) enlarged version of the XRD patterns in the range of 2θ from 31.61 to 32.81.

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from on increasing substitution concentration. Thus, it is interesting to observe that on increasing content of x controls the formation of impurities phases of BFO. Two most intense diffraction peaks (104) and (110) are clearly separated in all the samples (shown in Fig. 1b), but the peak position shifts towards the higher angle. The small shift may be due to the slight difference in ionic radius of Ca þ 2 (1.00 Å) and Mn þ 4 (0.53 Å) as compared to that of Bi þ 3 (1.03 Å) and Fe þ 3 (0.65 Å) respectively. This difference in radii of substituted ions creates a small unit cell distortion in terms of variation in Bi–O and Fe–O–Fe bond distances as well as the creation of defects. Bulk BFO phase is normally formed at temperatures above 825 1C when synthesized through a conventional solidstate reaction process [24]. The present study shows the formation of the samples at 700 1C which is much lower than the studies reported earlier. Due to high energy ball milling prior to calcinations, defect density in the samples increases which is in turn responsible for enhancing the diffusibility of the ingredients for phase formation. So mechanical activation process decreases the synthesis or phase formation temperature of the prepared system and also overcomes the micro structural problems [25,26]. Fig. 2 shows the room temperature FTIR spectra of BCFMO samples recorded in the wave number range of 400– 2000 cm
1. The band positions of metal oxygen bonds (Fe– O/Mn–O) were observed in the wave number range 400– 700 cm
1. The strong absorption bands near 552 cm
1 and another near 442 cm
1 correspond to Fe–O stretching and bending vibration in FeO6 octahedral unit in the perovskite structure respectively. These band positions are found to be in agreement with the characteristic infrared absorption bands of BFO reported earlier [27]. The broad nature of these vibration bands near 552 cm
1 arises due to overlapping of absorption peaks of both Fe–O and Bi–O at nearly the same wave number. With the increasing value of x, the absorption peaks shift slightly towards the higher wave number side due to the difference in ionic radius of substituted ions. The formation of perovskite structure in the samples is confirmed by the presence of metal–oxygen band (Fe–O–Fe) [28] and the shifting of the absorption bands toward higher wave number side is supported by the shifting of diffraction peaks towards higher diffraction angle for modified samples. The microstructures of the BCFMO sintered pellets are illustrated in Fig. 3. During the sintering process, a liquid phase is occurred, which may be due to the Bi2O3 vaporization at high temperature [29,30]. This liquid phase results rapid grain growth in BFO and the grain growth process is more or less looks completed. The microstructure of the samples looks dense and almost void free with less number of scattered pores. From the figure, pure BFO consist of very large grains with an average grain size of over 3 mm. The co-substituted samples have reduced grain size with homogeneous microstructure compared to that of pure BFO. A highly dense microstructure was obtained, because the substitution prevents the grain growth of BFO. On increasing value of x, the grains had more angular shape rather than circular, which may be due to the x content leading to an increase in surface energy. Therefore, the

Fig. 2. FTIR spectra of (Bi1
xCax)(Fe1
xMnx)O3 (x ¼0.00, 0.05, 0.10, and 0.15) samples at room temperature.

grains were found to have angular shape rather than curved surfaces. This was because the angular surface had surface energy less than curved surfaces [31]. Fig. 4 shows the EDX images of BCFMO samples. From the images, it is confirmed that all samples contain elements of Bi, Fe, Ca, Mn and O and there is no trace of other elements presence. The atomic ratio of Bi:Ca: Fe:Mn is closely matched with the expected stoichiometry ratio of the solid solutions. The EDX patterns also conclude the completion of chemical reaction of the precursor to form the desired materials without any loss of ingredients during the sintering process. 3.2. Electrical characteristics There are a number of methods, which are employed to study the dipolar dynamics of the ceramics. Among all the techniques, the dielectric spectroscopy is a most useful tool for the investigation of dipolar relaxation processes over a wide range of frequency domains. Again, among all types of relaxation, the dielectric relaxation has been the most focused subject of interest due to its technological and industrial importance [32]. Generally, the dielectric response involves effects from ‘dipolar’ and ‘charge carrier’ behaviors. The former one is responsible for the restoration of the zero residual polarization after discharging while the later one is associated with partial recovery on discharge but always leaves a finite polarization in the system [33]. 3.2.1. Dielectric studies with frequency Fig. 5 (a and b) shows the frequency dependence of the dielectric constant (εr) and loss tangent loss (tan δ) at 100 1C for BCFMO samples. It is observed that for all samples, the value of εr decreases rapidly in the low-frequency region whereas decrease is quite slow in the high-frequency region

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Fig. 3. SEM micrographs of (Bi1
xCax)(Fe1
xMnx)O3 (x ¼0.00, 0.05, 0.10, and 0.15) samples.

(almost approaching to frequency independent response). In the high-frequency region, the higher value of εr can be described in terms of a rapid polarization processes with no contribution of ionic motion. At higher frequency, the ions oscillate without reaching to the sample–electrode interface. The contribution of interfaces or grain boundaries to apparent dielectric constant in the spinel ferrites can be explained by the Maxwell–Wagner theory [34]. According to this theory, it is assumed that a dielectric medium is made up of highly conducting grains and poorly conducting grain boundaries. The grain boundaries are more effective at lower frequencies while the grains are found to be more effective at higher frequencies. As the measurements of studied materials were performed with silver paint (as electrode), the low-frequency high dielectric constant implies a main non-intrinsic contribution which comes from the electrode and sample interface [35]. Additionally, hopping conductivity can give frequency dependent contribution to apparent dielectric constant at higher frequencies [34]. In ceramics, the space charges polarization normally occurs due to defects or oxygen vacancies (VO) created during high-temperature sintering of pellet samples [36]. The presence of VO will lead to a change in the valence state of iron (Fe þ 3-Fe þ 2 ions) due to electron exchange. Therefore, the decrease in εr with the increase in frequency is due to the energetic hopping of electrons along the direction of the applied field between Fe þ 3 and Fe þ 2 ions at the octahedral

sites [37]. The decrease in εr with rise in frequency is a general trend because of the difference in the contribution of different types of polarization at different frequencies. Summarizing, both the intrinsic and non-intrinsic (grain boundary/interface) contributions are considered to the apparent or dielectric constant (more explanations are included in the Section 3.2.2).The value of tan δ decreases as the frequency increases because the frequency of charge carriers cannot follow the frequency of the applied field after certain range of frequency. At low-frequency range the tan δ value increases with the increase in value of x, and is nearly same in the high frequency region. The dielectric loss arises mainly due to impurities and imperfections in the crystal lattice, which cause polarization to lag behind with the applied alternating field. The density of the materials also plays an important role in the variation of dielectric constant and dielectric loss, and thus porous materials have low dielectric constants and dielectric losses. 3.2.2. Complex impedance studies Impedance spectroscopy (IS) is the most reliable and important technique to study the electrical properties and process of the materials. The IS technique is based on analyzing the ac response of a system to a sinusoidal perturbation, and subsequent calculation of impedance and related parameters as a function of frequency of the perturbation at different temperatures. The complex impedance can be expressed as

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Fig. 4. EDX images of (Bi1
xCax)(Fe1
xMnx)O3 (x¼ 0.00, 0.05, 0.10, and 0.15) samples.

|Z|¼ [(Z0 )2 þ (Z″)2]1/2; where Z0 and Z″ are the real and imaginary components of the complex impedance respectively. This complex impedance of the ceramic can be demonstrated as the sum of the single RC circuit with parallel combination as per the below relation: Z Yðτ; TÞdðτÞ ð1Þ Z nðTÞ ¼ Z 0 ðTÞ 1þ jωτ This complex equation can be represented as real and imaginary part and can be expressed as Z Yðτ; TÞdðτÞ Z'ðω; TÞ ¼ Z 0 ðTÞ ð2Þ 1 þ ω2 τ 2 Z ðωτÞnYðτ; TÞdðτÞ Z''ðω; TÞ ¼ Z 0 ðTÞ ð3Þ 1þ ω2 τ2

Here, τ ¼ Rb Cb represents the relaxation time, T ¼ time period and Y (τ, T) ¼ distribution function of relaxation time. The variation of imaginary part of complex impedance Z''ðω; TÞ provides information about the distribution function Y (τ, T). Fig. 6 shows the complex impedance spectrum at few selected temperatures for BCFMO samples. The impedance property of ceramics is characterized by the formation of semicircular arcs whose pattern of evolution changes with change in temperature. The extent of intercept of semicircles on the real axis (Z0 -axis) and its number in the spectrum provide information on the kind of electrical processes occurring within the material. The semicircular arcs of the impedance pattern can mainly be attributed to a parallel combination of resistance and capacitance. As temperature

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Fig. 5. Frequency dependence of (a) dielectric constant (εr) and (b) loss tangent (tan δ) at 100 1C for (Bi1
xCax)(Fe1
xMnx)O3 (x ¼0.00, 0.05, 0.10, and 0.15) samples.

increases from room temperature, the arc progressively becomes semicircular with the shift of its center below the real axis. For the sample (x ¼ 0), the presence of a single semicircle up to 300 1C suggests that the electrical processes in the material arise due to the contribution from bulk material, and thus can be modeled as an equivalent electrical circuit (inset) comprising of a parallel combination of bulk resistance (Rb) and bulk capacitance (Cb) [38] only. For the modified BFO samples in the temperature range of 200–300 1C, two semicircular arcs are observed in the high- and low-frequency regions. This type of electrical process can be modeled in terms of equivalent electrical circuit according to the “Brick layer model” [39], comprising of series combination of two parallel R–C circuits of both grain and grain boundary effects. But for the sample (x ¼ 0.15), a third semicircular arc (smaller) was also observed above 275 1C. This may be attributed to the beginning of the polarization effects (polarization at the material–electrode interface) [40]. The electrode polarization can be neglected as it generally dominates at very low frequencies. The experimental impedance data were fitted using a software ZSIMP WIN version 2. For ideal Debyelike response, an equivalent circuit consists of parallel combination of CQR and CR, where Q is known as constant phase element (CPE). The admittance of Q is defined as Y (CPE) ¼ A0 (jω)n ¼ Aωn þ jBωn, with A ¼ A0Cos(nπ/2) and B¼ A0Sin(nπ/2), where A0 and n are frequency independent but temperature dependent parameters. The magnitude of the dispersion is determined by knowing the value of A0. The value of n is 0 r n r 1 where n ¼ 1 corresponds to the behavior of an ideal capacitor and n¼ 0 represents that of an ideal resistor [37]. The diameter of the arc decreases with increasing temperature for all the samples; indicating the decrease of resistance with increase in temperature; which in turn, suggests the NTCR behavior of the compound. For an ideal Debye-type relaxation, a perfect semi-circle with its center at real Z-axis should be observed. But in the prepared system, the centers of semicircles exist well below the real axis indicating the non-

Debye-type distribution of relaxation time. This depicts that there is a distribution of relaxation time instead of a single relaxation time in the material [41]. The intercept of each semicircle on real Z'-axis gives the value of bulk and grain boundary contribution in the resistance/impedance. The semicircles in the impedance spectrum have a characteristic peak occurring at a unique relaxation frequency usually referred as resonance frequency (fr) (ωr ¼ 2πfr). It can be expressed as ωr RC ¼ ωrτ ¼ 1 and thus fr ¼ 1/2πRC, where τ is the relaxation time. The relaxation time due to bulk effect (τb) has been calculated using the equation ωrτb ¼ 1 or, τb ¼ 1/2πfr. 3.2.3. Polarization studies As we know, ferroelectric properties of ceramics are influenced by their composition, microstructure, crystal phase, and lattice defects like oxygen vacancies [42]. Ferroelectricity persists in samples due to the long-range polar orders of dipoles and will be affected if there will be certain degree of disruption in the polar order [43]. Fig. 7 shows room temperature ferroelectric hysteresis loops of BCFMO samples. It is observed from the loops that value of maximum polarization as well as the area of the hysteresis loops strongly depends on the value of x in the modified samples. With the increasing value of x, the behavior of the ferroelectric loops changes noticeably; the value of polarization increases. For the samples with x¼ 0.00, 0.05, 0.10 and 0.15, remnant polarization (2Pr) obtained was found to be 0.56, 0.64, 0.77 and 0.95 μC/cm2 respectively. The contributions to the apparent ferroelectric loop possibly arise from the leakage contributions. Usually bulk BFO endures a high-leakage current at room temperature. It is due to the involvement of the too high concentrations of oxygen vacancies induced by the loss of bismuth and the mixed valences of Fe (Fe þ 2 and Fe þ 3) [44,45]. Oxygen vacancies are one of the major sources of movable charges in BFO and the mixed valence Fe ions cause increased conductance, possibly through the double exchange mechanism via Fe þ 2–O–Fe þ 3 chains [46]. It is believed that

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Fig. 6. Complex impedance spectrums of (Bi1
xCax)(Fe1
xMnx)O3 (x ¼0.00, 0.05, 0.10, and 0.15) samples at different temperature with electrical equivalent circuit (inset).

co-substitution significantly reduces the movable charges and the contribution to leakage current density is also minimized, as a result ferroelectric behaviors of the samples were enhanced. 3.3. Magnetic characteristics The magnetic hysteresis (M–H) loops of BCFMO samples were investigated at room temperature using a VSM with an applied magnetic field of
10 kOe r H r 10 kOe, as shown in Fig. 8. The magnetic properties of materials are usually characterized by a hysteresis loop, which determines the behavior of material when excited by an external magnetic field. According to Fig. 8, it is clear that the BFO exhibits a typical anti-ferromagnetic behavior (i.e., M depending linearly on H and no hysteresis of M on H), it is consistent with other reports [47,48]. Although the crystal structure of BFO makes the appearance of weak ferromagnetism arising from the canting of the anti-ferromagnetic sub-lattices, the spiral spin structure leads to a cancellation of the macroscopic magnetization [49]. All the samples exhibit weak ferromagnetic behavior in contrast to BFO, which can be attributed to the size effect [50] and this magnetic behavior of small sized grains of anti-

ferromagnetic materials as discussed in the proposed theory [51]. In all the samples, saturation is achieved at 10 kOe. Thus saturation magnetization (Ms) of the above solid solution increases with increase in the value of x. The saturation magnetization (Ms) are 0.145, 0.542 and 0.784 emu/g for x¼ 0.05, 0.10 and 0.15, respectively. The remnant magnetizations (Mr) of these samples are 90  10–4, 1015  10
4and 1379  10
4 emu/g, respectively. The large enhancement in the value of magnetization is attributed to either the destroying of the spiral spin-modulated incommensurate structure, or the increasing of the spin canting angle resulting in the net macroscopic magnetization [52,53]. So, co-substitution can only suppress but cannot destruct the spin cycloid, which is responsible for the smooth increase of the Ms and Mr values. 4. Conclusion The high energy ball milling method followed by thermal treatment has been used to prepare (Bi1
xCax)(Fe1
xMnx)O3 (x¼ 0.00, 0.05, 0.10, and 0.15) systems, and their structural, dielectric, ferroelectric and magnetic properties have been studied using standard experimental techniques. Studies of basic structural parameters using x-ray diffraction technique reveal that all samples

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Fig. 7. Room temperature P–E loops of (Bi1
xCax)(Fe1
xMnx)O3 (x ¼0.00, 0.05, 0.10, and 0.15) samples.

the presence of uncompensated spin and the collapse of the cycloidal spin structure. A high value of saturation and remnant magnetization is observed in the co-substituted samples, which may help in finding a new way for improving magnetization in BFO materials. This may, in turn, will lead to many technological applications. Acknowledgments The authors is thankful to (i) Mr. B.B. Palei of IMMT, Bhubaneswar for his kind help in x-ray diffraction analysis (XRD), (ii) Dr. M.N. Goswami, Midnapore College, Midnapore for his kind help in carrying out FTIR analysis and (iii) Prof. P.K. Mahapatra of this department for his fruitful discussions. Fig. 8. Magnetic hysteresis loops of (Bi1
xCax)(Fe1
xMnx)O3 (x ¼0.00, 0.05, 0.10, and 0.15) samples at room temperature.

crystallized in perovskite form, and were confirmed by FTIR analysis. A highly dense SEM microstructure recorded on the pellet samples has confirmed the formation of good quality of the samples. The value of dielectric constant (εr) increases on increasing the value of x and the enhancement in dielectric properties will make the materials useful for energy storage devices in microelectronics. At room temperature, the modified samples exhibit weak ferromagnetic behavior which has been assigned to

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