Synthesis and characterization of La0.75Ca0.15Sr0.05Ba0.05MnO3–Ni0.9Zn0.1Fe2O4 multiferroic composites

Journal of Magnetism and Magnetic Materials ∎ (∎∎∎∎) ∎∎∎–∎∎∎

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 Q1 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 66

Contents lists available at ScienceDirect

Journal of Magnetism and Magnetic Materials journal homepage: www.elsevier.com/locate/jmmm

Synthesis and characterization of La0.75Ca0.15Sr0.05Ba0.05MnO3–Ni0.9Zn0.1Fe2O4 multiferroic composites Md. D. Rahaman a,n, S.H. Setu a, S.K. Saha b, A.K.M. Akther Hossain b a b

Deparment of Physics, University of Dhaka, Dhaka 1000, Bangladesh Department of Physics, Bangladesh University of Engineering and Technology (BUET), Ramna, 1000 Dhaka, Bangladesh

art ic l e i nf o

a b s t r a c t

Article history: Received 21 December 2014 Received in revised form 1 March 2015 Accepted 6 March 2015

In the present work, we report on structural, dielectric, impedance spectroscopic studies and magnetoelectric properties of (1
x) La0.75Ca0.15Sr0.05Ba0.05MnO3 (LCSBMO) þ(x) Ni0.9Zn0.1Fe2O4 (NZFO) (x ¼0.0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0) composites. The composites were prepared by the solid state reaction route. The coexistence of a cubic spinel NZFO phase and a tetragonal LCSBMO phase in the composites is confirmed by the X-ray diffraction measurement. Scanning electron microscopy images reveal that NZFO particles were distributed non-uniformly with some porosity in the LCSBMO matrix. Frequency dependent dielectric constant shows usual dielectric dispersion behavior, which may be attributed to the Maxwell–Wagner type interfacial polarization. At higher frequencies ( Z105 Hz), due to electronic and ionic polarizations only, the dielectric constant is independent of frequency. Complex impedance shows semicircular arc due to the domination of grain boundary resistance and electric modulus confirms the presence of hopping conduction. The AC conductivity (σ AC ) obeys the power law and the linearity of log ω2 versus log σ AC plots indicates that the conduction mechanism is due to small polaron hopping. Low frequency dispersion in permeability is due to domain wall motion and the frequency stability of permeability indicates that the arrangement of the magnetic moment in the polarization process can keep up with the external field. The maximum magnetoelectric voltage coefficient of  40 mV Oe
1 cm
1 for x ¼ 0.8. & 2015 Published by Elsevier B.V.

Keywords: Manganite–ferrite composites Sintering X-ray diffraction Dielectric response Magnetic property Magnetoelectric voltage coefficient

1. Introduction Magnetoelectric (ME) effect which occurs due to interaction between electric and magnetic dipole in ferroelectric–ferromagnetic multiferroic materials [1] has brought about much interest, not only for technological applications in possible miniaturization and integration for electronic devices such as waveguides, switches, transducers, modulators, magnetic storage systems, read heads for hard disks, sensing devices, electrically controlled microwave phase shifters or ferromagnetic resonance devices, magnetically controlled electro-optic or piezoelectric devices, broadband magnetic field sensors, and magnetoelectric memory cells due to their capability of monitoring magnetization/electric polarization as a function of external electric/magnetic field, but also from the view point of fundamental scientific understanding [2– 7]. With the development of science and technology, rapid development of miniaturized digital circuits, and excellent highfrequency characteristics of the microwave devices, single-phase n

Corresponding author. Fax: þ 880 28613642. E-mail address: [email protected] (Md.D. Rahaman).

multiferroic materials like Cr2O3, BiFeO3, Pb(Fe0.5Nb0.5)O3, etc. [3,8,9] cannot satisfy the requirement of much technological applications because they normally have a feeble ME effect at room temperature as well as a Neel temperature far below room temperature [10]. Then there have been resurgence among the researchers for the dual phase composite materials based on the concept of “product property” [11] to mitigate the requirements of miniaturization and excellent high frequency performance of the electronic devices. Although reported values for the dual phase composites are 10–100 times higher than the ME coefficient for single phase multiferroics [12] but their practical applications are still at infancy. In the recent years, large numbers of manganite–ferrite (hard and soft) composites have been produced and investigated such as La0.67Sr0.33MnO3/(Zn0.6Fe0.4)[Ni0.4
xCuxFe1.6]O4 [13], La0.7Ca0.2Sr0.1 MnO3/CoFe2O4 [14], LCMO/NiFe2O4 [15–17], La2/3(Ca0.6Ba0.4)1/3 MnO3/NiFe2O4 [18], La0.7Sr0.3MnO3/NiFe2O4 [19], La0.7Sr0.3MnO3/ CuFe2O4 [20, 21], La0.7Sr0.3MnO3/ZnFe2O4 [22], La0.67Ca0.33MnO3/ CuFe2O4 [23], La0.7Sr0.3MnO3/CoFe2O4 [24, 25], La0.67Ca0.33MnO3/ Fe3O4 [26, 27], Cu0.7Cd0.3Fe2O4/La0.67Sr0.33Mn0.98Co0.02O3 [28], L0.67Sr0.33MnO3/BaFe11.3(ZnSn)0.7O19 [29], La2/3Sr1/3MnO3/BaFe12 O19 [30], La0.7Sr0.3MnO3/SrFe12O19 [31], and SrFe12O19/La1
xCax

http://dx.doi.org/10.1016/j.jmmm.2015.03.024 0304-8853/& 2015 Published by Elsevier B.V.

Please cite this article as: Md.D. Rahaman, et al., Journal of Magnetism and Magnetic Materials (2015), http://dx.doi.org/10.1016/j. jmmm.2015.03.024i

67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93

Md.D. Rahaman et al. / Journal of Magnetism and Magnetic Materials ∎ (∎∎∎∎) ∎∎∎–∎∎∎

2

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 66

MnO3 [32] with a view to enhance the magnetoresistance because the introduction of high resistivity second phase in the manganites has the barrier effect to the tunneling process and also causes the magnetic disorder near the grain boundary. However, to our knowledge, no reports are available in the literature on magnetoelectric (ME) properties of (1
x) [La0.75Ca0.15 Sr0.05Ba0.05MnO3] (LCSBMO) þ (x) [Ni0.9Zn0.1Fe2O4] (NZFO) composites by the solid state reaction method. As a magnetostrictive phase, we have selected NZFO due to its mixed spinel (AB2O4) structure. NZFO is a combination of NiFe2O4 (NFO) which has inverse spinel structure while ZnFe2O4 (ZFO) has normal spinel structure. The incorporation of Zn2 þ into the NFO spinel cell pushes the inverse structure to acquire mixed spinel [33,34]. The NZFO ferrite is highly magnetostrictive and resistive, since a Jahn– Teller ion such as Zn2 þ has a high coupling coefficient [35]. From an application point of view, NZFO is one of the most versatile, reasonable material for general use in both high- and low-frequency devices because of their high electric resistivity (exceeds 107 Ω m), high value of magnetic permeability, low dielectric loss, good chemical stability, low coercivities, high mechanical hardness, low porosity, reasonable cost and high Curie temperature, which further depends on the compositions and heat treatment of the samples [36–43]. On the other hand, La0.67Sr0.33MnO3 (LSMO) is known for its simple perovskite crystal structure and dopant-controlled Curie temperature accompanied by a metal-to-insulator transition (MIT) above room temperature [44]. The doping at A site (the divalent alkali-earth metal ions) of these manganites can both lead to the appearance of Mn4 þ ions which produces Mn3 þ –O–Mn4 þ double-exchange function [45] and change the average radius of A-site ions to adjust the bond length and bond angle of ABO3 oxygen octahedrons and to influence double-exchange function [46,47]. We have selected La0.75Ca0.15Sr0.05Ba0.05MnO3 (LCSBMO) as the matrix which is a mixture of La0.7Ca0.3MnO3, LSMO, and La0.7Ba0.3MnO3, respectively because light doping of Sr2 þ and Ba2 þ for Ca2 þ ions caused a shift of paramagnetic (PM) to ferromagnetic (FM) transition temperature (Tc) to a higher value close to room temperature as well as improves the electrical and magetotransport properties [48–50]. On the other hand, the simultaneous of Sr–Ba results in a Mn3 þ /Mn4 þ mixed valence state creating mobile charge carriers and canting of Mn spins [51]. In the present study, we have studied the structural, morphological, electrical, magnetic, and magnetoelectric properties of (1
x) LCSBMOþ(x) NZFO composites.

2. Experimental details 2.1. Materials and method The (1
x) La0.75Ca0.15Sr0.05Ba0.05MnO3 þ(x) Ni0.9Zn0.1Fe2O4 (x ¼0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0) composites were prepared using the following steps. First, a pure polycrystalline La0.75Ca0.15Sr0.05Ba0.05MnO3 (LCSBMO) and Ni0.9Zn0.1Fe2O4 (NZFO) were prepared using the conventional solid state reaction method. For the preparation of LCSBMO powders, high purity (4 99%) powders of La2O3, CaCO3, SrCO3, BaCO3, Mn2O3 were mixed in stiochiometric proportions and ground for 4 h in order to achieve homogeneity in the mixed powder. Then the mixed powders were calcined in air at 900 °C for 6 h with a heating rate of 10 °C/min and cooling rate of 5 °C/min in a furnace to establish the course of nucleation for the grain growth and to felicitate the decomposition of the substituent oxides/carbonates. The calcined powders were again ground to make fine powders. A similar procedure has been followed to prepare NZFO ferrite powder using NiO, ZnO, and Fe2O3 (with a purity higher than 99%) as starting materials. For

preparing the composites, we have mixed LCSBMO and NZFO powders together according to the stoichiometric ratios and ground for 2 h to allow good mechanical mixing with small amount of diluted polyvinyl alcohol (PVA) as binder. The mixed powder was palletized into small disks (1.3 cm diameter and  0.2 cm thickness) and toroids (outer diameter 1.3 cm, inner diameter 0.7 cm and thickness  0.2 cm) using a steel die and a hydraulic press with a uniaxial pressure of 5000 psi and sintered 1200 °C for 5 h in a closed alumina crucible in atmospheric air followed and cooled at room temperature by adjusting the cooling rate to yield the final product. The heating rate and cooling rate was set at 10 °C/min and 5 °C/min, which is an essential and important process to achieve the required oxygen content in the material. Composites were sintered with different temperatures since the sintering process is key issue to develop the most appropriate structure for the application, complete the inter-diffusion of the component metal ions into the desired crystal lattice and to establish the appropriate valencies for the multi-valent ions by proper oxygen control [52,53]. 2.2. Characterizations The phase formation of all the composites were characterized by a Philips PanAnalytic Xpert Pro X-ray diffractometer with a Cu anode (Cu-Kα radiation source with λ ¼ 1.541 Å) with the step size 0.02 operating at 40 kV and 30 mA at room temperature and 1 °C/ min scanning speed, by collecting the data in a 2θ range of 10–70°. The lattice parameters were calculated from XRD data. The microstructure of the composites was analyzed using a scanning electron microscope (SEM) (JEOL JSM-6460, Japan). For electrical measurement samples were painted conducting silver paste on either side to ensure good electrical contacts. The toroids were wrapped with 4-turns copper wire to measure magnetic permeability. The frequency dependence of dielectric constant and magnetic permeability were measured using a Wayne Kerr Impedance Analyzer (B 6500 series) in the frequency region 100 Hz to 120 MHz [54]. The complex electric modulus (M* (ω) = 1/ε* = 1/ε′ − jε″) has been determined using real and imaginary components of the dielectric permittivity. The ME effect was obtained by applying an AC magnetic field superimposed on a DC magnetic field on the sample, and then measuring the output signal with applied DC magnetic field. An electromagnet was used to provide a dc magnetic field of up to 0.86 T. A signal generator was used to drive the Helmholtz coil to generate an ac magnetic field of 0.0008 T. The output voltage generated from the composite was measured by a Keithley electrometer (Model 2000) as a function of dc magnetic field. ME voltage coefficient, α ME , was calculated using relation, α ME = V /H AC × d , where V is the ME voltage across the sample surface and HAC is the amplitude of the sinusoidal magnetic field and d is the thickness of the sample [54].

3. Results and discussions 3.1. X-ray diffraction analysis Fig. 1(a) and (b) shows the XRD of La0.75Ca0.15Sr0.05Ba0.05MnO3 (LCSBMO) manganite and Ni0.9Zn0.1Fe2O4 (NZFO) ferrite sintered at 1200 °C. From Fig. 1(a), it is observed that the diffraction pattern for LCSBMO shows all the allowed reflection lines corresponding to single-phase tetragonal structure. The lattice parameters for LCSBMO are a¼ b¼3.8625 Å and c¼ 12.4822 Å. On the other hand, from Fig. 1(b), it is observed that the XRD diffraction pattern for NZFO shows all the allowed reflection lines corresponding to single-phase cubic spinel structure. The lattice constant of NZFO was determined by the method described elsewhere [54,55]. The

Please cite this article as: Md.D. Rahaman, et al., Journal of Magnetism and Magnetic Materials (2015), http://dx.doi.org/10.1016/j. jmmm.2015.03.024i

67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95 96 97 98 99 100 101 102 103 104 105 106 107 108 109 110 111 112 113 114 115 116 117 118 119 120 121 122 123 124 125 126 127 128 129 130 131 132

Md.D. Rahaman et al. / Journal of Magnetism and Magnetic Materials ∎ (∎∎∎∎) ∎∎∎–∎∎∎ # x=0

Intensity (a. u.)

Ts = 1200 ° C

(220)

(211)

(200) (202)

(110)

20

30

40 50 2θ (Degree)

(311)

10

60

x = 0.2 x = 0.4

*

*

x = 0.8

*

*

x=1

*

*

x = 0.6

20

30

*

#

* *

#

#

##

#

# #

#

#

#

#

*

#

*

* *# #

*

*# 40

*

##

50

# #

#

*

#

*# * *# * * *

#

60

70

2θ (Degree)

lattice parameters with doping suggests that the Fe, Ni, and Zn ions from NZFO may occupy the Mn3 þ site in the LCSBMO. Therefore, the increase in the lattice volume may result in a difference in the ionic radii because the ionic radii of Fe3 þ , Ni2 þ and Zn2 þ are bigger than the Mn3 þ ions, so the value of the total lattice volume increases with an increase in x.

(440)

(511) (422)

(222)

30 40 50 2θ (Degree)

*

#

#

Fig. 2. X-ray diffraction pattern for (1
x) La0.75Ca0.15Sr0.05Ba0.05MnO3 þ (x) Ni0.9Zn0.1Fe2O4 (for x ¼0.0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0) composites sintered at 1200 °C, along with individual ferrite and manganite phases, where (n) denotes the ferrite phase and (#) denotes the manganite phase.

Ni0.9Zn0.1Fe2O4

(400)

(220) (111)

20

#

70

Ts = 1200 ° C

10

#

* NZFO # LCSBMO # #

##

#

x = 0.1

(114)

Intensity (a.u.)

(110)

La0.75Ca0.15Sr0.05Ba0.05MnO3

10

Intensity (a. u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 66

3

3.2. Surface morphology

60

70

Fig. 1. XRD patterns of (a) La0.75Ca0.15Sr0.05Ba0.05MnO3 and (b) Ni0.9Zn0.1Fe2O4 sintered at 1200 °C, respectively.

calculated lattice constant (a0) was found to be 8.394 Å for NZFO ferrite. The calculated lattice constants are in good agreement with those obtained previously [36–43]. The crystallite size was calculated from the full width half maxima (FWHM) for all peaks using Debye Scherrer's formula, D = (0.9λ /β cos θ) [56], where β is the FWHM of the XRD peak, λ is the X-ray wavelength and θ is the diffraction angle and found to be 46 nm for NZFO ferrite and 30 nm for LCSBMO manganite. Fig. 2 shows the X-ray diffraction (XRD) patterns of the (1
x) LCSBMO/(x) NZFO) composites (x ¼0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0) sintered at 1200 °C for 5 h. The peaks corresponding to LCSBMO and NZFO are marked by ‘#’ and ‘n’ sign, respectively in the diffraction pattern. It can be clearly seen that the ferromagnetic NZFO phase with a typical spinel structure and the LSCBMO manganite phase with a typical tetragonal perovskite structure coexist in the composites for all sintering temperature under study. The XRD peaks are identified with no detection of intermediate or interfacial phases, which indicates that the reactions between LCSBMO and NZFO are negligible. The results reflect the success in synthesizing composites with simultaneous manganite and ferromagnetic phases using solid state reaction method. The lattice parameters of the constituent phases are almost the same in all the composites. This indicates that the no noticeable transformation in the main structure that consists of LCSBMO and NZFO structures even if the composition of the composites is varied. Meanwhile, with increase in the NZFO content x; the peak intensity of the NZFO phase increases, which suggests that the NZFO and LCSBMO are coexistent in the composites of LCSBMO–NZFO. During the progress of the composite, the variant nature of the

The representative SEM micrographs of composites with x ¼0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0 sintered at 1200 °C are shown in Fig. 3 (a)–(g), respectively. Since composite materials are multiphase, it is desirable to know the individual phases as observed in the micrograph. It is observed that the microstructures consist of small, randomly oriented, non-uniform (as shape and size) grains and a certain amount of intergranular porosity of all the composites. The LCSBMO manganite phase appears dark region with bigger size of grains while the grains of ferromagnetic NZFO phase appear whitish with small grain size. The presence of porosity in the composites suggests that grain growth of the NZFO has been greatly inhibited when LCSBMO exists as impurity phase. It is known that the distribution of NZFO and LCSBMO phases in a composite plays a crucial role not only in the combined properties (magnetization, dielectric constant, resistivity, etc.) but also in the product properties like the ME output. The grain sizes of the samples were calculated by the linear intercept technique. It is seen from the SEM images that number of white grains goes on increasing with increase in NZFO content, which is as per the expectation. The grain sizes of the composites vary from 0.37 to 0.58 μm. The increase of average grain size of the composites with decrease in porosity decreases the grain boundary area and acts as an obstacle for domain wall motion [57]. The behavior of grain growth reflects the competition between the driving force for grain boundary movement and the retarding force exerted by pores. When the driving force of the grain boundary in each grain is homogeneous, the sintered body attains a uniform grain size distribution; in contrast, discontinuous grain growth occurs if this driving force is inhomogeneous. As a result of an increase in the grain size, the population of grain boundaries decreases, resulting in a decrease in the porosity of the final compound. 3.3. Dielectric properties Fig. 4(a) and (b) represents the variation of the real part (ε′) of the complex dielectric and dielectric loss (tan δ) with frequency. It is observed that the value of ε′ for x¼ 0.1 is higher as compared to

Please cite this article as: Md.D. Rahaman, et al., Journal of Magnetism and Magnetic Materials (2015), http://dx.doi.org/10.1016/j. jmmm.2015.03.024i

67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95 96 97 98 99 100 101 102 103 104 105 106 107 108 109 110 111 112 113 114 115 116 117 118 119 120 121 122 123 124 125 126 127 128 129 130 131 132

4

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 66

Md.D. Rahaman et al. / Journal of Magnetism and Magnetic Materials ∎ (∎∎∎∎) ∎∎∎–∎∎∎

Fig. 3. Typical SEM micrographs of the composites sintered at 1200 °C with (a) x¼ 0.0, (b) x ¼ 0.1, (c) x¼ 0.2, (d) x¼ 0.4, (e) x¼ 0.6, (f) x ¼0.8 and (g) x¼ 1.0.

the other composites due to higher LCSBMO concentration. But on the other hand, the value of ε′ for x ¼0.2, 0.4, 0.6, 0.8 and 1.0 is lower than that of pristine LCSBMO. It is observed that ε′ decreases with increasing the frequency for all composites showing dispersion in the lower frequency region from 100 Hz to 0.1 MHz. The low frequency dielectric dispersion is observed due to the presence Maxwell–Wagner type interfacial polarization theory [58,59]. In general, there is a characteristic relaxation time for the

charge transport and therefore the values of ε′ depend on the applied frequency. The higher values of ε′ observed at lower frequencies are due to the presence of the hetrogeneity in the composites. Hetrogeneities in the composites are the interfaces between the ferrite and manganite phases, which give rise to space charge polarization and contribute towards the high values of dielectric constant. On the application of electric field, space charges provided by NZFO phase accumulates at the interface of the two

Please cite this article as: Md.D. Rahaman, et al., Journal of Magnetism and Magnetic Materials (2015), http://dx.doi.org/10.1016/j. jmmm.2015.03.024i

67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95 96 97 98 99 100 101 102 103 104 105 106 107 108 109 110 111 112 113 114 115 116 117 118 119 120 121 122 123 124 125 126 127 128 129 130 131 132

Md.D. Rahaman et al. / Journal of Magnetism and Magnetic Materials ∎ (∎∎∎∎) ∎∎∎–∎∎∎

300 TS = 1200° C

250

x = 0.0 x = 0.1 x = 0.2 x = 0.4 x = 0.6 x = 0.8 x = 1.0

ε'

200 150 100 50 0 2 10

3

10

4

5

6

7

10 10 10 Frequency (Hz)

8

10

10

2.5 T S = 1200°C

2.0

x = 0.0 x = 0.1 x = 0.2 x = 0.4 x = 0.6 x = 0.8 x = 1.0

1.5

tanδ

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 66

1.0 0.5 0.0 2 10

10

3

4

5

10 10 10 Frequency (Hz)

6

10

7

10

8

Fig. 4. (a) Variation of dielectric constant (ε′) and (b) dielectric loss (tan δ) with frequency in the (1
x) LCSBMOþ (x) NZFO composites with x ¼0.0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0 sintered at 1200 °C.

phases due to different conductivities and permittivities of the constituent phases [60], and results in space charge polarization. As ferrite content increases, amount of space charge provided by NZFO also increases, consequently giving a higher value of dielectric constant. In general, the dielectric structure is supposed to be composed of two layers according to Maxwell–Wagner two layer model in agreement with Koop's phenomenological theory [59]. Accordingly, the dielectric structure of polycrystalline materials is assumed to be made up of two layers. First layer being a conducting layer consists of large grains and the other being grain boundaries are poor conductor [54]. The grain boundaries were found to be more effective at lower frequencies while the ferrite grains are more effective at higher frequencies [61]. At high frequency region (40.1 MHz), the values of ε′ becomes independent of frequency for all composites. In the low frequency region, different types of polarizations such as dipolar, electronic, ionic and interfacial polarization are the main sources of dielectric constant [62]. With the increase of frequency beyond a certain limit, dipoles are not able to align themselves with the applied electric field and contribution from some of these polarizations cease except electronic polarization. Only electronic polarization has a significant contribution to the dielectric constant at higher frequencies. In electronic polarization, the electric dipoles are unable to follow the first alteration of the applied alternating electric field beyond certain critical frequency. As a consequence, for all sintering temperatures, the values of ε′ become frequency independent in the high frequency region for the composites. In the higher frequency region, the dielectric constant increases with the increase of NZFO content in the composites, which can be explained by the fact that the dielectric constant of NZFO is much higher than that

5

of LCSBMO. 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. It is observed that the dielectric loss for all composites decrease with increasing frequency and remains constant value at higher frequency. The dielectric loss decreases with the increase of frequency which is attributed due to the fact that the hopping frequency of the electron exchange between Fe2 þ and Fe3 þ ions at adjacent octahedral sites cannot follow the changes of the externally AC applied electric field beyond a certain frequency limit [63,64]. In the frequency region (20 kHz to 0.1 MHz), composite (x¼0.1) show the dielectric loss peaks, according to Debye relaxation theory. The peak in the dielectric loss is attributed to the hopping frequency of the electron exchange between Fe2 þ and Fe3 þ ions at the tetrahedral and octahedral is nearly equal to the frequency of externally applied electric field [65] and condition ωτ = 1(ω = 2πf ) is satisfied. There might be a another reason for the maximum dielectric loss which may be attributed due to the fact that period of relaxation process is the same as the period of applied field. In other words, when the relaxation time is large as compared to the period of the applied field, losses are small. Similarly, when relaxation process is rapid as compared to the frequency of the applied field, losses are small. At higher frequencies, the dielectric losses are found to be low since domain wall motion is inhibited and magnetization is forced to change by rotation. The dielectric loss increased as the NZFO content increased in LCSBMO, but decreased with the increase in frequency. All the composites present the values of tan δ below 1 at 10 kHz which is remarkably low. At higher frequency, the loss is small, which means present ceramics could possibly be used in practical applications such as magnetically tunable filters and oscillators. 3.4. Impedance spectra analysis Fig. 5(a) shows the variations in real part (Z′) of impedance with frequency. It is observed that the magnitude of Z′ gradually decreases with increasing frequency up to a certain limiting frequency (  0.1 MHz) except for x¼ 0.1. The decrease in Z′ indicates that the conduction is increasing with frequency and at a frequency (Z0.1 MHz) it becomes almost frequency independent. The higher values of Z′ at lower frequencies means the polarization in the composites is larger. The merger of Z′ at higher frequencies indicates possible release of space charge polarization/ accumulation at the boundaries of homogeneous phases in the composites under the applied external field [66,67]. Our results are in well agreement with the reported results in BaTiO3/La0.7Ca0.3MnO3 and (Bi0.5Na0.5)0.94Ba0.06TiO3–PVDF composites [68,69]. On the other hand with the increase of NZFO concentration, the values of Z′ decrease in the low frequency ranges (up to a certain frequency), and thereafter appears to merge in the frequency region (Z 0.1 MHz). The frequency at which the release of space charge occurs depends upon the NZFO concentration. The decreasing value of Z′ indicates the increasing loss in resistive property of the composites. Such behavior is usual owing to the attendance of space charge polarization in the material [70,71]. In order to distinguish between the grain and grain boundary contribution the Nyquist plot of composites is shown in Fig. 5(b). In general, the plot would be composed of three semicircles, depending upon the electrical properties of the material. The semicircle at lower frequency represents the sum of resistance of grains and grain boundaries, while the semicircle at higher frequency corresponds to the resistance of grains only [72]. The third semicircle is also observed in some materials which could be due to the electrode effect [73,74]. In the present investigation, it has been observed that the impedance spectra of the composites sintered at 1200 °C do not take the shape of the semicircle but

Please cite this article as: Md.D. Rahaman, et al., Journal of Magnetism and Magnetic Materials (2015), http://dx.doi.org/10.1016/j. jmmm.2015.03.024i

67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95 96 97 98 99 100 101 102 103 104 105 106 107 108 109 110 111 112 113 114 115 116 117 118 119 120 121 122 123 124 125 126 127 128 129 130 131 132

Md.D. Rahaman et al. / Journal of Magnetism and Magnetic Materials ∎ (∎∎∎∎) ∎∎∎–∎∎∎

6

0.75 T S = 1200°C

0.60

x = 0.0 x = 0.1 x = 0.2 x = 0.4 x = 0.6 x = 0.8 x = 1.0

M'

0.45 0.30 0.15 0.00 2 10

0.12 0.10 0.08

M"

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 66

0.06

10

3

4

5

6

10

6

10

10 10 10 Frequency (Hz)

x = 0.0 x = 0.1 x = 0.2 x = 0.4 x = 0.6 x = 0.8 x = 1.0

7

10

7

10

8

T S = 1200°C

0.04 0.02 0.00 2 10

Fig. 5. Impedance spectra (1
x) LCSBMOþ (x) NZFO composites with x ¼0.0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0: (a) real part (Z‵) and (b) Cole–Cole plot of complex impedance sintered at 1200 °C.

10

3

4

5

10 10 10 Frequency (Hz)

8

Fig. 6. Electric modulus spectra of (1
x) LCSBMO þ (x) NZFO composites with x¼ 0.0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0: (a) real part (M′) and (b) imaginary part (M ‵‵) sintered at 1200 °C.

Simplifying (1) and substituting ε″ by ε′ tan δ , we get

M * (ω) = M′ (ω) + jM ′′ (ω) = rather represent straight line with large slope. This shows that grain boundary impedance is out of the measurement scale, suggesting the insulating behavior of the composites. Furthermore, single-component of complex impedance plots, irrespective of NZFO concentration, indicates the good homogeneity of the dielectric and conductive properties. Since the complex impedance spectra exhibit the single semicircular arcs, it can be said that the electrodes and contacts do not have significant impedance because electrode-contact would result in an additional semicircle. A single semicircle indicates that only one primary mechanism is responsible for the electrical conduction within the sample. In other words, the absence of other semicircles in the complex impedance plots suggests the dominance of grain boundary effect in the composites. So the electrical properties of the composites take place predominantly through the grain boundary effect only. Our results are in good agreement with the earlier reports [75]. 3.5. Electric modulus spectra analysis We have studied the complex electric modulus because it is possible to separate the electrode polarization effect the grain boundary conduction process through complex electric modulus study. The analysis of electrical relaxation in this system is carried out using the dielectric modulus Mn as formulated by Macedo et al. [76,54] M * (ω) =

1 ε*

=

1 ε′ ε′′ = +j = M′ (ω) + jM′′ (ω) ε′ − jε′′ ε′2 + ε′′2 ε′2 + ε′′2

(1)

1 tan δ +j ε′ (1 + tan δ) ε′ (1 + tan2δ)

(2)

Fig. 6(a) shows the variation of real M′ (ω) part of the electric modulus as a function of frequency. M′ (ω) increases from the low frequency of zero towards a high frequency limit and the dispersion shift to the high frequency region as increasing NZFO concentration for all composites. The zero values of M′ (ω) in the low frequency region confirm the presence of an appreciable electrode and/or ionic polarization in the composites under the studied frequency ranges and a continuous dispersion on increasing the frequency may be contributed to the conduction phenomena due to short range mobility of carriers. This implies the lack of restoring force for flow of charge under the influence of a steady electric field [76]. The increase of NZFO in LCSBMO results in lower value of M′ (ω), implying that the real part of the permittivity increases with ferrite filler and negligible contributions arising from the electrode polarization. For all the samples, the reduction in M′ (ω) indicates an increase in charge carrier density due to NZFO as well as LCSBMO manganite segmental motion. Fig. 6(b) shows the variation imaginary M′′ (ω) part of dielectric modulus with frequency. The modulus curves indicate not only the considerable shift in the modulus peak towards higher frequency side but also broadening of peaks with change in the concentration of NZFO. It is observed that M ′′ (ω) increases in the lower frequency region and exhibits a single relaxation peak centered at the dispersion region of M′ (ω). In the higher frequency region M′′ (ω) decreases and becomes constant which may be attributed due to limited carriers in potential wells. The frequency region (low frequency region) below the peak in M′′ (ω) spectra determines the range in

Please cite this article as: Md.D. Rahaman, et al., Journal of Magnetism and Magnetic Materials (2015), http://dx.doi.org/10.1016/j. jmmm.2015.03.024i

67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95 96 97 98 99 100 101 102 103 104 105 106 107 108 109 110 111 112 113 114 115 116 117 118 119 120 121 122 123 124 125 126 127 128 129 130 131 132

Md.D. Rahaman et al. / Journal of Magnetism and Magnetic Materials ∎ (∎∎∎∎) ∎∎∎–∎∎∎

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 66

which charge carriers are mobile on long distances, i.e., in between grains and is associated with hopping conduction process and at frequency above the peak (high frequency region), the carriers are spatially confined to their potential wells, being mobile on short distances, i.e., inside the grains of the composites and associated with relaxation polarization process, i.e., the carriers can execute only localized motion. The asymmetric modulus peaks shift toward higher frequencies, indicating correlation between motions of mobile ion charges [77], while the asymmetric in peak broadening shows the spread of relaxation times with different time constant. Thus, the peak frequency, fm is an indicative of transition from long range to short range mobility. Therefore, the nature of modulus spectrum suggests the existence of hopping mechanism of electrical conduction in the composites. 3.6. AC conductivity Fig. 7(a) shows variation of AC conductivity (sAC) with frequency for all composites. It is observed that the increase of sAC is less at low frequency (r1MHz), but at a frequency ( Z1 MHz), it increases with the frequency for all compositions. The increase of σ AC with NZFO concentration is due to the generation of the charge carrier concentration [78]. Furthermore, the increase of σ AC with frequency is generally attributed to the dipole polarization, i.e., the rotation of diploes between two equilibrium positions is involved. It is the spontaneous alignment of dipoles in one of the equilibrium positions that give rise to the nonlinear polarization behavior of these compositions. In low frequency region, the conductivity is almost independent of frequency which corresponds to DC conductivity (σDC ) because the resistive grain boundaries are more active at low frequencies according to the Maxwell–Wagner

7

double layer model for dielectrics. On the other hand, in the high frequency region which is known as hopping region, AC conductivity increases [79] because at higher frequencies the conductive grains become more active thereby increasing hopping of charge carriers [80] and obeys the following Joncher's law: σ AC (ω) = σ0 + Aωs , where σ AC (ω) is the total electrical conductivity, σ0 is the frequency-independent dc conductivity, where A is a preexponential factor, which is temperature-dependent, known as the Universal Dynamic Response (UDR) [81] and s is the power law exponent which generally varies between 0 and 1 depending on the temperature. The exponent s represents the degree of interaction between mobile ions with the lattice around them. The prefactor A determines the strength of polarizability. The frequency at which the conductivity deviates from the plateau region is called hopping frequency. In the higher frequency region, σ AC increases due to the hopping of charge carrier in finite cluster. In the low frequency region, where the conductivity is almost constant which corresponds to long-range ionic displacements (in between grains). While, in the high frequency region, where the conductivity increases strongly with frequency, the transport takes place in short distance (inside the grains of the composites) through the hopping process [82,33]. Thus, in the low frequency region, where the conductivity is almost constant, the transport takes place on percolated paths. While, in the high frequency region, where the conductivity increases strongly with frequency, the transport phenomena are dominated by contribution from hopping clusters. Variation of log σ AC as a function of log ω2 is depicted in Fig. 7(b). In small polaron hopping, ac conductivity increases with frequency whereas in large polaron hopping ac conductivity decreases with frequency [56]. The electrical conduction mechanism in terms of the electron and polaron hopping model has been discussed by Austin and Mott [83]. From Fig. 7(b), it is evident that the conductivity increases with the increase of frequency for all composites and curves are linear which indicates the conduction in these composites is due to small polarons associated with lattice strain accompanied by free charges. The grain boundaries are more active and the role of polaron hopping process is minor when the frequency of the applied field is low. On the contrary, when the frequency of the applied field increases, the conductive phase becomes more active and thereby increasing the importance of small polaron hopping conduction. Therefore a gradual increase in conductivity was observed with frequency. As reported by Alder and Fienleib [84], the frequency-dependent conduction is mainly attributed to a small polaron-type hopping mechanism. It is also observed that the conductivity increases with the increase of NZFO concentration which is attributed to the fact that the NZFO particles make chains; the electrical resistivity of the composites is reduced significantly due to the low resistivity of the NZFO phase as well as the parallel connectivity between NZFO and LCSBMO grains in all composites for all sintering temperatures [75]. 3.7. Complex permeability spectra

Fig. 7. (a) Variation of AC conductivity (σ AC ) with frequency and (b) plots of log σ AC versus log ω2 for (1
x) LCSBMO þ (x) NZFO composites with x ¼0.0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0 sintered at 1200 °C.

Fig. 8(a) shows the variation of real part ( μ′) of initial permeability with frequency sintered at 1200 °C. It is clearly observed that the initial permeability of the composition x ¼0.0 is lower that the composition x¼ 1.0. This may be due to lower grain size of LCSBMO and the absence of NZFO in the composites. The low frequency dispersion is attributed to the domain wall motion. μ′ demonstrates excellent frequency stability throughout the whole measured frequency range which shows the compositional stability and quality of the composites, except for the variations associated with the parametric resonance of the measurement circuit that appeared in the x ¼0.4, 0.6 and 0.8 composites at frequencies above 100 MHz. The dispersion of the permeability for x¼ 0.4,

Please cite this article as: Md.D. Rahaman, et al., Journal of Magnetism and Magnetic Materials (2015), http://dx.doi.org/10.1016/j. jmmm.2015.03.024i

67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95 96 97 98 99 100 101 102 103 104 105 106 107 108 109 110 111 112 113 114 115 116 117 118 119 120 121 122 123 124 125 126 127 128 129 130 131 132

Md.D. Rahaman et al. / Journal of Magnetism and Magnetic Materials ∎ (∎∎∎∎) ∎∎∎–∎∎∎

8

40 35

x = 0.0, x = 0.4,

x = 0.1, x = 0.6,

μ'

x = 0.2, x = 0.8,

x = 1.0

Ts = 1200°C

30 25 20 15 10 3 10

10

4

5

2500 2000 1500 1000

6

10 10 Frequency (Hz)

10

7

10

8

T s = 1200°C x = 0.0 x = 0.1 x = 0.2 x = 0.4 x = 0.6 x = 0.8 x = 1.0

500 0 3 10

structure, and phase distribution [92]. Fig. 8(b) shows the frequency dependence of relative quality factor (RQF). RQF increases with increase in frequency showing a peak, and then decreases with further increase in frequency for all composites. It is observed that RQF deteriorates at the frequency (Z10 MHz), this means that the magnetic loss is minimum at the frequency (r10 MHz) and then it rises rapidly. The loss is due to lag of domain-wall motion with respect to the applied alternating magnetic field and is attributed to various domain defects [93], which include nonuniform and non-repetitive domain-wall motion, domain-wall bowing, localized variation of flux density and nucleation and annihilation of domain walls. The peak corresponding to maxima in RQF shifts to higher frequency range as NZFO content increases. The NZFO possesses the maximum value of RQF. 3.8. Magnetoelectric properties

3000

RFQ

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 66

10

4

5

6

10 10 Frequency (Hz)

10

7

10

8

Fig. 8. (a) Variation of real part (μ‵) of permeability and (b) relative quality factor (RQF) with frequency for (1
x) LCSBMO þ (x) NZFO composites with x ¼0.0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0 sintered at 1200 °C.

0.6 and 0.8 at frequencies above 100 MHz may be attributed due to the resonance type domain wall motion [85]. The frequency stability may be attributed due to the arrangement of the magnetic moment in the polarization process can keep up with the external field [86]. It is believed that the reversible rotation of the magnetization vector is another factor making the permeability constant in the high frequency [87]. It also observed that the real part of initial permeability increases with the increase of NZFO concentration because the existence of strong magnetic phase (NZFO) strengthened the continuity and interaction of magnetic grains, which leads to an increase of initial permeability [88,89]. For instance, the sample with x¼ 0 has an initial permeability of about 14, whereas the composite with x¼ 0.8 has an initial permeability of about 27 and remains almost constant over the measured frequency range. This means that the composites possess an advantage of much wider working frequency range. In other words, magnetic permeability can be actually controlled by the introduction of the NZFO phase into the composite system. The natural resonance frequency peak, which originates from the effective anisotropy field, produces magnetic losses and gives the frequency limit for the application of electronic devices [90]. According to Snoek's law [91], the product of initial permeability and resonance frequency is a constant for ferromagnetic material, i.e., (μi − 1) fr = γ /2πMs where μi is the initial permeability, fr is the natural resonance frequency, Ms is the saturation magnetization, and γ is the gyromagnetic ratio. When the concentration of NZFO is increased, the magnetization of the composites becomes stronger. As a consequence, the cut-off frequency (i.e., the frequency where value reaches the half of the initial value) is decreased which is consistent with our results. In general, the initial permeability of composites is influenced by many factors, such as nonmagnetic phase content, microstructure, magnetic domain

Fig. 9 shows the variation of magnetoelectric (ME) voltage coefficient (α ME ) with applied DC magnetic field (HDC) for all composites. α ME remains constant throughout the range of the magnetic field studied except for x¼ 0.8. The constant ME output indicates that the magnetostriction reaches a saturation state during the magnetic poling and produces a constant electric field in the manganite phase. Beyond a certain applied magnetic field, the magnetostriction gets saturated and it produces constant electric field in the manganite phase. Thus, constant value of the α ME may be attributed to the saturated domain wall movement [94]. Similar results have been obtained for the composites [95, 54]. The value of α ME for x ¼0.1 is smaller than that of x ¼0.0, 0.2, 0.4, 0.6, and 1, respectively. On the other hand, α ME has the lower value for x¼ 0.2, 0.4, 0.6 and 1 than x ¼0.0. The decrease may be attributed to the fact that large grains can be polydomain and small ones cannot, hence, less effective in inducing a high piezomagnetic and piezoelectric coefficient. Another reason for decrease of ME voltage coefficient is the increase of porosity because the increase of pores disturbs the local connectivity, giving rise to local demagnetization and depolarization due to which the magnitude of the ME coefficient decreases [96]. The decrease of ME output with the ferrite content is attributed to the lower resistivity of ferrite phase which provides a leakage path for the charges developed in the manganite phase during poling. For x ¼0.8, α ME increases linearly with the applied magnetic field which is attributed to the fact that the magnetostriction reaches its saturation value at the time of magnetic poling and produces a constant electric field in the manganite phase. The highest value of α ME of 40 mV cm
1 Oe
1 has been obtained for x¼0.8 sintered at 1200 °C which is attributed to the uniform distribution and proper

Fig. 9. Variation of ME voltage coefficient (α ME ) as a function of static magnetic field for (1
x) LCSBMO þ(x) NZFO composites with x ¼0.0, 0.1, 0.2, 0.4, 0.6, 0.8 and 1.0 sintered at 1200 °C.

Please cite this article as: Md.D. Rahaman, et al., Journal of Magnetism and Magnetic Materials (2015), http://dx.doi.org/10.1016/j. jmmm.2015.03.024i

67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95 96 97 98 99 100 101 102 103 104 105 106 107 108 109 110 111 112 113 114 115 116 117 118 119 120 121 122 123 124 125 126 127 128 129 130 131 132

Md.D. Rahaman et al. / Journal of Magnetism and Magnetic Materials ∎ (∎∎∎∎) ∎∎∎–∎∎∎

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 66

9

Table 1 Comparative study of magnetoelectric voltage of particulate composites. ME Composite

Preparation route

ME output (mV cm
1 Oe
1)

Reference

NiFe2O4–Pb(Zr0.47Ti0.53)O3 Ni0.8Zn0.2Fe2O4–Pb0.93La0.07(Zr0.60Ti0.40)O3 Ni0.5Zn0.5Fe2O4–PbZr0.8Ti0.2O3 Ni0.9Zn0.1Fe2O4–PbZr0.52Ti0.48O3 0.1Ni0.8Zn0.2Fe2O4–0.9Pb0.955Sm0.03Zr0.65Ti0.35O3 NiFe2O4–Pb(Mg1/3Nb2/3)0.67Ti0.33O3 NiFe2O4–PZT Ni0.8Zn0.2Fe2O4/Sr0.5Ba0.5Nb2O6 Ni0.8Zn0.2Fe2O4– PbZr0.52Ti0.48O3 La0.75Ca0.15Sr0.05Ba0.05MnO3 þ Ni0.9Zn0.1Fe2O4

Sol–gel Solid state Double sintering solid state Double sintering ceramic method Solid state Solid state Spark plasma sintering High energy ball-milling Tape casting Solid state

0.5 7.53 0.84 0.748 22.5 10.43 25 26.6 40 40

[98] [99] [100] [101] [102] [103] [104] [105] [106] Present study

grain sizes (small grain size) of the two phases as well as the stiffness of the two phases which are comparable in magnitude and the elastic interaction between the two phases is strongest near x ¼0.8. Furthermore, it is well known that α ME depends on the mechanical coupling, resistivity, mole fraction of the constituent phases as well as the magnetostriction of the ferrite phase. Zinc doped nickel ferrite (NZFO) high magnetostriction coefficient because a Jahn–Teller ion such as Zn2 þ has a high coupling coefficient and its contribution to magnetostriction is high. It has been reported that the composites prepared with a lower content of the ferrite phase results in the reduction of piezoelectricity or magnetostriction respectively, leading to a decrease in the static ME voltage coefficient as predicted theoretically [97]. Therefore, the high value of α ME for x ¼0.8 is in agreement with theoretical prediction. As the NZFO concentration increases the conductivity of the composites increases, giving rise to depolarization phenomenon which causes to reduce the voltage in the process of polarization, thus reduces the residual polarization. Therefore, low value of ME coefficient for high NZFO concentration as compared to ferrite/ferroelectric composites may be due to this reason. The comparative study on various ME composites are listed in Table 1, which is comparable to the values of reported bulk composites, such as ferrite/ferroelectric composites.

4. Conclusion In summary, multiferroic magnetoelectric (ME) composites consist of LCSBMO as a manganite phase and NZFO as a ferrite phase were successfully synthesized using the conventional solid state reaction. X-ray diffraction patterns confirm the formation of cubic spinel phase for the ferrite phase and tetragonal perovskite phase for the manganite phase. The frequency dependent dielectric properties show the usual dielectric dispersion behavior for all the composites which can be explained on the basis of Maxwell–Wagner interfacial polarization effect. At higher frequencies, the dielectric constant remains independent of frequency because electric dipoles with large relaxation time cease to respond. Complex impedance shows semicircular arc due to the domination of grain boundary resistance and electric modulus confirms the presence of hopping conduction. The ac electrical conductivity increases with increasing frequency which is generally attributed to the inhomogeneities in the composites, induced probably by the presence of space charge. Thus, in the low frequency region, where the conductivity is almost constant, the transport takes place on percolated paths. While, in the high frequency region, where the conductivity increases strongly with frequency, the transport phenomena are dominated by contribution from hopping clusters. The linearity of log σ AC versus log ω2 plots indicates that the conduction mechanism is due to small polaron hopping associated with lattice strain accompanied by

free charges. Moreover stable frequency dependence of the permeability in a large frequency range which may be attributed due to the arrangement of the magnetic moment in the polarization process can keep up with the external field. The increment of initial permeability with the NZFO concentration is due to the existence of strong magnetic phase (NZFO) strengthened the continuity and interaction of magnetic grains, which leads to an increase of initial permeability. The applied magnetic field independent ME voltage coefficient of the composites is due to the saturated domain wall motion. The composite of x¼0.8 exhibits a maximum ME voltage coefficient of  40 mV Oe
1 cm
1 which is attributed to the uniform distribution and proper grain sizes (small grain size) of the two phases as well as the stiffness of the two phases which are comparable in magnitude and the elastic interaction between the two phases. Therefore, the preliminary investigations on AC electrical properties (impedance, conductivity, and dielectric), magnetic and ME voltage coefficient indicate that studied composites may open the possibility to explore new family of composites and also possess a useful figure of merit for technological application in devices.

Acknowledgements The authors are thankful to the Experimental Solid State Physics Laboratory of BUET for allowing us to carry out this research. The authors would also like to thank to the authorities of the Center for Advanced Research in Sciences (CARS), University of Dhaka and Atomic Energy Center, Dhaka for allowing us to use SEM and XRD facility, respectively.

Reference [1] R. Rani, J.K. Juneja, S. Singh, K.K. Raina, C. Prakash, Adv. Mat. Lett. 5 (2014) 229. [2] M. Fiebig, J. Phys, D: Appl. Phys. 38 (2005) R123. [3] W. Eerenstein, N.D. Mathur, J.F. Scott, Nature 442 (2006) 759. [4] M. Atif, M. Nadeem, R. Grossinger, R.S. Turtelli, J. Alloy. Compd. 509 (2011) 5720. [5] H. Zhang, X. Chen, T. Wang, F. Wang, W. Shi, J. Alloy. Compd. 500 (2010) 46. [6] Y. Bai, F. Xu, L. Qiao, J. Zhau, L. Li, J. Magn. Magn. Mater. 321 (2009) 148. [7] V.E. Wood, A.E. Austin, in: A.J. Freeman, H. Schmid (Eds.), Magnetoelectric Interaction Phenomena in Crystals, Gordon and Breach, 1975. [8] N.A. Hill, J. Phys. Chem. B 104 (2000) 6694. [9] P. Guzdek, M. Sikora, L. Gora, Cz Kapusta, J. Eur. Ceram. Soc. 32 (2012) 2007. [10] J. Ryu, A.V. Carzo, K.J. Uchino, H.R. Kim, Jpn. J. Appl. Phys. 40 (2001) 4948. [11] J.V. Suchetelene, Philips Res. Rep. 27 (1972) 28. [12] Ce-Wen Nan, M.I. Bichurin, S. Dong, D. Viehland, G. Srinivasan, J. Appl. Phys. 103 (2008) 031101. [13] D. Jinliang, L. Difei, W. Bolong, W. Xiaolong, Electron. Comp. Mater. 29 (2010) 8. [14] C.S. Xiong, F.F. Wei, Y.H. Xiong, L.J. Li, Z.M. Ren, X.C. Bao, Y. Zeng, Y.B. Pi, Y. P. Zhou, X. Wu, C.F. Zheng, J. Alloy. Compd. 474 (2009) 316. [15] B.B. Nayak, A. Mondal, S. Vitta, D. Bahadur, IEEE Trans. Magn. 47 (2011) 2728. [16] B.B. Nayak, S. Vitta, D. Bahadur, Mater. Sci. Eng. B 139 (2007) 171.

Please cite this article as: Md.D. Rahaman, et al., Journal of Magnetism and Magnetic Materials (2015), http://dx.doi.org/10.1016/j. jmmm.2015.03.024i

67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95 96 97 98 99 100 101 102 103 104 105 106 107 108 109 110 111 112 113 114 115 116 117 118 119 120 121 122 Q3123 124 125 126 127 128 129 130 131 132

10

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 66

Md.D. Rahaman et al. / Journal of Magnetism and Magnetic Materials ∎ (∎∎∎∎) ∎∎∎–∎∎∎

[17] G.W. Kim, S. Kumar, Y.J. Seo, K. Wang, C.G. Lee, B.H. Koo, Metals Mater. Int. 16 (2010) 807. [18] X. Yan, C. Shuiyuan, Y. Qingying, H. Zhigao, Z. Tanfeng, C. Qin, J. Fuqing Branch Fujial Norm. Univ., 2009-S1. [19] W.J. Feng, D. Geng, J.J. Ceng, J.Z. Sheng, S. Zhang, H. Yang, Integr. Ferroelectr.: Int. J. 129 (2011) 215. [20] Y.J. Seo, G.W. Kim, C.H. Sung, C.G. Lee, B.H. Koo, J. Phys.: Conf. Ser. 266 (2011) 012106. [21] Y.J. Kim, S. Kumar, Y.J. Seo, J. Chang, C.G. Lee, B.H. Koo, Surf. Rev. Lett. 17 (2010) 33. [22] Y.J. Seo, G.W. Kim, C.H. Sung, C.G. Lee, B.H. Koo, Korean J. Mater. Res. 20 (2010) 137. [23] Z.M. Tian, S.L. Yuan, Y.Q. Wang, L. Liu, S.Y. Yin, P. Li, K.L. Liu, J.H. He, J.Q. Li, Mater. Sci. Eng. B 150 (2008) 50. [24] C.H. Yan, Z.G. Xu, T. Zhu, Z.M. Wang, F.X. Cheng, Y.H. Huang, C.S. Liao, J. Appl. Phys. 87 (2000) 5588. [25] H. Yunhui, Y. Chunhua, W. Zheming, X. Zhigang, Z. Tao, L. Chunsheng, G. Song, X. Guangxian, Sci. China Ser. B 43 (2000) 561. [26] Z.Q. Xue, D.M. Xing, W.R. Hui, F.X. Nan, W. Yang, X.Z. Cai, Chin. Phys. Lett. 23 (2006) 2194. [27] Z.C. Xia, S.L. Yuan, G.H. Zhang, L.J. Zhang, J. Tang, W. Feng, J. Liu, G. Peng, L. Liu, Z.Y. Li, Q.H. Zheng, L. Cheng, C.Q. Tang, S. Liu, C.S. Xiong, J. Phys. D: Appl. Phys. 36 (2003) 217. [28] M.A. Ahmed, S.T. Bishay, S.M. Salem-Gaballah, J. Magn. Magn. Mater. 334 (2013) 96. [29] Q. Huang, J. Li, X.J. Huang, C.K. Ong, X.S. Gao, J. Appl. Phys. 90 (2001) 2924. [30] Z. Zhou, X. Wu, J. Du, S. Xu, X. Bai, G. Luo, F. Jiang, Acta Metall. Sin. (Engl. Lett.) 22 (2009) 51. [31] Z. Zi, Y. Fu, Q. Liu, J. Dai, Y. Sun, J. Magn. Magn. Mater. 324 (2012) 1117. [32] D.T.M. Hue, P. Lampen, T.V. Manh, V.D. Viet, H.D. Chinh, H. Srikanth, M. H. Phan, J. Appl. Phys. 114 (2013) 123901. [33] A. Sakanas, R. Grigalaitis, J. Banys, L. Mitoseriu, V. Buscaglia, P. Nanni, J. Alloy. Compd. 602 (2014) 241. [34] N. Adhlakha, K.L. Yadav, Smart Mater. Struct. 21 (2012) 115021. [35] R.S. Devan, C.M. Kanamadi, S.A. Lokare, B.K. Chougule, Smart Mater. Struct. 15 (2006) 1877. [36] A. Muhammad, M. Asghari, Mater. Sci. Eng. B 139 (2007) 164. [37] Q. Liu, L. Lv, J.P. Zhou, X.M. Chen, X.B. Bian, P. Liu, J. Ceram. Process. Res. 13 (2012) 110. [38] A.K.M. Akther Hossain, S.T. Mahmud, M. Seki, T. Kawai, H. Tabata, J. Magn. Magn. Mater. 312 (2007) 210. [39] M. Rahimi, P. Kameli, M. Ranjbar, H. Hajihashemi, H. Salamati, J. Mater. Sci. 48 (2013) 2969. [40] Y. Liu, Z. Yi, J.J. Li, T. Qiu, F.F. Min, M.X. Zhang, J. Am. Ceram. Soc. 96 (2013) 151. [41] P. Gao, X. Hua, V. Degirmenci, D. Rooney, M. Khraisheh, R. Pollard, R. M. Bowman, E.V. Rebrov, J. Magn. Magn. Mater. 348 (2013) 44. [42] A.C.F.M. Costa, M.R. Morelli, R.H.G.A. Kiminami, J. Mater. Sci. 42 (2007) 779. [43] M. Jalaly, M.H. Enayati, P. Kameli, F. Karimzadeh, Physica B 405 (2010) 507. [44] C. Krishnamoorthy, K. Sethupathi, V. Sankaranarayanan, Mater. Lett. 61 (2007) 3254. [45] C. Zener, Phys. Rev. 82 (1951) 403. [46] H.Y. Hwang, S.W. Cheong, P.G. Radaelli, Phys. Rev. Lett. 75 (1995) 914. [47] Y. Moritomo, H. Kuwahara, Y. Tomioka, Phys. Rev. B 55 (1999) 7549. [48] B. Munirathinam, M. Krishnaiah, U. Devarajan, S.E. Muthu, S. Arumugam, J. Phys. Chem. Solids 73 (2012) 925. [49] B. Munirathinam, M. Krishnaiah, M.M. Raja, S. Arumugam, K. Porsezian, Bull. Mater. Sci. 34 (2011) 121. [50] B. Munirathinam, M. Krishnaiah, S. Arumugam, M.M. Raja, J. Phys. Chem. Solids 71 (2010) 1763. [51] U.V. Chhaya, R.G. Kulkarni, Mater. Lett. 39 (1999) 91. [52] Alex Goldman, Modern Ferrite Technology, 2nd ed., Springer, New York, 2006. [53] P. Kumar, S.K. Sharma, M. Knobel, M. Singh, J. Alloy. Compd. 508 (2010) 115. [54] Md. D. Rahaman, S.K. Saha, T.N. Ahmed, D.K. Saha, A.K.M. Akther Hossain, J. Magn. Magn. Mater. 371 (2014) 112. [55] J.B. Nelson, D.P. Riley, Proc. Phys. Soc. 57 (1945) 160. [56] M.S. Khandekar, R.C. Kambale, J.Y. Patil, Y.D. Kolekar, S.S. Suryavanshi, J. Alloy. Compd. 509 (2011) 1861. [57] P.B. Belavi, G.N. Chavan, L.R. Naik, V.L. Mathe, R.K. Kotnala, Int. J. Sci. Technol.

Res. 2 (2013) 298. [58] J.C. Maxwell, Electricity and Magnetism, Oxford University Press, London, 1973. [59] C.G. Koops, Phys. Rev. 83 (1951) 121. [60] Y. Liu, Y. Wu, D. Li, Y. Zhang, J. Zhang, J. Yang, J. Mater. Sci.: Mater. Electron. 24 (2013) 1900. [61] B.K. Kunar, G.S. Srivastava, J. Appl. Phys. 75 (1994) 6115. [62] A. Sharma, R.K. Kotnala, N.S. Negi, J. Alloy. Compd. 582 (2014) 628. [63] D. Ravinder, K.V. Kumar, Bull. Mater. Sci. 5 (2001) 505. [64] A.R. James, C. Prakash, G. Prasad, J. Phys. D: Appl. Phys. 39 (2006) 1635. [65] E. Rezlescu, N. Rezlescu, P.D. Popa, L. Rezlescu, C. Pasnicu, Phys. Status Solidi A 162 (1997) 673. [66] B. Behera, P. Nayak, R.N.P. Choudhary, Cent. Eur. J. Phys. 6 (2008) 289. [67] J. Plocharski, W. Wieczoreck, Solid State Ion. 28–30 (1988) 979. [68] M.H. Khan, S. Pal, E. Bose, Cand, J. Phys. 91 (2013) 1029. [69] A.K. Roy, A. Singh, K. Kumari, K. Prasad, A. Prasad, IOSR J. Appl. Phys. 3 (2013) 47. [70] S.G. Doh, E.B. Kim, B.H. Lee, J.H. Oh, J. Magn. Magn. Mater. 272–276 (2004) 2238. [71] M. Hashim, Alimuddin, S.E. Shirsath, S. Kumar, R. Kumar, A.S. Roy, J. Shah, R. K. Kotnala, J. Alloy. compd. 549 (2013) 348. [72] H. Inaba, J. Mater. Sci. 32 (1997) 1867. [73] M.H. Abdullah, A.N. Yusuff, J. Mater. Sci. 32 (1997) 5817. [74] N.V. Prasad, G. Prasad, T. Bhimasankaram, S.V. Suryanarayana, G.S. Kumar, Bull. Mater. Sci. 24 (2001) 487. [75] C.E. Ciomaga, A.M. Neagu, M.V. Pop, M. Airimioaei, S. Tascu, G. Schileo, C. Galassi, L. Mitoseriu, J. Appl. Phys. 113 (2013) 074103. [76] P.B. Macedo, C.T. Moynihan, R. Bose, Phys. Chem. Glasses 13 (1972) 171. [77] E. Boucher, B. Guiffard, L. Lebrun, D. Guyomar, Ceram. Int. 32 (2006) 479. [78] S. Panteny, C.R. Bowen, R. Stevens, J. Mater. Sci. 41 (2006) 3845. [79] O. Bidault, P. Goux, M. Kchikech, M. Belkaoumi, M. Maglione, Phys. Rev. B 49 (1994) 7868. [80] K. Verma, A. Kumar, D. Varshney, J. Alloy. Compd. 526 (2012) 91. [81] A.K. Jonscher, Nature 267 (1977) 673. [82] L.P. Curecheriu, M.T. Buscaglia, V. Buscaglia, L. Mitoseriu, P. Postolache, A. Ianculescu, P. Nanni, J. Appl. Phys. 107 (2010) 104106. [83] I.G. Austin, N.F. Mott, Adv. Phys. 18 (1969) 41. [84] D. Alder, J. Fienleib, Phys. Rev. B 2 (1970) 3112. [85] J. Slama, P. Siroky, M. Soka, A. Gruskova, V. Jancarik, J. Electr. Eng. 60 (2009) 39. [86] M. Yu, J. Hu, J. Liu, S. Li, J. Magn. Magn. Mater. 326 (2013) 31. [87] X. Lu, G. Liang, Y. Zhang, W. Zhang, Nanotechnology 18 (2007) 015701. [88] J. Shen, Y. Bai, J. Zhou, L. Li, J. Am. Ceram. Soc. 88 (2005) 3440. [89] L. Jia, S. Chen, H. Zhang, Z. Zhong, Mater. Chem. Phys. 114 (2009) 697. [90] X. Qi, J. Zhou, Z. Yue, Z. Gui, L. Li, S. Buddhudu, Adv. Funct. Mater. 14 (2004) 920. [91] J.L. Snoek, Physica 14 (1948) 207. [92] J. Wei, J. Liu, S. Li, J. Magn. Magn. Mater. 312 (2007) 414. [93] K. Overshott, IEEE Trans. Magn. 17 (1981) 2698. [94] Y.H. Tang, X.M. Chen, Y.J. Li, X.H. Zheng, Mater. Sci. Eng. B 116 (2005) 150. [95] S.U. Durgadsimi, S.S. Chougule, B.K. Chougule, C.H. Bhosale, S.S. Bellad, Mater. Chem. Phys. 131 (2011) 199. [96] V.M. Petrov, G. Srinivasan, U. Laletsin, M.I. Bichurin, D.S. Tuskov, N. Paddubnaya, Phys. Rev. B 75 (2007) 174422. [97] Ce-Wen Nan, Phys. Rev. B 50 (1994) 6082. [98] C.E. Ciomaga, M. Airimioaei, V. Nica, L.M. Hrib, O.F. Caltun, A.R. Iordan, C. Galassi, L. Mitoseriu, M.N. Palamaru, J. Eur. Ceram. Soc. 32 (2012) 3325. [99] A.S. Fawzi, A.D. Sheikh, V.L. Mathe, Mater. Res. Bull. 45 (2010) 1000. [100] B.K. Bammannavar, L.R. Naik, J. Magn. Magn. Mater. 324 (2012) 944. [101] S.S. Chougule, B.K. Chougule, J. Alloy. Compd. 456 (2008) 441. [102] R. Rani, J.K. Juneja, S. Singh, K.K. Raina, C. Prakash, Ceram. Int. 39 (2013) 7845. [103] A.D. Sheikh, V.L. Mathe, Smart Mater. Struct. 18 (2009) 065014. [104] Q.H. Jiang, Z.J. Shen, J.P. Zhou, Z. Shi, Ce-Wen Nan, J. Eur. Ceram. Soc. 27 (2007) 279. [105] Y.J. Li, X.M. Chen, Y.Q. Lin, Y.H. Tang, J. Eur. Ceram. Soc. 26 (2006) 2839. [106] Y.K. Fetisova, K.E. Kamentsev, A.Y. Ostashchenko, G. Srinivasan, Solid State Commun. 132 (2004) 13.

Please cite this article as: Md.D. Rahaman, et al., Journal of Magnetism and Magnetic Materials (2015), http://dx.doi.org/10.1016/j. jmmm.2015.03.024i

67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95 96 97 98 99 100 101 102 103 104 105 106 107 108 109 110 111 112 113 114 115 116 117 118 119 120 121 122 123 124 125 126 127 128 129 130 131 132