Ac electrical conductivity and magnetic properties of BiFeO3–CoFe2O4 nanocomposites

Journal of Alloys and Compounds 599 (2014) 32–39

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Ac electrical conductivity and magnetic properties of BiFeO3–CoFe2O4 nanocomposites H.B. Sharma a, K. Nomita Devi a, V. Gupta b, J.H. Lee c, S. Bobby Singh a,⇑ a

Department of Physics, Manipur University, Canchipur 795003, India Department of Physics and Astrophysics, University of Delhi, Delhi 110017, India c Department of BIN Fusion Technology Chonbuk National University, Jeonju, Jeonbuk 561-756, South Korea b

a r t i c l e

i n f o

Article history: Received 20 September 2013 Received in revised form 29 January 2014 Accepted 6 February 2014 Available online 15 February 2014 Keywords: Nanocomposite Ceramics Dielectric Antiferromagnetic Multiferroic

a b s t r a c t In this paper we report the synthesis of multiferroic BiFeO3 (BFO)–CoFe2O4 (CFO) nanocomposites. The XRD and SEM data confirm the formation of both the parent compounds and the composite. All the composite samples show typical frequency dispersion at low frequency. The composite samples possess low dielectric loss. The frequency dependence of ac electrical conductivity curves of the composite samples were fitted using the relation rðf Þ ¼ rdc þ AðTÞf nðTÞ . It is observed that at lower frequency range the curve is not well fitted suggesting the contribution of small ac conductivity due to electrode polarization effect along with dc conductivity; however at higher frequency the curves are well fitted confirming the frequency dependence of ac conductivity at higher frequency. Frequency dependent of ac conductivity indicates that conduction occurs due to hopping of small polarons among the localized states. BFO shows weak ferromagnetism with coercivity, Hc = 374 Oe and magnetization, Ms = 0.06 emu/g at 21 kOe. All the composite samples show ferrimagnetic nature and found that there is a decrease in both saturation and remanence magnetization with increase in BFO content. However, there is a variation in coercivity with increase in BFO content and the sample 0.8(BFO)–0.2(CFO) (BF8–CF2) shows maximum Hc. Ó 2014 Elsevier B.V. All rights reserved.

1. Introduction Multiferroic materials exhibit two or three of the properties of ferroelectricity, ferromagnetism and ferroelasticity in the same phase. In recent years, this class of materials has attracted extensive interest in the field of solid state physics and in numerous technological fields. Cross-coupling between the magnetic and electrical orders, termed as magnetoelectric (ME) coupling, is of particular interest [1]. This coupling enables electrical polarization to be controlled using a magnetic field and, conversely, the manipulation of magnetization by varying an electrical field [2–4]. BiFeO3 (BFO) is a promising candidate as a single-phase multiferroic material because it is antiferromagnetic with a relatively high Neel temperature (TN = 380 °C) and ferroelectric with a high Curie temperature (TC = 810 °C) [5]. BFO and BFO-based single-phase multiferroics have been widely studied in both ceramic and film forms [6–10]. However, in spite of their excellent ferroelectric property, BFO-based single-phase multiferroics generally have weak macroscopic magnetic properties. Although relatively enhanced magnetization can be observed in highly constrained epitaxial films or chemically substituted films and ceramics, the ⇑ Corresponding author. Tel.: +91 385 2435833; fax: +91 385 2435831. E-mail address: [email protected] (S. Bobby Singh). http://dx.doi.org/10.1016/j.jallcom.2014.02.024 0925-8388/Ó 2014 Elsevier B.V. All rights reserved.

enhancement is limited. This leads to weak ME effect in these materials. This problem prevents the application of BFO-based single-phase multiferroics as large ME effect is desirable for real applications. The reason for low magnetization of BFO-based single-phase multiferroics is the weak magnetization due to the intrinsic spatially modulated, incommensurate cycloidal spin structure and antiferroelectric nature in BFO [11]. Although chemical substitution and strain engineering are effective ways to suppress or destroy the spin cycloid and thus enhance magnetization and ME effect to some extent, they cannot change the intrinsic antiferroelectric nature of BFO-based single-phase multiferroics [6–9]. One effective way of increasing the magnetic property of the BFO based multiferroics is to introduce suitable ferrites to form composites. In this paper, we report the synthesis of BiFeO3, CoFe2O4 and (1  x)(BiFeO3)–x(CoFe2O4) nanocomposite powders. The structural, dielectric and magnetic properties of the composite ceramics have been studied.

2. Experiment Multiferroic nanocomposite powders of (1  x)(BFO)–x(CFO) with x = 0.0, 0.2, 0.5, 0.7 and 1.0 were prepared by sol–gel auto-combustion method. Nanocomposite powders were synthesized using analytically pure grade bismuth nitrate

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Intensity (arb.unit)

(311)

BF3-CF7

30

50

(214)

(018)

(116)

(202) (006)

* *

*

40

BFO

(012)

* *

*

BF1

(122)

BF2

BF8-CF2

(024)

*

*

BF3

BF5-CF5

(104) (110)

*

BF4

*

20

CFO (222)

(220)

(214)

(018)

(116)

(122)

(024)

(006) (202)

(110)

BF5

*

Intensity (arb.unit)

(012)

(104)

Fig. 1 shows the XRD patterns of BiFeO3 powders annealed at different temperatures such as 300 °C (as-fired), 400 °C, 500 °C and 550 °C respectively. The samples will be respectively called as BF1, BF2, BF3 and BF4, respectively hereafter. All the samples show characteristic XRD patterns of BiFeO3 with the presence of some impure phases such as Bi2Fe4O9, Bi46Fe2O72 and Bi24Fe2O39 (as indicated by * mark). It is observed from Fig. 1 that as the annealing temperature increases, the intensity of the impurity peaks decreases, while the XRD pattern of the characteristic BiFeO3 peaks became more intense and sharper, showing enhanced crystallinity of the sample. It is also observed that at 550 °C almost all the impurity phases of BiFeO3 were melt out, only the impure phase at 2h = 28.02 still present with a very low intensity. To remove the unwanted impure phases of bismuth ferrite, the sample BF4 was then leached with dilute nitric acid (60%) followed by washing with deionized water. BF5 shows the XRD pattern of the leached BF4 sample and observed that the leached powder developed a polycrystalline single-phase BiFeO3 powder. Fig. 2 shows the XRD patterns of the (1  x)(BFO)–x(CFO) nanocomposite samples annealed at 700 °C for 4 h in free air. As evident

(511)

3.1. Structural properties

(422)

3. Results and discussions

from the figure, the BFO powder shows polycrystalline singlephase perovskite rhombohedral structure with the preferred orientation along (1 1 0) plane. The XRD peaks were indexed using ICDD Card No. 82-1254. The diffraction pattern of the CFO powder confirms the formation of cubic spinel type lattice of CoFe2O4 (space group Fd3m), which matches well with the standard XRD pattern (ICDD card No. 22-1086). Using Miller indices of the respective planes, the lattice parameters a, b and c of the unit cell were evaluated and the average crystallite size of the samples were determined from the full width at half maximum (FWHM) of the major peak using the Scherrer’s formula and are reported in Table 1. For the composite samples, only the peaks of perovskite BFO and spinel CFO are present without any additional peaks or phase including untreated oxides or impurities. This means that there is no reaction between BFO and CFO [10,11]. A slight change in the peak positions from that of pure BFO and CFO is observed. It is observed that for composition x = 0.3 and 0.7, the corresponding XRD peaks of BFO and CFO shift to smaller 2h values while for 50–50 composition the peaks are more or less same to their parents compounds. This might suggests a small lattice distortion from the rhombohedral structure of BiFeO3 occurring with the presence of CoFe2O4. The lattice parameters for the composite sample are reported in Table 1. The composite samples have smaller BFO grain sizes than that of the pure BFO powder. Fig. 3 shows the SEM images of (1  x)(BFO)–x(CFO) nanocomposite powders annealed at 550 °C. From the images it can be seen that all the samples have a narrow size distribution of particles. Some moderately agglomerated particles are also present in the images. It can be hardly to determine the parent compounds from the SEM images. Fig. 4 shows the SEM images of (1  x)(BFO)–x(CFO) composite ceramics, made on the cross section of the ceramic samples. From the SEM images it can be seen that pure BFO grains show rhombohedral in structure with different sizes. It is evident from the SEM image of CFO that the grains have a narrow size distribution of particles. From the SEM images it is observed that the composite samples have smaller grain sizes than that of the parent compounds. Energy dispersive analysis of X-ray (EDAX) study was carried out and the typical EDAX spectra for (1  x)BFO–x(CFO) are presented in Fig. 5. Energy dispersive X-ray results confirmed the ratio of the transition metal atoms in each material according to the nominal stoichiometry. The atomic ratio of Fe:Bi, Fe:Co and Bi:Co for the BFO, CFO,BF8–CF2, BF5–CF5 and BF3–CF7 samples (Fig. 5) are maintained at 1:1, 2:1 and 8:2, 1:1 and 3:7 respectively.

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pentahydrate [Bi(NO3)35H2O], cobalt nitrate hexahydrate [Co(NO3)26H2O], iron nitrate nonahydrate [Fe(NO3)39H2O] and citric acid [C6H8O7] as chemical reagents. Stoichiometric amount of bismuth nitrate say (1  x) mole, cobalt nitrate (x) mole and iron nitrate [(1  x) + 2x] mole were dissolved in deionized water under constant stirring on a hotplate at 50 °C. Citric acid of 1:1 M ratio with total metal nitrate was then added to the above mixture solution to chelate the metal ions. The solution was then allowed to evaporate on a hot plate maintaining the solution temperature at 80–90 °C until the formation of gel. The temperature was then increased (T > 250 °C) which leads the gels to convert directly into the corresponding nanocrystalline composite powders. The resulting composite powders were grounded to get fine powder samples. The grinded powders were then calcined at 550 °C in air for 4 h, respectively followed by leaching with dilute nitric acid. The powders were then pressed to form pellets of 15.10 mm diameter and 2 mm thickness by applying a pressure of 300 kg/cm2. The pellets were sintered at 700 °C for 4 h with a heating rate of 5 °C/min to make dense ceramics for electrical and magnetic measurements. The composite samples of (1  x)(BFO)–x(CFO) with x = 0.0, 0.2, 0.5, 0.7 and 1.0 will be respectively called as BFO, BF8-CF2, BF5-CF5, BF3-CF7 and CFO hereafter. The crystal structures of the samples were determined by X-ray diffraction (XRD) h–2h scan with scanning rate 2°/min using PANalytical X’PERT PRO diffractometer. Surface morphologies were determined using a FEI QUANTA 250 scanning electron microscope (SEM) with energy dispersive X-ray spectroscopy (EDAX). Dielectric properties of the samples were measured using an Agilent 4284A LCR meter. The room-temperature magnetization was determined as a function of field using a vibrating sample magnetometer.

60

2θ (Degrees) Fig. 1. X-ray diffraction pattern of BiFeO3 powders annealed at different temperatures.

20

30

40

50

60

2θ (Degrees) Fig. 2. X-ray diffraction pattern of (1  x)BFO–xCFO nanocomposites.

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H.B. Sharma et al. / Journal of Alloys and Compounds 599 (2014) 32–39

Table 1 Microstructure parameters of BFO–CFO nanocomposites. Sample

BFO BF8–CF2 BF5–CF5 BF3–CF7 CFO

Lattice constant

Crystallite size (nm) CFO

Crystallite size (nm) BFO

CFO

BFO

a (Å)

a (Å)

c (Å)

(XRD)

(XRD)

– 8.34 8.35 8.34 8.36

5.56 5.55 5.56 5.55 –

13.80 13.79 13.80 13.79 –

74.30 44.06 50.43 47.90 –

– 41.93 46.46 45.83 47.5

These results revealed the formation of BFO–CFO composite at low calcination temperature. 3.2. Dielectric properties Fig. 6(a) and (b) displays the variation of dielectric constant and dissipation factor for (1  x)(BFO)–x(CFO) samples as a function of frequency at room temperature from 100 Hz to 1 MHz. It is observed that the dielectric constant decreases steeply at lower frequencies and remains constant at higher frequencies, indicating

the usual dielectric dispersion. This may be attributed to the polarizations due to the changes in valence states of cations and space charge polarization (i.e. Maxwell–Wagner) [12–14], in agreement with Koop’s phenomenological theory, as the dielectric constant is a combined effect of dipolar, electronic, ionic and interfacial polarizations. It is also reported that the polarization in ferrites is mainly contributed by the space charge polarization, which is governed by the number of space charge carriers and the conductivity in materials [15] and hopping exchange of the charges between two localized states governed by density of the localized state and resultant displacement of charges w.r.t. the external field. So, the decrease in dielectric constant with increasing frequency may be attributed to the two factors (i) the electron exchange between Fe2+ and Fe3+ ions cannot follow the change of the external applied field beyond certain frequency [16] and (ii) the dipoles due to the space-charge polarization or Maxwell–Wagner type interfacial polarization do not respond at higher frequencies. At lower frequencies the grain boundaries are more effective or the contribution from the space charge polarization is high. The high dielectric loss at low frequency in this sample confirmed the presence of space charges at the grain boundaries. The space charge polarization is inherently related to non-uniform charge accumulation at grain boundaries mainly due to the vacancies of oxygen.

Fig. 3. SEM photograph of (1  x)BFO–xCFO nanocomposite powders annealed at 550 °C.

H.B. Sharma et al. / Journal of Alloys and Compounds 599 (2014) 32–39

35

Fig. 4. SEM photograph of (1  x)BFO–xCFO nanocomposite ceramic sintered at 700 °C.

Similar low values of the dielectric constant were found in Mn substituted Ni–Zn ferrites prepared by citrate method [17] and Cd, Cr substituted Co ferrite [18]. In composites, the higher value of dielectric constant at lower frequencies is associated with heterogeneous conduction [19], but sometimes the polaron hopping mechanism results in electronic polarization contributing to low frequency dispersion. The dielectric behavior in composites can also be explained on the basis of polarization mechanism in ferrites because the conduction beyond the phase percolation limits due to the presence of ferrite phase [20]. In the present work, the presence of Bi2+/Bi3+, and Co2+/ Co3+ ions would give rise to p-type carriers. The local displacement of p-type carriers in the external electric field direction contributes to the net polarization in additional to that of n-type carriers. However, the p-type carrier contribution is smaller than that from the electronic exchange between Fe3+ and Fe2+ ions and it is opposite in sign [21]. Since the mobility of p-type carriers is smaller than n-type carriers, their contribution to polarization decreases more rapidly and then decreases with increase in frequency as it is observed in the present compositions. The dielectric loss factor is considered to be the most important part of the total core loss in ferrites [17,18]. From Fig. 6(b), it can be

seen that the dielectric loss possesses similar behavior to that of dielectric constant (e0 ) with respect to frequency. The dielectric loss originates from two mechanisms: resistive loss and relaxation loss. In the resistive loss mechanism, energy is consumed by mobile charges in the ceramic. In the case of relaxation loss mechanism, it is relaxation of the dipoles that dissipates the energy. At lower frequencies, the maximum dielectric loss may be attributed to the fact that the hopping frequency of electrons between ferrous and ferric ions at adjacent octahedral sites is same as the period of applied field. At higher frequencies the hopping frequency of the electron exchange between ferrous and ferric ions cannot follow applied field beyond certain critical frequency and the loss is minimum. For comparison, the values of e0 and tan d at three frequencies for BFO–CFO composite ceramics are presented in Table 2. 3.3. AC conductivity To understand the conduction mechanism in composites the ac conductivity measurement was carried out at room temperature in the frequency range from 100 Hz to 1 MHz. Fig. 7 shows the variation of ac conductivity (rac) with frequency for BFO–CFO

36

H.B. Sharma et al. / Journal of Alloys and Compounds 599 (2014) 32–39

Fig. 5. EDAX spectra of (1  x)BFO–xCFO nanocomposites.

composites at room temperature. At low frequency range the ac conductivity is nearly independent of the frequency and has been attributed to the dc conductivity of the sample. With increasing frequency the ac conductivity increases, at first slowly and then rapidly at about 100 kHz. The conductivity r(f) at a particular temperature over a wide range of frequencies can be expressed as [22–24],

rðf Þ ¼ rdc þ rac ðf Þ ¼ rdc þ AðTÞf nðTÞ

ð1Þ

where rdc is the dc conductivity (frequency independent plateau in the low frequency region), A(T) and n(T) are constants at a certain temperature. The rac(f) versus ln (f) curves for BFO–CFO composite samples were fitted according to Eq. (1) and the fitting curves are shown as continuous line in Fig. 7. The values of rdc, A(T) and n(T) can be obtained from the fitted curves and are reported in Table 2. It is observed that at lower frequency range the curves are not well fitted suggesting the contribution of small ac conductivity due to

electrode polarization effect along with dc conductivity; however at higher frequency the curves are well fitted confirming the frequency dependence of ac conductivity at higher frequency. Frequency dependence of ac conductivity indicates that conduction occurs due to hopping of small polarons among the localized states. Hopping conduction is favored in ionic lattice in which the same kind of cations exists in two oxidation states [25]. Also, for ionic solids, the concept of small polaron conduction is valid. While Alder and Fienleib [26], have shown that the conduction in ferrite is due to hopping of charges which gives rise to the linearity between ac conductivity and angular frequency and as already stated, the concept of small polaron is valid in ferrites. The electrical conductivity is due to migration of ions and this ionic transport depends on the angular frequency. Thus it has been shown that the ac conductivity is proportional to angular frequency. The treatment of conduction by polaron is discussed by Austin and Mott [27]. In large polaron model, ac conductivity is due to band mechanism at all temperatures and decreases with increase in frequency. In

37

H.B. Sharma et al. / Journal of Alloys and Compounds 599 (2014) 32–39

(a)

300.0µ

BFO BF8-CF2 BF5-CF5 BF3-CF7 CFO

BFO BF8-CF2 BF5-CF5 BF3-CF7 CFO fitted curve

250.0µ -1

700

ac conductivity (σacSm )

Dielectric constant (ε')

800

600 500 400 300 200

200.0µ 150.0µ 100.0µ 50.0µ 0.0

100

1k

10k

100k

100

1M

1k

(b)

BFO BF8-CF2 BF5-CF5 BF3-CF7 CFO

14.0µ

0.6

dc conductivity n

-1

0.2

0.0 10k

100k

0.96

10.0µ 0.92

8.0µ

n

0.4

1k

1M

1.00

12.0µ

100

100k

Fig. 7. Variation of ac conductivity with frequency at room temperature for (1  x)BFO–xCFO nanocomposites.

dc conductivity (σdcSm )

Dissipation factor (tanδ)

0.8

10k

Frequency (Hz)

Frequency (Hz)

6.0µ

0.88

4.0µ 0.84 2.0µ

1M

Frequency (Hz)

0.80 0.0

0.2

Fig. 6. Variation of (a) dielectric constant and (b) dissipation factor with frequency for (1  x)BFO–xCFO nanocomposites.

case of small polaron model, the ac conductivity increases with increase in frequency [28]. All the plots at higher frequency are linear confirming that the conduction in all composites is due to small polaron hopping. The compositional dependence of fitting parameters of rdc and n is shown in Fig. 8. It is observed that BFO sample shows high dc conductivity while the composite sample BF8–CF2 and BF5–CF5 show small dc conductivity. The compositional dependence of n is just opposite to the dc conductivity showing high values of n for composite samples BF8–CF2 and BF5–CF5. Composition dependence of ac conductivity for composite samples was studied at different frequencies. The variation of conductivity with composition for composite samples at room temperature is shown in Fig. 9. It is observed that at low frequency, 10 kHz, the ac conductivity is almost independent of composition, however at higher frequency above 100 kHz, the composition

0.4

0.6

0.8

1.0

Composition (x) Fig. 8. Variation of fitting parameters rdc and n with composition for (1  x)BFO– xCFO nanocomposites.

dependence of ac conductivity is very clear. It is found that the ac conductivity for BiFeO3 (i.e. x = 0) is greater than the other compositions. As the conductivity in ferrites is mainly due to the hole hopping in metallic ions (Bi ions) and electron hopping in ferrous ions on the octahedral sites, in the BFO–CFO composites when x = 0, maximum number of Bi2+ and Bi3+ ions are available for hole hopping. Hence the main contribution of conductivity may be due to hole and electron hopping. As the concentration of CoFe2O4 (x) increases, the number of Bi2+ and Bi3+ ions on B sites might decrease. Also small decrease in conductivity for composition x = 0.2 and x = 0.5 is observed and this may be due to electron–hole compensation at the B site. Further increase of CFO concentration diminishes the hole hopping by decreasing the no. of Bi2+ and Bi3+ ions on the B site. Thus the electron hopping becomes

Table 2 Dielectric and electrical properties of BFO–CFO nanocomposites. Sample

BFO BF8–CF2 BF5–CF5 BF3–CF7 CFO

Dielectric constant (e0 )

Dissipation factor (tan d)

10 kHz

100 kHz

1 MHz

10 kHz

100 kHz

1 MHz

269 231 259 264 203

252 225 251 251 196

248 225 250 249 195

0.121 0.026 0.031 0.055 0.041

0.041 0.015 0.014 0.024 0.019

0.020 0.011 0.010 0.014 0.011

rac (lS m1) 1 MHz

rdc (lS m1)

A(T) (106)

n(T)

277 143 146 192 124

12.5 2.1 3.1 5.3 3.9

30.3 2.06 2.66 9.43 3.06

0.82 0.97 0.95 0.88 0.93

H.B. Sharma et al. / Journal of Alloys and Compounds 599 (2014) 32–39

250.0µ 200.0µ 150.0µ 100.0µ 50.0µ

40

Magnetization (emu/gm)

10kHz 100 kHz 500 kHz 1 MHz

-1

ac conductivity (σac Sm )

300.0µ

30 20

Magnetization (emu/gm)

38

BFO

CFO

0.06 0.04 0.02

-30k

-20k

0.00 -10k 0 -0.02

10k

20k

BF3-CF7

30k

BF5-CF5

-0.04 -0.06

Applied Magnetic Field (Oe)

10

BF8-CF2 BFO

0 -10 -20 -30 -40

0.0 0.0

0.2

0.4

0.6

0.8

1.0

-20k

Composition (x)

-10k

0

10k

20k

Applied Magnetic Field (Oe)

Fig. 9. Variation of ac conductivity with composition for (1  x)BFO–xCFO nanocomposites at different frequency.

predominant and it increases the conductivity. Thus a high conductivity for BF3–CF7 is observed. For x = 1, that is for CoFe2O4 the conductivity again decreases. 3.4. Magnetic properties A great deal of information about the magnetic properties of a material can be obtained by studying its hysteresis loop. Magnetic properties of a material are strongly dependent on chemical composition, sintering temperature, grain size, crystal structure and porosity of the material [29]. Ferrite being magnetic materials, it becomes necessary to study the magnetic properties of these ferrites individually and in composite forms. Fig. 10 shows the magnetization M versus applied field H of BFO–CFO composite samples at room temperature. BFO shows weak ferromagnetism with coercivity, Hc = 374 Oe and magnetization, Ms = 0.06 emu/g at 21 kOe (inset in Fig. 10). CFO exhibits typical ferrimagnetic behavior with a saturation magnetization Ms and coercivity Hc of 40.4 emu/g and 1.8 kOe, respectively. The obtained Ms-value of CFO is comparable with those earlier reported [30–33]. The high coercivity and low saturation magnetization in our CFO sample might be due to small grain size. It is observed that all the composite samples show ferrimagnetic nature, which is similar to CFO sample and there is a decrease in both saturation and remanence magnetization, Mr with increase in BFO content. However, there is a variation in coercivity with increase in BFO content and the sample BF8–CF2 shows maximum Hc. The values of saturation magnetization (Ms), remanence magnetization (Mr), coercivity (Hc) and squareness (S = Mr/Ms) of the loop for BFO–CFO composite sample are reported in Table 3 for comparison. To see whether the magnetization in our composite samples is solely from CFO, we calculate the magnetization (M 0s ) of each composite by multiplying the magnetization (Ms) of CFO with their respective weight % of CFO present in each sample and it is found that the calculated

Fig. 10. Room temperature M–H loop of (1  x)BFO–xCFO nanocomposites.

values of magnetization (M 0s ) are lower than those obtained from the M–H loop (Table 3). Accordingly, the magnetic moment of the composite is unlikely to arise only from CFO. The enhanced magnetization in our composite samples may be due to two factors (i) although pure BFO is a canted antiferromagnet with a weak magnetization of 0.06 emu/g at 21 kOe, the spin orientations of the BFO grains are very likely changed by the strain in the grain boundaries between the BFO and CFO in the composite samples and (ii) magnetic interaction between the two compounds cause the increase of the magnetization of the composite, i.e. the compounds are exchange coupled. Thus the magnetic properties of the BFO can be enhanced by making composite with CFO. 4. Conclusion In summary, multiferroic composites consisting of BFO and CFO were successfully synthesized using the sol–gel method. The XRD and SEM data confirm the formation of both the parent compounds and the composite. All the composite samples show typical frequency dispersion at low frequency and the composites BF5–CF5 and BF3–CF7 exhibits high dielectric constant. The composite samples possess low dielectric loss. The frequency dependence of ac electrical conductivity curves of the composite samples were fitted using the relation rðf Þ ¼ rdc þ AðTÞf nðTÞ . It is observed that at lower frequency range the curve is not well fitted suggesting the contribution of small ac conductivity due to electrode polarization effect along with dc conductivity; however at higher frequency the curves are well fitted confirming the frequency dependence of ac conductivity at higher frequency. Frequency dependent of ac conductivity indicates that conduction occurs due to hopping of small polarons among the localized states. BFO shows weak ferromagnetism with coercivity, Hc = 374 Oe and magnetization, Ms = 0.06 at 21 kOe. All the composite samples show ferrimagnetic nature and found that there is a decrease in both saturation and

Table 3 Magnetic properties of BFO–CFO nanocomposites. Sample

Equivalent weight % of CFO

Saturation magnetization (Ms) (emu/g)

Remanence magnetization (Mr) (emu/g)

Coercivity (Hc) (Oe)

Squareness (S = Mr/Ms)

Calculated magnetization (M 0s ) (emu/g)

BFO BF8–CF2 BF5–CF5 BF3–CF7 CFO

0 15.78% 42.8% 63.62% 100%

0.06 11.3 24.3 25.9 40.4

0.002 5.62 12.31 13.40 20.59

374 2557 2348 2534 1841

0.033 0.497 0.506 0.517 0.509

– 6.34 17.29 25.69 –

H.B. Sharma et al. / Journal of Alloys and Compounds 599 (2014) 32–39

remanence magnetization with increase in BFO content. However, there is a variation in coercivity with increase in BFO content and the sample BF8–CF2 shows maximum Hc.

[12] [13] [14] [15]

Acknowledgements

[16]

One of the authors, S. Bobby Singh is thankful to the CSIR, New Delhi, for providing financial assistance as Research Associateship. The authors are also thankful to the Department of Physics and Astrophysics, University of Delhi for extending the facility for magnetic measurements. References [1] W. Eerenstein, M.D. Mathur, J.F. Scott, Nature (London) 442 (2006) 759–765. [2] M. Fiebig, Th. Lottermoser, D. Fröhlich, A.V. Goltsev, R.V. Pisarev, Nature (London) 419 (2002) 818–820. [3] J. Wang, J.B. Neaton, H. Zheng, V. Nagarajan, Science 299 (2003) 1719–1722. [4] M. Fiebig, J. Phys. D 38 (2005) R123–R152. [5] P. Monica, D. Crespo, M. Jose, C.-M.S. Preda, V. Fruth, J. Am. Ceram. Soc. 90 (9) (2007) 2723–2727. [6] Y.K. Jun, W.T. Moon, C.M. Chang, H.S. Kim, H.S. Ryu, J.W. Kim, K.H. Kim, H.S. Hong, Solid State Commun. 135 (2005) 133. [7] G.L. Yuan, S.W. Or, Y.P. Wang, Z.G. Liu, J.M. Liu, Solid State Commun. 138 (2006) 76–81. [8] Y.K. Jun, S.H. Hong, Solid State Commun. 144 (2007) 329–333. [9] S.T. Zhang, L.H. Pang, Y. Zhang, M.H. Lu, Y.F. Chen, J. Appl. Phys. 100 (2006) 114108. 6pp.. [10] H. Zheng, Q. Zhan, F. Zavaliche, M. Sherburne, F. Straub, M.O. Cruz, L.Q. Chen, U. Dahmen, R. Ramesh, Nano Lett. 6 (2006) 1401–1407. [11] L. Yan, Z. Wang, Z. Xing, J. Li, D. Viehland, J. Appl. Phys. 107 (2010) 064106. 5pp..

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