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Improvement of room temperature electric polarization and ferrimagnetic properties of Co1.25Fe1.75O4 ferrite by heat treatment
Author’s Accepted Manuscript Improvement of room temperature electric polarization and ferrimagnetic properties of Co1.25Fe1.75O4 ferrite by heat treatment R.N. Bhowmik, A.K. Sinha www.elsevier.com/locate/jmmm
PII: DOI: Reference:
S0304-8853(16)31674-2 http://dx.doi.org/10.1016/j.jmmm.2016.08.014 MAGMA61699
To appear in: Journal of Magnetism and Magnetic Materials Received date: 14 April 2016 Revised date: 22 June 2016 Accepted date: 3 August 2016 Cite this article as: R.N. Bhowmik and A.K. Sinha, Improvement of room temperature electric polarization and ferrimagnetic properties of Co1.25Fe1.75O ferrite by heat treatment, Journal of Magnetism and Magnetic Materials, http://dx.doi.org/10.1016/j.jmmm.2016.08.014 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Improvement of room temperature electric polarization and ferrimagnetic properties of Co1.25Fe1.75O4 ferrite by heat treatment R.N. Bhowmika*, and A.K. Sinhab a
Department of Physics, Pondicherry University, R.V. Nagar, Kalapet, Pondicherry -605014,
India b
Indus Synchrotrons Utilization Division, Raja Ramanna Centre for Advanced Technology,
Indore- 452013, India *
Corresponding author (RNB): Tel.: +91-9944064547; Fax: +91-413-2656988, E-mail: [email protected]
Abstract We report the improvement of electric and magnetic properties of the chemical routed samples of Co1.25Fe1.75O4 ferrite. Synchrotron X-ray diffraction pattern confirmed the cubic spinel structure in all samples. Raman spectroscopy provided the information of mixed-type cubic spinel structure, where Co2+ and Fe3+ ions order in octahedral and tetrahedral sites. The thermal heat treatment has brought a systematic change in electric polarization and magnetic parameters. At a critical annealing temperature (TAN) of 800 °C, the sample has shown the highest magnetic coercivity ~ 596 Oe and relatively large squareness ~ 0.5. The electrical conductivity, dielectric loss and dielectric constant have decreased with increase of TAN. The dielectric constant becomes stable for TAN above 800 °C. The samples annealed in the temperature range 900 0C -1000 0C have shown a good signature of ferroelectric polarization, although it is under saturated and leakage of polarization due to conductive effect is observed. This work shows the scope of the improvement of electric polarization due to capacitive effect on increasing the annealing temperature of as-prepared Co ferrite sample. Keywords: Magnetic materials; Chemical synthesis; X-ray diffraction; Raman spectroscopy; Ferroelectricity. 1
1. Introduction Cobalt substituted magnetite (CoxFe3-xO4) belongs to the class of cubic spinel structured ferrite with formula unit AB2O4, where A and B represent the tetrahedral and octahedral coordinated interstitial lattice sites for metal ions in oxygen environment [1]. The variation of Co content, and distribution and charge states of Co and Fe ions have played an important role for tuning the structure, ferrimagnetic, dielectric, and ferroelectric properties in Cobalt substituted ferrite [1-7]. The coexistence of ferrimagnetic and ferroelectric properties in ferrite structure could enhance its applications in many modern areas of microelectronics, data storage devices, microwave devices, and sensors [8-10]. The establishment of a coupling between ferroelectric and ferrimagnetic properties in Co ferrite may be an additional advantage for exploring the multi-functional applications of ferrites in non-volatile memory devices, magnetic and electric field controlled sensors, and spintronics [11]. We have noted that chemical routed samples of CoFe2O4 and Co1.5Fe1.5O4 ferrites formed single phased cubic spinel structure for a wide range of annealing temperature [4, 1213]. The Co and Fe ions in such single phased spinel ferrites can be distributed as (Co2+1-b Fe3+b)A[Co2+bFe3+3-x-bCo3+x-1]O4 ; 0 ≤ b ≤ 1 [14-15]. The distribution of magnetic ions (Co2+ highly anisotropic and Fe3+) in A and B sites of the spinel structure quantifies the magnetic moment (M = MB – MA, where MB and MA are the net moments in B and A sites) per formula unit of Co ferrite. Substitution of high spin Fe3+ moments at B sites by the low spin (nonmagnetic) Co3+ decreases the magnetization and magneto-crystalline anisotropy in CoxFe3xO4 ferrite
[2, 15]. The increase of Co atoms in CoxFe3-xO4 ferrite for x > 1.5 showed the
structural phase destabilization, randomization and re-configuration of the charge states of Co and Fe ions [2, 15-20]. The Co content in CoxFe3-xO4 ferrite also determines the nature of semiconductor properties in ferrite structure whether p type for x > 1 (hole hopping through Co2+-O2—Co3+ superexchange paths) or n type for x £ 1 (electron hopping through Fe3+-O2-
2
Fe2+ superexchange paths) [1, 14, 18]. Recently, some reports showed ferroelectric properties in CoFe2O4 ferrite [3, 21]. CoFe2O4 ferrite may exhibit reasonably large magnetization and coercivity, but it is a poor ferroelectric material with huge leakage of polarization. The ferroelectric polarization in Co ferrite is obviously small due to cubic structure with semiconductor nature and presence of magnetic ions. Hence, improvement of ferroelectric polarization with low leakage in ferrite is a challenging task. Such ferroelectric materials, different from traditional ferroelectric BaTiO3, may be classified as type –II multiferroics, where spontaneous electric polarization is extremely low, but there is a possibility of finite coupling between electric polarization and ferro/ferrimagnetic order [22]. Some of the recent works on Co rich ferrite [4, 6, 13, 17, 23-24] showed a lot of scope for achieving moderate magnetization, high squareness and coercivity with moderate dielectric constant and low dielectric loss. A detailed study of the dielectric, ferroelectric and ferrimagnetic properties in Co rich side (x > 1) of CoxFe3-xO4 ferrite could be useful for basic understanding and development ferrite based multiferroic materials. Taking a clue from the observation of ferroelectric and ferrimagnetic properties in Co rich spinel ferrites [24-25], this work is dedicated for searching the signature of ferroelectric properties in Co1.25Fe1.75O4 ferrite. Subsequently, efforts are made to improve the ferroelectric polarization and ferrimagnetic properties by adopting thermal treatment of the chemically routed as-prepared material. This ferrite composition belongs to slightly Co rich side (x > 1) of CoxFe3-xO4 ferrite and semiconductor in nature [14]. This ferrite composition is stabilized in single phase for annealing temperature up to 1000 0 C and excludes the effects of structural complexity as found for Co content more than 1.5 [15, 17, 19]. 2. Experimental Chemical co-precipitation route was used to prepare the Co1.25Fe1.75O4 ferrite. The stoichiometric amounts of high purity (> 99.999%) Co(NO3)2.6H2O and Fe(NO3)3.9H2O were
3
dissolved in distilled water to prepare a nitrate solution (pH value ~ 1.6). NaOH solution with pH ~ 13.5 was used as a precipitating agent and added drop wise to nitrate solution until pH reached to 11. The mixed solution was kept at 80 °C for 2 h with continuous stirring and pH ~ 11 is maintained during chemical reaction at 80°C by adding required amount of NaOH solution. Then, the resultant solution was cooled to room temperature and observed a black coloured precipitation. The transparent solution above the precipitate was carefully removed. The by-product mixed with the precipitate was removed by repeated washing processes. During washing process, sufficient amount of distilled water was added with stirring condition and the mixture was heated at 100 °C for 10-15 min. This was followed by cooling to room temperature and allowed to precipitate. The water solution above the precipitate was carefully removed. This washing process was repeated several times to ensure the removal of the by-products, mainly NaNO3 [13]. The absence of any white coloured powder due to formation of NaNO3 or non-absorbance of moisture from air indicated the black coloured ferrite nanoparticles free from by-products. The resultant black powder of ferrite was heated at 200 °C for 2 h and ground well to get black coloured ferrite powder. The ferrite (CF) powder was used to make several pellets of diameter 13 mm. The pellet from of the asprepared ferrite powder after 2 h air annealing (heating and cooling @ 5 °C/min) is denoted as CF20, CF50, CF60, CF80, CF90, and CF100 for annealing temperature (TAN) at 200 °C, 500 °C, 600 °C, 800 °C, 900 °C, and 1000 °C, respectively. The characterization and measurements of the samples were carried out at room temperature. The synchrotron X-ray powder diffraction (SXRD) measurement was performed in the 2θ range 5-40 ° using angle dispersive X-ray diffraction beam line (BL-12) at a wave length 0.7763 Å at Indus-2 synchrotron source, India [19]. The wavelength was calibrated using SXRD pattern of LaB6 NIST standard. We analysed the SXRD pattern of the samples using FULLPROF profile fit program. The Raman spectra of the samples were recorded
4
using Renishaw’s Invia Raman microscope (Renishaw, UK) with 488 nm laser light. Laser power was optimized to 3 mW on the sample surface for acquisition time 30 s. Surface morphology of the samples was checked using scanning electron microscope (SEM) (HITACHI S-3400N, Japan). Energy dispersive x-ray (EDX) analyser was used to estimate elemental composition. The dielectric properties of the samples were measured using a Broadband Dielectric Spectrometer (Novocontrol tech., Germany) in the frequency range 1 Hz to 7 MHz at an ac electrical signal of 1 V. For dielectric measurement, the disc shaped samples (Æ ~ 13 mm, t ~ 1 mm) were used and sand witched between two gold coated electrodes. Electric field dependent polarization (P-E) loop of the thin disc shaped samples (Æ ~ 13 mm, t~ 0.36 -0.42 mm) were measured using Precision Premire II loop tracer (Radiant Tech., USA). Magnetic hysteresis loop of the samples was recorded within field ±70 kOe using Physical properties measurement system (EC2-Quantum Design, USA). 3. Results and discussion 3.1. Structural properties Fig. 1 shows the SXRD patterns, along with profile fit data, of the samples. SXRD patterns confirmed the cubic spinel structure with space group Fd3m in all samples. A careful analysis of the SXRD peaks indicated asymmetric shape for the samples annealed at lower temperatures (200-500 0C), which could be fitted with secondary spinel phase. We have plotted the (440) SXRD peak alone in Fig. 2(a) to show the refinement of crystalline structure with the increase of TAN, because the diffraction peaks at higher 2θ are better choice to observe the second phase, if anything exists. The peaks achieved symmetric shape for TAN in the range 600-1000 0C. Since the SXRD pattern does not show the presence of any impurity phase like CoO or Fe doped CoO, as reported earlier [2, 16], the observed minor peak asymmetry for the samples (CF20 and CF50) prepared at low annealing temperature could be assigned to non-equilibrium spinel phase [13]. The lattice parameter shows a non-monotonic
5
increase with TAN. The lattice parameter for the samples with TAN up to 600 0C may be affected by the structural heterogeneity and non-equilibrium cubic spinel structure. The lattice parameter increases at higher TAN due to improvement of micro-structure (refinement of grain boundary, local re-ordering of Co and Fe ions among A and B sites) [4, 18]. The refinement of crystalline structure with increasing TAN is understood from the evolution of peak profile. Fig. 2(a) shows a broad peak profile for samples at lower TAN (200600 0C) in comparison with sharp peak profile for samples at higher TAN. The (440) peak profile of the samples are analysed by Gaussian curve equation. I = I + I ∗ e
(!)∗"
#$#% * ' &
(1)
In equation (1), xC (= 2θC) is the peak position, Ih is the peak intensity above the back ground intensity Io and β is the full width at half maximum (FWHM) of the peak intensity. These parameters from major peaks were used to calculate the grain size () and lattice strain (erms) of the samples using Williamson Hall equation. bcosθ- =
0.89*λ
+2εrms sinθ-
(2)
Here, λ is the wavelength of X-ray radiation (0.7763 Å). We have used b and θC values from profile fitting of seven high intensity peaks (111, 220, 311, 400, 422, 511, and 440) to plot the bcosθ- vs. sinθ- data (Fig. 2(b)). The linear intercept on the bcosθ- axis and its slope provided
0.89*λ
and 2εrms , respectively. Fig. 2(c) shows the variation of grain size () and
lattice strain (erms) with TAN. The grain size is in the range 12-14 nm and the lattice-strain is relatively high for TAN at £ 600 0C. It is due to surface defects in the material [4, 13]. Thermal heating of the samples for TAN ³ 800 0C increases grain size (range 23-35 nm) and decreases lattice strain. This is supported by a rapid increase of SXRD peak sharpness. We have used SEM images (Fig. 3(a-d)) to show two distinct regimes of particle morphology. The clusters of nearly spherical shaped particles dominated for the samples at 6
TAN ≤ 800 0C and high crystalline structured polygonal shaped particles are observed for the samples at TAN between 900-1000 0C. Estimated sizes of the clusters are ~ 250 nm, 350 nm, 820 nm, and 980 nm for the samples CF20, CF80, CF90, and CF100, respectively. These values are nearly 15-25 times larger than the grain size (13-35 nm) from SXRD pattern. This means clusters in SEM images consist of many small particles/grains surrounding a core particle. The particles (grains) are melted at higher annealing temperature to form a polygonal shaped particle, as shown for CF100 sample (Fig. 3(d)). The point-shot (Fig. 3 (e-f)) and line scan (Fig. 3(g-h)) analysis in the elemental specific EDX spectra confirmed the presence of Co, Fe and O elements in the samples. The average atomic ratio of Co:Fe (1.25 ±0.05: 1.75; Co content was calculated by normalizing Fe content to 1.75) is found close to the expected atomic ratio in Co1.25Fe1.75O4 ferrite. 3-4 areas were randomly selected for recording the EDX spectra. In each selected area, spectrum was recorded for whole area (100 mm x 100 mm) and also at 7 arbitrary points using point-shot analysis. A minor heterogeneity in the distribution of Co and Fe atoms could be attributed to surface roughness of the samples. We recorded Raman spectra to get an insight of the micro-structure of the samples. Raman spectra (Fig. 4) indicated five phonon modes (A1g+ 1Eg+ 3T2g), representing cubic spinal structure [18, 26]. The peaks indicate asymmetric shape with doublet structure (higher frequency bands) or satellite peaks about the main peak position (lower frequency bands). Such characteristic peaks represented mixed type spinel structure (i.e., Fe3+ ions occupy both A and B sites) in NiFe2O4 [27]. We assigned the peak profiles centred at 190 ± 10, 300 ± 20, 460 ± 20, 560 ± 20, and 600-700 cm-1 under the bands of T2g(1), Eg, T2g(2), T2g(3) and A1g, respectively. The T2g(1) band arises due to translational movement of tetrahedron (AO4). The low frequency bands (Eg and T2g(2)) are associated with symmetric and antisymmetric stretching vibration at B sites. The band profile of T2g(3) is caused by asymmetric bending of oxygen with respect to metal ion. The high frequency band A1g(1) is associated with
7
symmetric stretching (breathing) mode of vibration at A sites [19]. The peak profiles around 620 cm-1 can be assigned to A1g(2). This sub-band of A1g(1) may be directly related to Co-O bonds in A sites. The sub-bands or splitting of the bands represents the mixed ordering of Co and Fe ions in spinel structure of Co1.25Fe1.75O4 ferrite. Such site disorder of cations is expected due to competition for stabilizing this slightly Co rich ferrite either to inverse spinel structure as in CoFe2O4 [1, 26] or to normal spinel structure as in Co3O4 [18,28]. The peak profiles were fitted with Gaussian shape (Fig. 4) and analysis of the peak parameters (width, position, and intensity) guides the micro-structural changes induced by heat treatment. Fig. 5(a-c) shows the annealing temperature effects on the peak parameters of the major intensity component of the bands (A1g(1), A1g(2), T2g(2) and Eg). Fig. 5(a) shows that peak width of the bands is relatively broad for TAN at £ 600. The peak width reduces for TAN at 800-1000 0C. The increasing trend of peak width for the Eg band is most probably affected by the gradual reduction of peak intensity with increasing TAN. Fig. 5(b) showed a minor decreasing trend of peak position of the bands on increasing TAN up to 600 0C. Then, peak position shifted to higher wave numbers on increasing TAN up to 1000 °C. The peak position (n4 =
5
!67
;?$@
:m
?$@
) is correlated to the spring constant kM-O (¥ rM-O-3; length rM-O) and
effective mass (mM-O) of the metal-oxygen (M-O) bonds. It appears that the peak broadening due to confinement of nano-sized grains and non-equilibrium structure [13, 26] affected the samples at lower TAN. At higher TAN (higher grain size), a notable shift of the peak position in Raman active bands could be affected by re-ordering of Co2+ and Fe3+ ions among A and B sites of the spinel structure. The re-ordering of Co2+ and Fe3+ ions can be correlated with the variation of peak intensity (Fig. 5(c)). The peak intensity for the A1g(1) increased throughout the annealing temperature. The peak intensity for A1g(2), T2g(2) and Eg bands have increased up to TAN = 900 0C, followed by a significant decrease for the sample annealed at 1000 0C.
8
According to perturbation theory [29], the intensity of the scattered light in Raman spectrum can be written as below. Im = KI0ln4 4½αml(n4 )½2
(m, l = x, y, z)
(3)
Here, K is constant, αml(n4 ) is dynamic molecular polarizability tensor, I0l is the intensity of incident light with polarization direction l and Im is the intensity of scattered light with polarization direction m. Since spectrum for all samples were recorded in identical conditions, the parameters K, I0l and l are assumed to be constant for all samples. The intensity variation of the samples is determined mainly by peak positionn4 and polarizability αml(n4 ). A non linear variation (not shown in Figure) of the intensity with n4 for major peaks suggests that αml(n4 ) played a significant role on determining the peak intensity of the samples. We understand the site disorder of Co2+ and Fe3+ ions by combining the results of SXRD pattern and Raman spectrum. According to earlier report [30], the (331) peak in XRD pattern (as in Fig. 5(d)) indicates the presence of Co2+ ions in A sites of the cubic spinel structure. The intensity ratio (I331/I400) of (331) and (400) peaks determines the type of spinel structure, i.e., 10% ratio for normal spinel structure (A sites fully occupied by Co2+ ions) and 0.1 % ratio for perfect inverse spinel structure (A sites fully occupied by Fe3+ ions). Fig. 5(e) shows the I331/I400 ratio in the range 0.5-1.5 (%) and suggests mixed type spinel structure in Co1.25Fe1.75O4 ferrite, where both Co2+ and Fe3+ ions occupy A sites [14, 15, 18]. On the other hand, the relative intensity of A1g(1) peak with respect to A1g(2) peak (I A1g(1)/I A1g(2)) in Raman spectrum can provide a qualitative information of the exchange of cations (Fe3+ and Co2+) among A and B sites [26]. As shown in Fig. 5(f), the A1g(1) and A1g(2) intensity ratio has shown a general decreasing trend with the increase of TAN, although the structural heterogeneity for the samples with lower annealing temperatures can influence the accuracy of such peak ratio. A general decreasing trend of the SXRD peak intensity (I331/I400) ratio and Raman peak intensity (I A1g(1)/I A1g(2)) ratio suggest the decreasing population of Co2+ ions at A sites 9
with the increase of annealing temperature, although an intermediate fluctuation is observed. Generally, ferroelectric polarization is not expected in ideal cubic spinel structure of ferrite. Although there may not be a direct correlation between molecular polarizability αml(n4 ), which is induced by radiation (electric field), and ferroelectric polarization, which may be spontaneous or field induced [22], but atomic scale disorder can produce non-negligible electric polarization. Myoung et al. [31] discussed multiferroic properties in Co0.5Fe0.5Cr2O4 ferrite in terms of local non-cubic distortion due to site exchange of cations and spin ordering in lattice structure. Such lattice distortion is not within the detection limit of XRD pattern. However, the differences in ionic radii and mass of the Co and Fe ions give rise a distribution of length (rM-O) and effective mass (mM-O) of the metal-oxygen (M-O) bonds in both A and B sites. Such atomic level disorder was realized from Raman spectra [30] in the form of either peak splitting (higher frequency bands) or satellite peaks (low frequency bands). It is reported that magnetostriction in association with ferrimagnetic spin order can produce ferroelectric polarization in spinel structure [32]. Recently, some reports [16, 20] have shown considerably large magnetostriction and ferroelectric properties in Co ferrites [3, 24-25]. This motivates us for searching and improvement of electric polarization and ferrimagnetic properties at room temperature in the present ferrite samples. 3.2. Electric field dependence of electric polarization Fig. 6 shows the electric field dependence of polarization (P-E) of the samples. The P-E loops were measured using standard bipolar sine voltage at frequency 1 kHz. The samples at TAN ≤ 600 0 showed completely round shaped loop. For these samples, maximum polarization (Pmax) is observed at zero electric field. The remanent polarization (PR), which is equal to Pmax in these samples, does not sustain with the increase of electric field and becomes zero at the highest electric field. Such P-E loop does not carry true electric polarization and it is artefact due to conductive effect [33]. We see the improvement in P-E loop with non-zero remanent
10
polarization (PR) at higher electric field for the samples at TAN ³ 800 ºC. The feature of the PE loop for these samples is different from the character of both ideal resistor (elliptical shape with major axis along field axis) and ideal dielectric (straight line without loop due to capacitance response). The tilted P-E loops for the CF80, CF90 and CF100 samples may represent the response of a non-linear dielectric (assumed to be consisting of capacitive and conductive effects) or under saturated loop of a ferroelectric material [33]. The notable feature is that the capacitor component increases with the increase of TAN of the samples. We attempted to distinguish the purely conductive effect from capacitive effect by analysing the leakage of polarization for electric fields at above and below of the field (Em) corresponding to Pmax. The loss of polarization (LP) due to purely conductive effect at electric fields above Em is defined as LP (%) =
(ABCD AEBCD )∗5FF ABCD
, where PEmax is the polarization at maximum
applied electric field (Emax). The retaining of polarization (RP) due to capacitive effect at zero electric field during decrease of field below Em is defined as RP (%) =
AG ∗5FF ABCD
. We observe
(LP) 100% leakage of polarization due to conductive effect for the samples at low TAN. Although LP ~ 11 % is observed in the sample with TAN at 800 ºC (Fig. 6(c)), the Pmax now differed from PR. This shows that conductive effect may dominate at higher electric fields for CF80 sample, but capacitive effect also improved, which gives rise to RP ~ 77 % and remaining 23% of the Pmax has leaked through intrinsic (grain boundary) disorder in the sample. There is no decrease of polarization at higher electric field (maximum ~ 14 kV/cm, corresponds to applied voltage 500 V) for the samples with TAN at 900-1000 ºC (Fig. 6 (d-e)). A significant decrease of the minor diameter length (shown by red arrow for CF80 sample) indicates the decrease of conductive effect at higher TAN. However, a fraction of polarization leaked through grain boundary and retained ~ 42 % and ~ 46 % of the Pmax at zero field for the samples CF90 and CF100, respectively. It may be mentioned that ferroelectric
11
polarization in BaTiO3 is determined by the switchable polarization PS = 2PR. In semiconductors, leakage component is added in net polarization. The leakage component depends on R-C time constant and electrical conductivity of the samples [33]. Fig. 6(f) compares the P-E loops of the sample CF100, measured at frequencies 1 kHz, 200 Hz and 100 Hz of applied voltage 500 V. The decrease of frequency increases both RP and LP of the sample, e.g., RP increases up to 78 % and 89 %, and LP increases up to 34% and 41 % for frequency at 200 Hz and 100 Hz, respectively. Such increase of the retaining polarization on lowering frequency is controlled by conductivity effects at grain boundaries. We understand such artefact effects on the shape of P-E loop through dielectric properties of the samples. Fig. 7(a-c) shows the frequency (range 1-107 Hz) variation of ac conductivity (σ/), dielectric loss (tand), and dielectric constant (e/). Fig. 7(a) shows that conductivity of the samples increases with measurement frequency. The conductivity of the CF20 and CF50 samples (TAN £ 500 0C) are significantly high throughout the applied frequencies in comparison to the conductivity of the samples annealed at higher temperatures. The frequency activated conductivity in Co ferrite can be explained in terms of hopping of charge carriers via superexchange paths (electron hopping via Fe3+-O2—Fe2+ and hole hopping via Co2+-O2—Co3+) in octahedral lattice sites of the spinel structure [7, 24]. At low frequency regime, the dc limit of conductivity (sdc = nemn + pemh) depends on the density (n: electron density in conduction band, p: hole density in valence band) and mobility of charge carriers (mn: electron mobility, mh: hole mobility). The high conductivity for the samples at low TAN (£ 500 0C) can be attributed to high carrier density and mobility, arising due to surface defects [13]. Fig. 7(b) shows dielectric loss of the samples below 20. The dielectric loss in the high conductive samples is above 1 over a large frequency range. The samples at higher TAN (600-1000 ºC) showed relatively low dielectric loss, typically less than 1 for frequencies above 100 Hz. The dielectric loss exhibited a peak like behaviour in the frequency domain 12
that matches with the average hopping frequency of the charge carriers. The peak frequency can be used to roughly calculate the relaxation time of the R-C circuit for the samples. The decrease of dielectric loss at higher frequencies is expected due to mismatch of the applied frequency and hopping frequency. A rapid increase of the dielectric loss at lower frequencies indicates the space charge contribution to electric polarization at the interfaces of grains and grain boundaries [7, 13]. The dielectric constant (e/ = C/Cair) is another important parameter, which is directly related to the capacitance (C) of the sample. Fig. 7(c) showed a relatively high dielectric constant for the samples at low TAN. Such high value of dielectric constant is contributed by space charge polarization effect. The space charge polarization significantly reduces by increasing the annealing temperature of the samples. The dielectric constant of the CF90 and CF100 samples is nearly stable, indicating extremely low space charge polarization effect and development of polarization in the samples from electric dipoles. Fig. 7 (d-f) has compared the dielectric parameters at 1 kHz (sufficiently away from low frequency space charge regime) for all samples. The most striking feature is that electrical conductivity, dielectric lass and dielectric constant rapidly decreases with the increase of annealing temperature. The conductivity of the material at 1 kHz is extremely small (10-10 S/cm) for the samples with TAN in the range 900 0C-1000 0C, where as dielectric loss and dielectric constant attained the values 0.04-0.1 and 7-8, respectively. The results indicate that the conductive effect rapidly reduces with the increase of annealing temperature. The phase shift (φ) may be a good tool to differentiate capacitive effect from the conductive effect in the samples. This is obtained from measurement of complex impedance components (tanf = J // J/
L
= M , where IC and IR are the current components passing through the capacitance (C) and LG
resistance (R), respectively in RllC circuit model). In the case of purely conductive effect, f is 0 0 (IC = 0). In the case of pure capacitive effect, f is 90 0 (IR = 0). Fig. 7(e)-right Y axis shows that f is small (9-23 0) for the CF20 and CF50 samples. The f has increased to 67 0 for 13
CF60 sample and stabilized in the range 72-88 0 for TAN ³ 800 0C. The high conductivity, high dielectric loss, low phase shift and 100 % leakage of polarization suggest that P-E loops for the samples annealed at £ 600 0C are dominated by conductive effect. On the other hand, capacitive effect has improved electric polarization in the P-E loop of CF90 and CF100 samples, although conductive effect is not completely suppressed in these two samples. We have examined the possibility of further improvement of electric polarization by increasing the annealing time of CF90 sample at 900 0 C from 2 h to 4 h. The P-E loops in Fig. 8 confirm the improvement of stable electric polarization in CF90 sample by increasing its annealing time to 4 h. We obtained a good shaped P-E loop with negligible leakage of polarization (ac voltage up to 1 kV and measurement frequency down to 50 Hz). Although the loops are under saturated, but some of the loops showed a tendency of saturation, irrespective of the magnitude of applied frequency (Fig. 8(a-c)) and voltage (Fig. 8(d-f)). Non-linear dielectrics may exhibit such characteristic features at frequency which matched near to the critical frequency of the R-C equivalent circuit of the sample [33]. In our sample, a good quality P-E loop has been observed for a wide range of frequencies measured below 1 kHz. Second important point is that magnitude of the electric polarization increases with the increase of applied voltage, without loss of polarization at higher electric fields. These typical characters of the P-E loops indicate a possible ferroelectric polarization in the sample, although a fraction of polarization cannot be retained at zero electric field due to intrinsic (grain boundary) disorder and semiconductor nature of the sample. Fig. 9 plots the derived parameters from loops (Pmax, PEmax, PR, electric coercivity (EC), retaining of polarization (RP) at zero field and loss of polarization (LP) at maximum field with respect to the Pmax in positive electric field side). The general feature is that Pmax, PR, EC and RP(%) rapidly decreased at low frequencies (up to 67 Hz) and more or less stabilized at higher frequency. The Pmax is found in the range 4-22 nC/cm2 (Fig. 9(a)). The PR is found in the range 2-11
14
nC/cm2 (Fig. 9(b)). The EC is found in the range 2-15 kV/cm (Fig. 9(c)). The retaining of polarization (RP) at zero field is found in the range 30-60 % (Fig. 9(d)), where as the leakage of polarization at highest field is found in the range 0-8% mostly below 67 Hz (Fig. 9(e)). Since the loops are under saturated, the observed range of ferroelectric parameters (Fig. 9) may be less than actual values for the sample. The ferroelectric polarization in the present material is low due to many reasons. The primary reason is that it is not a traditional ferroelectric, like BaTiO3, where displacement of d0 shell ions (Ti4+) in non-cubic (tetragonal) structure produces large electric polarization [22]. The ferrite is having partially filled 3d shell ions (Fe and Co) that act as the hopping centres for charge transfer and it increases dielectric loss and electrical conductivity. The present material belongs to the class of nontraditional ferroelectric material [3, 21, 30, 32-34]. In such ferroelectrics, a non-negligible displacement of B sites cations and subsequently, charge imbalance between lattice sites can be anticipated during hopping and site exchange of cations. Even, magnetic spin disorder and spin-lattice coupling can produce small ferroelectric polarization [33, 34]. The P-E loops in Co1.25Fe1.75O4 samples are better than that reported for CoFe1.8Ti0.2O4 ferrite [3]. The ferroelectric parameters in our material are comparable to that found in the composite of Ba0.6Sr0.4TiO3 and CoFe2O4 [35] or similar composites [36-39]. The ferroelectric parameters in our sample at room temperature are magnitude wise better than that in a good ferroelectric material MnWO4 with PR (0-40 mC/m2 = 0-4 nC/cm2) and EC (0-300 kV/m = 0-3 kV/cm) at extremely low temperature (< 12 K) [40] and such system may be interesting for basic theory, but does not have much importance for room temperature applications. 3.3. Magnetic properties The magnetic field dependence of magnetization (M(H)) curves in Fig. 10 confirms the ferrimagnetic nature with the exhibition of hysteresis loop in all samples. We have calculated the ferrimagnetic parameters (Coercivity (Hc), spontaneous magnetization (Ms), and
15
squareness (SQ: ratio of remanent magnetization (MR) with Ms) using hysteresis loop. The linear extrapolation of the M2 vs. H/M curve from higher field side on M2 axis was used to determine MS of the samples. The Coercivity (Hc) and remanent magnetization (MR) of the samples were calculated by magnifying the loop in low field scale (inset of Fig. 10(a)). As shown in Fig. 10(b), an initial decreasing trend of MS with increase of annealing temperature could be related to non-equilibrium spinel structure in the samples [13]. The increase of MS value with annealing temperature at ³ 600 0C could be related to the refinement of grain boundary structure. The re-ordering of Co2+ and Fe3+ ions among A and B sites of the spinel structure also affects on the improvement of magnetization with annealing temperature. The observed maximum MS (~ 69 emu/g) in CF100 sample is considerably less than the value of MS ~ 80 emu/g in CoFe2O4 [2], but higher than the MS ~ 55 emu/g in Co1.5Fe1.5O4 [4]. This is due to increasing population of non-magnetic Co3+ ions at B sites as the Co content increases in CoxFe3-xO4 ferrite for x > 1 [14]. Fig. 10(b) shows a monotonic increase of the squareness (SQ) to achieve the maximum value (~ 0.5) for TAN = 900 0C, then it significantly in CF100 sample. The observed SQ values are comparable to ~ 0.47-0.54 reported for CoFe2O4 [2] and ~ 0.33 found in Co1.5Fe1.5O4 [4]. Fig.10(c) shows that magnetic coercivity of the material increased from 159 Oe in CF20 to the maximum value ~ 580 Oe in CF80 sample and then, decreases to 160 Oe in CF100 sample. It can be attributed to transformation of magnetic domain structure from single domain state for TAN below 800 °C (grain size 10-23 nm) to multi domain state for TAN in the range 900-1000 °C (grain size 28-35 nm) by comparing similar variation of squareness and coercivity with annealing temperature in Co ferrite samples [2, 4, 12]. The maximum magnetic coercivity in the present Co1.25Fe1.75O4 ferrite is smaller than the coercivity ~1.2-1.4 kOe for CoFe2O4 at TAN = 800 ºC [12]. It is due to presence of Co3+ ions that reduce magneto-crystalline anisotropy in the spinel structure [16]. . 16
4. Conclusions Heat treatment of the as-prepared sample has brought several changes in the electric and ferrimagnetic properties of Co1.25Fe1.75O4 ferrite. The samples annealed at low temperatures exhibited non-equilibrium spinel structure. The samples annealed at higher temperatures indicated well crystalline spinel structure. The improved micro-structure (large grain size, grain boundary refinement and re-ordering of Co2+ and Fe3+ ions among A and B sites of the spinel structure) by heat treatment played an important role for development of electric polarization and ferrimagnetic parameters in the samples. The present ferrite system belongs to the class of non-traditional ferroelectrics, where hopping of charge carriers between lattice sites and semiconductor nature of the samples introduces electrical conductive paths for leakage of polarization. In this paper, we have shown the decrease of polarization leakage due to conductive effect and improvement of electric polarization due to capacitive effect by increasing the annealing temperature and time. We believe the range of temperature where the material exhibits ferroelectric polarization is more important than the magnitude of electric polarization. The leakage of polarization in the present ferrite samples can be further improved by adopting 3 different techniques, (1) high annealing temperature and time, (2) dispersion of ferrite particles in insulating (ferroelectric) matrix like BaTiO3, (3) formation of laminated composites of ferrite and BaTiO3. We are working on these three aspects for developing the ferrite based multiferroic and multifunctional materials, where ferrite will used as the source of magnetic and ferroelectric component and BaTiO3 will be used for resisting leakage of polarization through grain boundary of ferrite particles and additional source of ferroelectric polarization. Further, the establishment of a coupling between ferrimagnetic order and electric polarization by controlling interfacial lattice structure and interactions in such hybrid multi-ferroics could be useful for the application in magnetic and electric field controlled devices.
17
Acknowledgements We acknowledge the support of Central Instrumentation Facility (CIF) of Pondicherry University for material characterization. The authors thank Ms. V. Vasanthi and Mr. M.R. Panda for technical assistance. RNB thanks DAE-BRNS, Govt. Of India for supporting research Grant (NO. 2011/37P/45/BRNS/2628). RNB thanks UGC, Govt. Of India for supporting research Grant (F.No. 42-804/2013 (SR). References [1] J.A. Moyer, C.A.F. Vaz, E. Negusse, D.A. Arena, and V. E. Henrich, Phys. Rev. B 83, 035121(2011) . [2] A. Franco Jr, and V. Zapf, J. Magn. Magn. Mater. 320, 709 (2008). [3] G. D. Dwivedi, A. G. Joshi, H. Kevin , P. Shahi , A. Kumar, A. K. Ghosh, H. D. Yang, and S. Chatterjee, Solid State Comm. 152, 360 (2012). [4] R. N. Bhowmik, A.T. Sathya, and A. Bharathi, J. Alloys Compd. 559, 134 (2013). [5] W. S. Chiu, S. Radiman, R. A. Shukor, M. H. Abdullah, and P. S. Khiew, J. Alloys Compd. 459, 291 (2008). [6] R.N. Bhowmik, V. Vasanthi, and A. Poddar, J. Alloys Compd. 578, 585 (2013). [7] M. George, S.S. Nair, K.A. Malini, P.A. Joy, and M.R. Anantharaman, J. Phys. D: Appl. Phys. 40, 1593 (2007). [8] V.G. Harris, A. Geiler, Y. Chen, S.D. Yoon, M. Wu, A. Yang, Z. Chen, P. He, P.V. Parimi, X. Zuo, C.E. Patton, M. Abe, O. Acher, and C. Vittoria, J. Magn. Magn. Mater. 321, 2035 (2009). [9] G.S.N. Rao, O.F. Caltun, K.H. Rao, P.S.V. Subba Rao, and B.P. Rao, J. Magn. Magn. Mater. 341, 60 (2013). [10] K. Inomata, N. Ikeda, N. Tezuka, R. Goto, S. Sugimoto, M. Wojcik, and E. Jedryka, Sci. Technol. Adv. Mater. 9, 014101 (2008).
18
[11] K.F. Wang, J.-M. Liu, and Z.F. Ren, Advances in Physics 58, 321 (2009). [12] A. Lunhong, and J. Jiang, Current Appl. Phys. 10, 284 (2010). [13] R.N. Bhowmik, Mater. Res. Bull. 50, 476 (2014). [14] G.H. Jonker, J. Phys. Chem. Solids 9 (1959) 165. [15] H. L. Trong, L. Presmanes, E. D. Grave, A. Barnabe, C. Bonningue, and Ph. Tailhades, J. Magn. Magn. Mater. 334, 66 (2013). [16] I. C. Nlebedim, A. J. Moses, and D. C. Jiles, J. Magn. Magn. Mater. 343, 49 (2013). [17] I. P. Muthuselvam, and R.N. Bhowmik, Solid State Sci. 11, 719 (2009). [18] N. Bahlawane, P.V.T. Ngamou, V. Vannier, T. Kottke, J. Heberle, and K.K. -Hoinghaus, Phys. Chem. Chem. Phys. 11, 9224 (2009). [19] M.R. Panda, R.N Bhowmik, H. Singh, M.N. Singh, and A K Sinha, Mater. Res. Exp. 2, 036101 (2015). [20] I. C. Nlebedim, N. Ranvah, P. I. Williams, Y. Melikhov, J. E. Snyder, A. J. Moses, and D. C. Jiles, J. Magn.Magn. Mater. 322, 1929 (2010). [21] A. Das, G.D. Dwivedi, P. Kumari, P. Shahi, H. D. Yang, A. K. Ghosh, and S. Chatterjee, J. Magn. Magn. Mater. 379, 6 (2015). [22] D. Khomskii, Physics 2, 20 (2009). [23] R.N. Bhowmik, M.R. Panda, S.M. Yusuf, M.D. Mukadam, A.K. Sinha, J. Alloys Compd. 646, 161 (2015). [24] R.N. Bhowmik, and I. Panneer Muthuselvam, J. Alloys Compd. 589, 247 (2014). [25] P.L. Meena, R. Kumar, C.L. Prajapat, K. Sreenivas, and V. Gupta, J. Appl. Phys. 106, 024105 (2009). [26] P. Chandramohan, M.P. Srinivasan, S. Velmurugan, and S.V. Narasimhan, J. Solid State Chem. 184, 89 (2011).
19
[27] Z. Ž. Lazarević, Č. Jovalekić, D. Sekulić, M. Slankamenac, M. Romčević, A. Milutinović, and N. Ž. Romčević, Science of Sintering 44, 331 (2012). [28] S. Grace Victoria, A. M. E. Raj, and C. Ravidhas, Mater. Chem. Phys. 162, 852 (2015). [29] A.C. Albrecht, The J. Chem. Phys. 34, 1476 (1961). [30] J. R.-Fuertes, T. Bernert, M. He, B. Winkler, V.L. Vinograd, and V. Milman, Appl. Phys. Lett. 105, 071911 (2014). [31] B.R. Myoung, and C.S. Kim, J. Appl. Phys. 117, 17B741 (2015). [32] S.V. Streltsov, A.I. Poteryaev, and A.N. Rubtsov, J. Phys.: Condens. Mater 27, 165601 (2015). [33] K.M. Rabe, M. Dawber, C. Lichtensteiger, C.H. Ahn, and J.-M. Triscone (Eds.): Physics of Ferroelectrics: A Modern Perspective, Topics in Applied Physics 105, 1-30 (2007) [34] A.G. Lone, and R.N. Bhowmik, J. Magn. Magn. Mater. 379, 244 (2015). [35] W.Y. An, W. Y. Bo, R. Wei, G. J. Xiong, and Z. W. Li, Chin. Phys. Lett. 29, 067701 (2012). [36] L. B. Hao, D. X. Zhou, S. P. Gong, Q. Y. Fu, W. Luo , G. Jian, F. Xue, and L. Zhou, J Mater Sci: Mater Electron 24, 2351(2013). [37] S. T. Zhang , Y. Zhang, Z. L. Luo, M. H. Lu, Z. B. Gu, and Y. F. Chen, Applied Surface Science 255, 5092 (2009). [38] J. Nie, G. Xu, Y. Yang, and C. Cheng, Mater. Chem. Phys. 115, 400 (2009). [39] L.Y. Ding, F. X. Wu, Y. B. Chen, Z. B. Gu, and S.T. Zhang, Appl. Surf. Sci. 257, 3840 (2011). [40] B. Kundys, C. Simon, and C. Martin, Phys. Rev. B 77, 172402 (2008).
20
Figure captions: Fig. 1 (Colour online) SXRD pattern and profile fitted data of the annealed samples. Fig. 2 (Colour online) (a) Fitting of the (440) peak using Gaussian curves. (b) fitting of the peak parameters according to Williamson-Hall equation. (c) variation of grain size and lattice strain (%) with annealing temperature of the sample. Fig.3 (Colour online) SEM images (a-d), EDX data- point analysis (e-f) and EDX data- line scan analysis (g-h) shown for selected samples. Fig. 4 (Colour online) Raman spectra of the samples, along with fitted data (lines). Fig. 5 (Colour online) (a-c) peak parameters from Raman spectra. (d) SXRD (331) peak, (e) (331) and (400) SXRD peak ratio, (f) A1g(1) and A1g(2) peak ratio (Raman spectra). Fig. 6 (Colour online) (a-e) P-E loop measured for different samples at 1 kHz loop frequency. (e) P-E loop measured for CF100 sample at different loop frequencies. Fig. 7 (Colour online) (a-c) frequency dependence of dielectric parameters for the samples annealed at different temperatures. (d-f) dielectric parameters of the samples at 1 kHz and 1 Hz with the variation of annealing temperature. Fig. 8 (Colour online) P-E loops of the CF90 sample annealed for 4 hrs in different modes by varying the measurement loop frequency and applied voltage. Fig. 9 (Colour online) Variation of the parameters with measurement loop frequencies and the parameters have been derived from the P-E loops in Fig. 8. Fig. 10 (Colour online) (a) Field dependence of magnetization curves. Annealing temperature dependence of magnetic parameters (b-c).
21
Highlights Ø Ø Ø Ø Ø
Nanoparticles of Co1.25Fe1.75O4 ferrite have been prepared in alkaline medium. Nanoparticles have been annealed to tailor magnetic and ferroelectric properties. Room temperature ferrimagnetism is confirmed in all samples. Some of the samples showed good signature of ferroelectric properties. Role of conductive effect and capacitive effect understood.
22
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PII: DOI: Reference:
S0304-8853(16)31674-2 http://dx.doi.org/10.1016/j.jmmm.2016.08.014 MAGMA61699
To appear in: Journal of Magnetism and Magnetic Materials Received date: 14 April 2016 Revised date: 22 June 2016 Accepted date: 3 August 2016 Cite this article as: R.N. Bhowmik and A.K. Sinha, Improvement of room temperature electric polarization and ferrimagnetic properties of Co1.25Fe1.75O ferrite by heat treatment, Journal of Magnetism and Magnetic Materials, http://dx.doi.org/10.1016/j.jmmm.2016.08.014 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Improvement of room temperature electric polarization and ferrimagnetic properties of Co1.25Fe1.75O4 ferrite by heat treatment R.N. Bhowmika*, and A.K. Sinhab a
Department of Physics, Pondicherry University, R.V. Nagar, Kalapet, Pondicherry -605014,
India b
Indus Synchrotrons Utilization Division, Raja Ramanna Centre for Advanced Technology,
Indore- 452013, India *
Corresponding author (RNB): Tel.: +91-9944064547; Fax: +91-413-2656988, E-mail: [email protected]
Abstract We report the improvement of electric and magnetic properties of the chemical routed samples of Co1.25Fe1.75O4 ferrite. Synchrotron X-ray diffraction pattern confirmed the cubic spinel structure in all samples. Raman spectroscopy provided the information of mixed-type cubic spinel structure, where Co2+ and Fe3+ ions order in octahedral and tetrahedral sites. The thermal heat treatment has brought a systematic change in electric polarization and magnetic parameters. At a critical annealing temperature (TAN) of 800 °C, the sample has shown the highest magnetic coercivity ~ 596 Oe and relatively large squareness ~ 0.5. The electrical conductivity, dielectric loss and dielectric constant have decreased with increase of TAN. The dielectric constant becomes stable for TAN above 800 °C. The samples annealed in the temperature range 900 0C -1000 0C have shown a good signature of ferroelectric polarization, although it is under saturated and leakage of polarization due to conductive effect is observed. This work shows the scope of the improvement of electric polarization due to capacitive effect on increasing the annealing temperature of as-prepared Co ferrite sample. Keywords: Magnetic materials; Chemical synthesis; X-ray diffraction; Raman spectroscopy; Ferroelectricity. 1
1. Introduction Cobalt substituted magnetite (CoxFe3-xO4) belongs to the class of cubic spinel structured ferrite with formula unit AB2O4, where A and B represent the tetrahedral and octahedral coordinated interstitial lattice sites for metal ions in oxygen environment [1]. The variation of Co content, and distribution and charge states of Co and Fe ions have played an important role for tuning the structure, ferrimagnetic, dielectric, and ferroelectric properties in Cobalt substituted ferrite [1-7]. The coexistence of ferrimagnetic and ferroelectric properties in ferrite structure could enhance its applications in many modern areas of microelectronics, data storage devices, microwave devices, and sensors [8-10]. The establishment of a coupling between ferroelectric and ferrimagnetic properties in Co ferrite may be an additional advantage for exploring the multi-functional applications of ferrites in non-volatile memory devices, magnetic and electric field controlled sensors, and spintronics [11]. We have noted that chemical routed samples of CoFe2O4 and Co1.5Fe1.5O4 ferrites formed single phased cubic spinel structure for a wide range of annealing temperature [4, 1213]. The Co and Fe ions in such single phased spinel ferrites can be distributed as (Co2+1-b Fe3+b)A[Co2+bFe3+3-x-bCo3+x-1]O4 ; 0 ≤ b ≤ 1 [14-15]. The distribution of magnetic ions (Co2+ highly anisotropic and Fe3+) in A and B sites of the spinel structure quantifies the magnetic moment (M = MB – MA, where MB and MA are the net moments in B and A sites) per formula unit of Co ferrite. Substitution of high spin Fe3+ moments at B sites by the low spin (nonmagnetic) Co3+ decreases the magnetization and magneto-crystalline anisotropy in CoxFe3xO4 ferrite
[2, 15]. The increase of Co atoms in CoxFe3-xO4 ferrite for x > 1.5 showed the
structural phase destabilization, randomization and re-configuration of the charge states of Co and Fe ions [2, 15-20]. The Co content in CoxFe3-xO4 ferrite also determines the nature of semiconductor properties in ferrite structure whether p type for x > 1 (hole hopping through Co2+-O2—Co3+ superexchange paths) or n type for x £ 1 (electron hopping through Fe3+-O2-
2
Fe2+ superexchange paths) [1, 14, 18]. Recently, some reports showed ferroelectric properties in CoFe2O4 ferrite [3, 21]. CoFe2O4 ferrite may exhibit reasonably large magnetization and coercivity, but it is a poor ferroelectric material with huge leakage of polarization. The ferroelectric polarization in Co ferrite is obviously small due to cubic structure with semiconductor nature and presence of magnetic ions. Hence, improvement of ferroelectric polarization with low leakage in ferrite is a challenging task. Such ferroelectric materials, different from traditional ferroelectric BaTiO3, may be classified as type –II multiferroics, where spontaneous electric polarization is extremely low, but there is a possibility of finite coupling between electric polarization and ferro/ferrimagnetic order [22]. Some of the recent works on Co rich ferrite [4, 6, 13, 17, 23-24] showed a lot of scope for achieving moderate magnetization, high squareness and coercivity with moderate dielectric constant and low dielectric loss. A detailed study of the dielectric, ferroelectric and ferrimagnetic properties in Co rich side (x > 1) of CoxFe3-xO4 ferrite could be useful for basic understanding and development ferrite based multiferroic materials. Taking a clue from the observation of ferroelectric and ferrimagnetic properties in Co rich spinel ferrites [24-25], this work is dedicated for searching the signature of ferroelectric properties in Co1.25Fe1.75O4 ferrite. Subsequently, efforts are made to improve the ferroelectric polarization and ferrimagnetic properties by adopting thermal treatment of the chemically routed as-prepared material. This ferrite composition belongs to slightly Co rich side (x > 1) of CoxFe3-xO4 ferrite and semiconductor in nature [14]. This ferrite composition is stabilized in single phase for annealing temperature up to 1000 0 C and excludes the effects of structural complexity as found for Co content more than 1.5 [15, 17, 19]. 2. Experimental Chemical co-precipitation route was used to prepare the Co1.25Fe1.75O4 ferrite. The stoichiometric amounts of high purity (> 99.999%) Co(NO3)2.6H2O and Fe(NO3)3.9H2O were
3
dissolved in distilled water to prepare a nitrate solution (pH value ~ 1.6). NaOH solution with pH ~ 13.5 was used as a precipitating agent and added drop wise to nitrate solution until pH reached to 11. The mixed solution was kept at 80 °C for 2 h with continuous stirring and pH ~ 11 is maintained during chemical reaction at 80°C by adding required amount of NaOH solution. Then, the resultant solution was cooled to room temperature and observed a black coloured precipitation. The transparent solution above the precipitate was carefully removed. The by-product mixed with the precipitate was removed by repeated washing processes. During washing process, sufficient amount of distilled water was added with stirring condition and the mixture was heated at 100 °C for 10-15 min. This was followed by cooling to room temperature and allowed to precipitate. The water solution above the precipitate was carefully removed. This washing process was repeated several times to ensure the removal of the by-products, mainly NaNO3 [13]. The absence of any white coloured powder due to formation of NaNO3 or non-absorbance of moisture from air indicated the black coloured ferrite nanoparticles free from by-products. The resultant black powder of ferrite was heated at 200 °C for 2 h and ground well to get black coloured ferrite powder. The ferrite (CF) powder was used to make several pellets of diameter 13 mm. The pellet from of the asprepared ferrite powder after 2 h air annealing (heating and cooling @ 5 °C/min) is denoted as CF20, CF50, CF60, CF80, CF90, and CF100 for annealing temperature (TAN) at 200 °C, 500 °C, 600 °C, 800 °C, 900 °C, and 1000 °C, respectively. The characterization and measurements of the samples were carried out at room temperature. The synchrotron X-ray powder diffraction (SXRD) measurement was performed in the 2θ range 5-40 ° using angle dispersive X-ray diffraction beam line (BL-12) at a wave length 0.7763 Å at Indus-2 synchrotron source, India [19]. The wavelength was calibrated using SXRD pattern of LaB6 NIST standard. We analysed the SXRD pattern of the samples using FULLPROF profile fit program. The Raman spectra of the samples were recorded
4
using Renishaw’s Invia Raman microscope (Renishaw, UK) with 488 nm laser light. Laser power was optimized to 3 mW on the sample surface for acquisition time 30 s. Surface morphology of the samples was checked using scanning electron microscope (SEM) (HITACHI S-3400N, Japan). Energy dispersive x-ray (EDX) analyser was used to estimate elemental composition. The dielectric properties of the samples were measured using a Broadband Dielectric Spectrometer (Novocontrol tech., Germany) in the frequency range 1 Hz to 7 MHz at an ac electrical signal of 1 V. For dielectric measurement, the disc shaped samples (Æ ~ 13 mm, t ~ 1 mm) were used and sand witched between two gold coated electrodes. Electric field dependent polarization (P-E) loop of the thin disc shaped samples (Æ ~ 13 mm, t~ 0.36 -0.42 mm) were measured using Precision Premire II loop tracer (Radiant Tech., USA). Magnetic hysteresis loop of the samples was recorded within field ±70 kOe using Physical properties measurement system (EC2-Quantum Design, USA). 3. Results and discussion 3.1. Structural properties Fig. 1 shows the SXRD patterns, along with profile fit data, of the samples. SXRD patterns confirmed the cubic spinel structure with space group Fd3m in all samples. A careful analysis of the SXRD peaks indicated asymmetric shape for the samples annealed at lower temperatures (200-500 0C), which could be fitted with secondary spinel phase. We have plotted the (440) SXRD peak alone in Fig. 2(a) to show the refinement of crystalline structure with the increase of TAN, because the diffraction peaks at higher 2θ are better choice to observe the second phase, if anything exists. The peaks achieved symmetric shape for TAN in the range 600-1000 0C. Since the SXRD pattern does not show the presence of any impurity phase like CoO or Fe doped CoO, as reported earlier [2, 16], the observed minor peak asymmetry for the samples (CF20 and CF50) prepared at low annealing temperature could be assigned to non-equilibrium spinel phase [13]. The lattice parameter shows a non-monotonic
5
increase with TAN. The lattice parameter for the samples with TAN up to 600 0C may be affected by the structural heterogeneity and non-equilibrium cubic spinel structure. The lattice parameter increases at higher TAN due to improvement of micro-structure (refinement of grain boundary, local re-ordering of Co and Fe ions among A and B sites) [4, 18]. The refinement of crystalline structure with increasing TAN is understood from the evolution of peak profile. Fig. 2(a) shows a broad peak profile for samples at lower TAN (200600 0C) in comparison with sharp peak profile for samples at higher TAN. The (440) peak profile of the samples are analysed by Gaussian curve equation. I = I + I ∗ e
(!)∗"
#$#% * ' &
(1)
In equation (1), xC (= 2θC) is the peak position, Ih is the peak intensity above the back ground intensity Io and β is the full width at half maximum (FWHM) of the peak intensity. These parameters from major peaks were used to calculate the grain size (
0.89*λ
+2εrms sinθ-
(2)
Here, λ is the wavelength of X-ray radiation (0.7763 Å). We have used b and θC values from profile fitting of seven high intensity peaks (111, 220, 311, 400, 422, 511, and 440) to plot the bcosθ- vs. sinθ- data (Fig. 2(b)). The linear intercept on the bcosθ- axis and its slope provided
0.89*λ
and 2εrms , respectively. Fig. 2(c) shows the variation of grain size (
lattice strain (erms) with TAN. The grain size is in the range 12-14 nm and the lattice-strain is relatively high for TAN at £ 600 0C. It is due to surface defects in the material [4, 13]. Thermal heating of the samples for TAN ³ 800 0C increases grain size (range 23-35 nm) and decreases lattice strain. This is supported by a rapid increase of SXRD peak sharpness. We have used SEM images (Fig. 3(a-d)) to show two distinct regimes of particle morphology. The clusters of nearly spherical shaped particles dominated for the samples at 6
TAN ≤ 800 0C and high crystalline structured polygonal shaped particles are observed for the samples at TAN between 900-1000 0C. Estimated sizes of the clusters are ~ 250 nm, 350 nm, 820 nm, and 980 nm for the samples CF20, CF80, CF90, and CF100, respectively. These values are nearly 15-25 times larger than the grain size (13-35 nm) from SXRD pattern. This means clusters in SEM images consist of many small particles/grains surrounding a core particle. The particles (grains) are melted at higher annealing temperature to form a polygonal shaped particle, as shown for CF100 sample (Fig. 3(d)). The point-shot (Fig. 3 (e-f)) and line scan (Fig. 3(g-h)) analysis in the elemental specific EDX spectra confirmed the presence of Co, Fe and O elements in the samples. The average atomic ratio of Co:Fe (1.25 ±0.05: 1.75; Co content was calculated by normalizing Fe content to 1.75) is found close to the expected atomic ratio in Co1.25Fe1.75O4 ferrite. 3-4 areas were randomly selected for recording the EDX spectra. In each selected area, spectrum was recorded for whole area (100 mm x 100 mm) and also at 7 arbitrary points using point-shot analysis. A minor heterogeneity in the distribution of Co and Fe atoms could be attributed to surface roughness of the samples. We recorded Raman spectra to get an insight of the micro-structure of the samples. Raman spectra (Fig. 4) indicated five phonon modes (A1g+ 1Eg+ 3T2g), representing cubic spinal structure [18, 26]. The peaks indicate asymmetric shape with doublet structure (higher frequency bands) or satellite peaks about the main peak position (lower frequency bands). Such characteristic peaks represented mixed type spinel structure (i.e., Fe3+ ions occupy both A and B sites) in NiFe2O4 [27]. We assigned the peak profiles centred at 190 ± 10, 300 ± 20, 460 ± 20, 560 ± 20, and 600-700 cm-1 under the bands of T2g(1), Eg, T2g(2), T2g(3) and A1g, respectively. The T2g(1) band arises due to translational movement of tetrahedron (AO4). The low frequency bands (Eg and T2g(2)) are associated with symmetric and antisymmetric stretching vibration at B sites. The band profile of T2g(3) is caused by asymmetric bending of oxygen with respect to metal ion. The high frequency band A1g(1) is associated with
7
symmetric stretching (breathing) mode of vibration at A sites [19]. The peak profiles around 620 cm-1 can be assigned to A1g(2). This sub-band of A1g(1) may be directly related to Co-O bonds in A sites. The sub-bands or splitting of the bands represents the mixed ordering of Co and Fe ions in spinel structure of Co1.25Fe1.75O4 ferrite. Such site disorder of cations is expected due to competition for stabilizing this slightly Co rich ferrite either to inverse spinel structure as in CoFe2O4 [1, 26] or to normal spinel structure as in Co3O4 [18,28]. The peak profiles were fitted with Gaussian shape (Fig. 4) and analysis of the peak parameters (width, position, and intensity) guides the micro-structural changes induced by heat treatment. Fig. 5(a-c) shows the annealing temperature effects on the peak parameters of the major intensity component of the bands (A1g(1), A1g(2), T2g(2) and Eg). Fig. 5(a) shows that peak width of the bands is relatively broad for TAN at £ 600. The peak width reduces for TAN at 800-1000 0C. The increasing trend of peak width for the Eg band is most probably affected by the gradual reduction of peak intensity with increasing TAN. Fig. 5(b) showed a minor decreasing trend of peak position of the bands on increasing TAN up to 600 0C. Then, peak position shifted to higher wave numbers on increasing TAN up to 1000 °C. The peak position (n4 =
5
!67
;?$@
:m
?$@
) is correlated to the spring constant kM-O (¥ rM-O-3; length rM-O) and
effective mass (mM-O) of the metal-oxygen (M-O) bonds. It appears that the peak broadening due to confinement of nano-sized grains and non-equilibrium structure [13, 26] affected the samples at lower TAN. At higher TAN (higher grain size), a notable shift of the peak position in Raman active bands could be affected by re-ordering of Co2+ and Fe3+ ions among A and B sites of the spinel structure. The re-ordering of Co2+ and Fe3+ ions can be correlated with the variation of peak intensity (Fig. 5(c)). The peak intensity for the A1g(1) increased throughout the annealing temperature. The peak intensity for A1g(2), T2g(2) and Eg bands have increased up to TAN = 900 0C, followed by a significant decrease for the sample annealed at 1000 0C.
8
According to perturbation theory [29], the intensity of the scattered light in Raman spectrum can be written as below. Im = KI0ln4 4½αml(n4 )½2
(m, l = x, y, z)
(3)
Here, K is constant, αml(n4 ) is dynamic molecular polarizability tensor, I0l is the intensity of incident light with polarization direction l and Im is the intensity of scattered light with polarization direction m. Since spectrum for all samples were recorded in identical conditions, the parameters K, I0l and l are assumed to be constant for all samples. The intensity variation of the samples is determined mainly by peak positionn4 and polarizability αml(n4 ). A non linear variation (not shown in Figure) of the intensity with n4 for major peaks suggests that αml(n4 ) played a significant role on determining the peak intensity of the samples. We understand the site disorder of Co2+ and Fe3+ ions by combining the results of SXRD pattern and Raman spectrum. According to earlier report [30], the (331) peak in XRD pattern (as in Fig. 5(d)) indicates the presence of Co2+ ions in A sites of the cubic spinel structure. The intensity ratio (I331/I400) of (331) and (400) peaks determines the type of spinel structure, i.e., 10% ratio for normal spinel structure (A sites fully occupied by Co2+ ions) and 0.1 % ratio for perfect inverse spinel structure (A sites fully occupied by Fe3+ ions). Fig. 5(e) shows the I331/I400 ratio in the range 0.5-1.5 (%) and suggests mixed type spinel structure in Co1.25Fe1.75O4 ferrite, where both Co2+ and Fe3+ ions occupy A sites [14, 15, 18]. On the other hand, the relative intensity of A1g(1) peak with respect to A1g(2) peak (I A1g(1)/I A1g(2)) in Raman spectrum can provide a qualitative information of the exchange of cations (Fe3+ and Co2+) among A and B sites [26]. As shown in Fig. 5(f), the A1g(1) and A1g(2) intensity ratio has shown a general decreasing trend with the increase of TAN, although the structural heterogeneity for the samples with lower annealing temperatures can influence the accuracy of such peak ratio. A general decreasing trend of the SXRD peak intensity (I331/I400) ratio and Raman peak intensity (I A1g(1)/I A1g(2)) ratio suggest the decreasing population of Co2+ ions at A sites 9
with the increase of annealing temperature, although an intermediate fluctuation is observed. Generally, ferroelectric polarization is not expected in ideal cubic spinel structure of ferrite. Although there may not be a direct correlation between molecular polarizability αml(n4 ), which is induced by radiation (electric field), and ferroelectric polarization, which may be spontaneous or field induced [22], but atomic scale disorder can produce non-negligible electric polarization. Myoung et al. [31] discussed multiferroic properties in Co0.5Fe0.5Cr2O4 ferrite in terms of local non-cubic distortion due to site exchange of cations and spin ordering in lattice structure. Such lattice distortion is not within the detection limit of XRD pattern. However, the differences in ionic radii and mass of the Co and Fe ions give rise a distribution of length (rM-O) and effective mass (mM-O) of the metal-oxygen (M-O) bonds in both A and B sites. Such atomic level disorder was realized from Raman spectra [30] in the form of either peak splitting (higher frequency bands) or satellite peaks (low frequency bands). It is reported that magnetostriction in association with ferrimagnetic spin order can produce ferroelectric polarization in spinel structure [32]. Recently, some reports [16, 20] have shown considerably large magnetostriction and ferroelectric properties in Co ferrites [3, 24-25]. This motivates us for searching and improvement of electric polarization and ferrimagnetic properties at room temperature in the present ferrite samples. 3.2. Electric field dependence of electric polarization Fig. 6 shows the electric field dependence of polarization (P-E) of the samples. The P-E loops were measured using standard bipolar sine voltage at frequency 1 kHz. The samples at TAN ≤ 600 0 showed completely round shaped loop. For these samples, maximum polarization (Pmax) is observed at zero electric field. The remanent polarization (PR), which is equal to Pmax in these samples, does not sustain with the increase of electric field and becomes zero at the highest electric field. Such P-E loop does not carry true electric polarization and it is artefact due to conductive effect [33]. We see the improvement in P-E loop with non-zero remanent
10
polarization (PR) at higher electric field for the samples at TAN ³ 800 ºC. The feature of the PE loop for these samples is different from the character of both ideal resistor (elliptical shape with major axis along field axis) and ideal dielectric (straight line without loop due to capacitance response). The tilted P-E loops for the CF80, CF90 and CF100 samples may represent the response of a non-linear dielectric (assumed to be consisting of capacitive and conductive effects) or under saturated loop of a ferroelectric material [33]. The notable feature is that the capacitor component increases with the increase of TAN of the samples. We attempted to distinguish the purely conductive effect from capacitive effect by analysing the leakage of polarization for electric fields at above and below of the field (Em) corresponding to Pmax. The loss of polarization (LP) due to purely conductive effect at electric fields above Em is defined as LP (%) =
(ABCD AEBCD )∗5FF ABCD
, where PEmax is the polarization at maximum
applied electric field (Emax). The retaining of polarization (RP) due to capacitive effect at zero electric field during decrease of field below Em is defined as RP (%) =
AG ∗5FF ABCD
. We observe
(LP) 100% leakage of polarization due to conductive effect for the samples at low TAN. Although LP ~ 11 % is observed in the sample with TAN at 800 ºC (Fig. 6(c)), the Pmax now differed from PR. This shows that conductive effect may dominate at higher electric fields for CF80 sample, but capacitive effect also improved, which gives rise to RP ~ 77 % and remaining 23% of the Pmax has leaked through intrinsic (grain boundary) disorder in the sample. There is no decrease of polarization at higher electric field (maximum ~ 14 kV/cm, corresponds to applied voltage 500 V) for the samples with TAN at 900-1000 ºC (Fig. 6 (d-e)). A significant decrease of the minor diameter length (shown by red arrow for CF80 sample) indicates the decrease of conductive effect at higher TAN. However, a fraction of polarization leaked through grain boundary and retained ~ 42 % and ~ 46 % of the Pmax at zero field for the samples CF90 and CF100, respectively. It may be mentioned that ferroelectric
11
polarization in BaTiO3 is determined by the switchable polarization PS = 2PR. In semiconductors, leakage component is added in net polarization. The leakage component depends on R-C time constant and electrical conductivity of the samples [33]. Fig. 6(f) compares the P-E loops of the sample CF100, measured at frequencies 1 kHz, 200 Hz and 100 Hz of applied voltage 500 V. The decrease of frequency increases both RP and LP of the sample, e.g., RP increases up to 78 % and 89 %, and LP increases up to 34% and 41 % for frequency at 200 Hz and 100 Hz, respectively. Such increase of the retaining polarization on lowering frequency is controlled by conductivity effects at grain boundaries. We understand such artefact effects on the shape of P-E loop through dielectric properties of the samples. Fig. 7(a-c) shows the frequency (range 1-107 Hz) variation of ac conductivity (σ/), dielectric loss (tand), and dielectric constant (e/). Fig. 7(a) shows that conductivity of the samples increases with measurement frequency. The conductivity of the CF20 and CF50 samples (TAN £ 500 0C) are significantly high throughout the applied frequencies in comparison to the conductivity of the samples annealed at higher temperatures. The frequency activated conductivity in Co ferrite can be explained in terms of hopping of charge carriers via superexchange paths (electron hopping via Fe3+-O2—Fe2+ and hole hopping via Co2+-O2—Co3+) in octahedral lattice sites of the spinel structure [7, 24]. At low frequency regime, the dc limit of conductivity (sdc = nemn + pemh) depends on the density (n: electron density in conduction band, p: hole density in valence band) and mobility of charge carriers (mn: electron mobility, mh: hole mobility). The high conductivity for the samples at low TAN (£ 500 0C) can be attributed to high carrier density and mobility, arising due to surface defects [13]. Fig. 7(b) shows dielectric loss of the samples below 20. The dielectric loss in the high conductive samples is above 1 over a large frequency range. The samples at higher TAN (600-1000 ºC) showed relatively low dielectric loss, typically less than 1 for frequencies above 100 Hz. The dielectric loss exhibited a peak like behaviour in the frequency domain 12
that matches with the average hopping frequency of the charge carriers. The peak frequency can be used to roughly calculate the relaxation time of the R-C circuit for the samples. The decrease of dielectric loss at higher frequencies is expected due to mismatch of the applied frequency and hopping frequency. A rapid increase of the dielectric loss at lower frequencies indicates the space charge contribution to electric polarization at the interfaces of grains and grain boundaries [7, 13]. The dielectric constant (e/ = C/Cair) is another important parameter, which is directly related to the capacitance (C) of the sample. Fig. 7(c) showed a relatively high dielectric constant for the samples at low TAN. Such high value of dielectric constant is contributed by space charge polarization effect. The space charge polarization significantly reduces by increasing the annealing temperature of the samples. The dielectric constant of the CF90 and CF100 samples is nearly stable, indicating extremely low space charge polarization effect and development of polarization in the samples from electric dipoles. Fig. 7 (d-f) has compared the dielectric parameters at 1 kHz (sufficiently away from low frequency space charge regime) for all samples. The most striking feature is that electrical conductivity, dielectric lass and dielectric constant rapidly decreases with the increase of annealing temperature. The conductivity of the material at 1 kHz is extremely small (10-10 S/cm) for the samples with TAN in the range 900 0C-1000 0C, where as dielectric loss and dielectric constant attained the values 0.04-0.1 and 7-8, respectively. The results indicate that the conductive effect rapidly reduces with the increase of annealing temperature. The phase shift (φ) may be a good tool to differentiate capacitive effect from the conductive effect in the samples. This is obtained from measurement of complex impedance components (tanf = J // J/
L
= M , where IC and IR are the current components passing through the capacitance (C) and LG
resistance (R), respectively in RllC circuit model). In the case of purely conductive effect, f is 0 0 (IC = 0). In the case of pure capacitive effect, f is 90 0 (IR = 0). Fig. 7(e)-right Y axis shows that f is small (9-23 0) for the CF20 and CF50 samples. The f has increased to 67 0 for 13
CF60 sample and stabilized in the range 72-88 0 for TAN ³ 800 0C. The high conductivity, high dielectric loss, low phase shift and 100 % leakage of polarization suggest that P-E loops for the samples annealed at £ 600 0C are dominated by conductive effect. On the other hand, capacitive effect has improved electric polarization in the P-E loop of CF90 and CF100 samples, although conductive effect is not completely suppressed in these two samples. We have examined the possibility of further improvement of electric polarization by increasing the annealing time of CF90 sample at 900 0 C from 2 h to 4 h. The P-E loops in Fig. 8 confirm the improvement of stable electric polarization in CF90 sample by increasing its annealing time to 4 h. We obtained a good shaped P-E loop with negligible leakage of polarization (ac voltage up to 1 kV and measurement frequency down to 50 Hz). Although the loops are under saturated, but some of the loops showed a tendency of saturation, irrespective of the magnitude of applied frequency (Fig. 8(a-c)) and voltage (Fig. 8(d-f)). Non-linear dielectrics may exhibit such characteristic features at frequency which matched near to the critical frequency of the R-C equivalent circuit of the sample [33]. In our sample, a good quality P-E loop has been observed for a wide range of frequencies measured below 1 kHz. Second important point is that magnitude of the electric polarization increases with the increase of applied voltage, without loss of polarization at higher electric fields. These typical characters of the P-E loops indicate a possible ferroelectric polarization in the sample, although a fraction of polarization cannot be retained at zero electric field due to intrinsic (grain boundary) disorder and semiconductor nature of the sample. Fig. 9 plots the derived parameters from loops (Pmax, PEmax, PR, electric coercivity (EC), retaining of polarization (RP) at zero field and loss of polarization (LP) at maximum field with respect to the Pmax in positive electric field side). The general feature is that Pmax, PR, EC and RP(%) rapidly decreased at low frequencies (up to 67 Hz) and more or less stabilized at higher frequency. The Pmax is found in the range 4-22 nC/cm2 (Fig. 9(a)). The PR is found in the range 2-11
14
nC/cm2 (Fig. 9(b)). The EC is found in the range 2-15 kV/cm (Fig. 9(c)). The retaining of polarization (RP) at zero field is found in the range 30-60 % (Fig. 9(d)), where as the leakage of polarization at highest field is found in the range 0-8% mostly below 67 Hz (Fig. 9(e)). Since the loops are under saturated, the observed range of ferroelectric parameters (Fig. 9) may be less than actual values for the sample. The ferroelectric polarization in the present material is low due to many reasons. The primary reason is that it is not a traditional ferroelectric, like BaTiO3, where displacement of d0 shell ions (Ti4+) in non-cubic (tetragonal) structure produces large electric polarization [22]. The ferrite is having partially filled 3d shell ions (Fe and Co) that act as the hopping centres for charge transfer and it increases dielectric loss and electrical conductivity. The present material belongs to the class of nontraditional ferroelectric material [3, 21, 30, 32-34]. In such ferroelectrics, a non-negligible displacement of B sites cations and subsequently, charge imbalance between lattice sites can be anticipated during hopping and site exchange of cations. Even, magnetic spin disorder and spin-lattice coupling can produce small ferroelectric polarization [33, 34]. The P-E loops in Co1.25Fe1.75O4 samples are better than that reported for CoFe1.8Ti0.2O4 ferrite [3]. The ferroelectric parameters in our material are comparable to that found in the composite of Ba0.6Sr0.4TiO3 and CoFe2O4 [35] or similar composites [36-39]. The ferroelectric parameters in our sample at room temperature are magnitude wise better than that in a good ferroelectric material MnWO4 with PR (0-40 mC/m2 = 0-4 nC/cm2) and EC (0-300 kV/m = 0-3 kV/cm) at extremely low temperature (< 12 K) [40] and such system may be interesting for basic theory, but does not have much importance for room temperature applications. 3.3. Magnetic properties The magnetic field dependence of magnetization (M(H)) curves in Fig. 10 confirms the ferrimagnetic nature with the exhibition of hysteresis loop in all samples. We have calculated the ferrimagnetic parameters (Coercivity (Hc), spontaneous magnetization (Ms), and
15
squareness (SQ: ratio of remanent magnetization (MR) with Ms) using hysteresis loop. The linear extrapolation of the M2 vs. H/M curve from higher field side on M2 axis was used to determine MS of the samples. The Coercivity (Hc) and remanent magnetization (MR) of the samples were calculated by magnifying the loop in low field scale (inset of Fig. 10(a)). As shown in Fig. 10(b), an initial decreasing trend of MS with increase of annealing temperature could be related to non-equilibrium spinel structure in the samples [13]. The increase of MS value with annealing temperature at ³ 600 0C could be related to the refinement of grain boundary structure. The re-ordering of Co2+ and Fe3+ ions among A and B sites of the spinel structure also affects on the improvement of magnetization with annealing temperature. The observed maximum MS (~ 69 emu/g) in CF100 sample is considerably less than the value of MS ~ 80 emu/g in CoFe2O4 [2], but higher than the MS ~ 55 emu/g in Co1.5Fe1.5O4 [4]. This is due to increasing population of non-magnetic Co3+ ions at B sites as the Co content increases in CoxFe3-xO4 ferrite for x > 1 [14]. Fig. 10(b) shows a monotonic increase of the squareness (SQ) to achieve the maximum value (~ 0.5) for TAN = 900 0C, then it significantly in CF100 sample. The observed SQ values are comparable to ~ 0.47-0.54 reported for CoFe2O4 [2] and ~ 0.33 found in Co1.5Fe1.5O4 [4]. Fig.10(c) shows that magnetic coercivity of the material increased from 159 Oe in CF20 to the maximum value ~ 580 Oe in CF80 sample and then, decreases to 160 Oe in CF100 sample. It can be attributed to transformation of magnetic domain structure from single domain state for TAN below 800 °C (grain size 10-23 nm) to multi domain state for TAN in the range 900-1000 °C (grain size 28-35 nm) by comparing similar variation of squareness and coercivity with annealing temperature in Co ferrite samples [2, 4, 12]. The maximum magnetic coercivity in the present Co1.25Fe1.75O4 ferrite is smaller than the coercivity ~1.2-1.4 kOe for CoFe2O4 at TAN = 800 ºC [12]. It is due to presence of Co3+ ions that reduce magneto-crystalline anisotropy in the spinel structure [16]. . 16
4. Conclusions Heat treatment of the as-prepared sample has brought several changes in the electric and ferrimagnetic properties of Co1.25Fe1.75O4 ferrite. The samples annealed at low temperatures exhibited non-equilibrium spinel structure. The samples annealed at higher temperatures indicated well crystalline spinel structure. The improved micro-structure (large grain size, grain boundary refinement and re-ordering of Co2+ and Fe3+ ions among A and B sites of the spinel structure) by heat treatment played an important role for development of electric polarization and ferrimagnetic parameters in the samples. The present ferrite system belongs to the class of non-traditional ferroelectrics, where hopping of charge carriers between lattice sites and semiconductor nature of the samples introduces electrical conductive paths for leakage of polarization. In this paper, we have shown the decrease of polarization leakage due to conductive effect and improvement of electric polarization due to capacitive effect by increasing the annealing temperature and time. We believe the range of temperature where the material exhibits ferroelectric polarization is more important than the magnitude of electric polarization. The leakage of polarization in the present ferrite samples can be further improved by adopting 3 different techniques, (1) high annealing temperature and time, (2) dispersion of ferrite particles in insulating (ferroelectric) matrix like BaTiO3, (3) formation of laminated composites of ferrite and BaTiO3. We are working on these three aspects for developing the ferrite based multiferroic and multifunctional materials, where ferrite will used as the source of magnetic and ferroelectric component and BaTiO3 will be used for resisting leakage of polarization through grain boundary of ferrite particles and additional source of ferroelectric polarization. Further, the establishment of a coupling between ferrimagnetic order and electric polarization by controlling interfacial lattice structure and interactions in such hybrid multi-ferroics could be useful for the application in magnetic and electric field controlled devices.
17
Acknowledgements We acknowledge the support of Central Instrumentation Facility (CIF) of Pondicherry University for material characterization. The authors thank Ms. V. Vasanthi and Mr. M.R. Panda for technical assistance. RNB thanks DAE-BRNS, Govt. Of India for supporting research Grant (NO. 2011/37P/45/BRNS/2628). RNB thanks UGC, Govt. Of India for supporting research Grant (F.No. 42-804/2013 (SR). References [1] J.A. Moyer, C.A.F. Vaz, E. Negusse, D.A. Arena, and V. E. Henrich, Phys. Rev. B 83, 035121(2011) . [2] A. Franco Jr, and V. Zapf, J. Magn. Magn. Mater. 320, 709 (2008). [3] G. D. Dwivedi, A. G. Joshi, H. Kevin , P. Shahi , A. Kumar, A. K. Ghosh, H. D. Yang, and S. Chatterjee, Solid State Comm. 152, 360 (2012). [4] R. N. Bhowmik, A.T. Sathya, and A. Bharathi, J. Alloys Compd. 559, 134 (2013). [5] W. S. Chiu, S. Radiman, R. A. Shukor, M. H. Abdullah, and P. S. Khiew, J. Alloys Compd. 459, 291 (2008). [6] R.N. Bhowmik, V. Vasanthi, and A. Poddar, J. Alloys Compd. 578, 585 (2013). [7] M. George, S.S. Nair, K.A. Malini, P.A. Joy, and M.R. Anantharaman, J. Phys. D: Appl. Phys. 40, 1593 (2007). [8] V.G. Harris, A. Geiler, Y. Chen, S.D. Yoon, M. Wu, A. Yang, Z. Chen, P. He, P.V. Parimi, X. Zuo, C.E. Patton, M. Abe, O. Acher, and C. Vittoria, J. Magn. Magn. Mater. 321, 2035 (2009). [9] G.S.N. Rao, O.F. Caltun, K.H. Rao, P.S.V. Subba Rao, and B.P. Rao, J. Magn. Magn. Mater. 341, 60 (2013). [10] K. Inomata, N. Ikeda, N. Tezuka, R. Goto, S. Sugimoto, M. Wojcik, and E. Jedryka, Sci. Technol. Adv. Mater. 9, 014101 (2008).
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[11] K.F. Wang, J.-M. Liu, and Z.F. Ren, Advances in Physics 58, 321 (2009). [12] A. Lunhong, and J. Jiang, Current Appl. Phys. 10, 284 (2010). [13] R.N. Bhowmik, Mater. Res. Bull. 50, 476 (2014). [14] G.H. Jonker, J. Phys. Chem. Solids 9 (1959) 165. [15] H. L. Trong, L. Presmanes, E. D. Grave, A. Barnabe, C. Bonningue, and Ph. Tailhades, J. Magn. Magn. Mater. 334, 66 (2013). [16] I. C. Nlebedim, A. J. Moses, and D. C. Jiles, J. Magn. Magn. Mater. 343, 49 (2013). [17] I. P. Muthuselvam, and R.N. Bhowmik, Solid State Sci. 11, 719 (2009). [18] N. Bahlawane, P.V.T. Ngamou, V. Vannier, T. Kottke, J. Heberle, and K.K. -Hoinghaus, Phys. Chem. Chem. Phys. 11, 9224 (2009). [19] M.R. Panda, R.N Bhowmik, H. Singh, M.N. Singh, and A K Sinha, Mater. Res. Exp. 2, 036101 (2015). [20] I. C. Nlebedim, N. Ranvah, P. I. Williams, Y. Melikhov, J. E. Snyder, A. J. Moses, and D. C. Jiles, J. Magn.Magn. Mater. 322, 1929 (2010). [21] A. Das, G.D. Dwivedi, P. Kumari, P. Shahi, H. D. Yang, A. K. Ghosh, and S. Chatterjee, J. Magn. Magn. Mater. 379, 6 (2015). [22] D. Khomskii, Physics 2, 20 (2009). [23] R.N. Bhowmik, M.R. Panda, S.M. Yusuf, M.D. Mukadam, A.K. Sinha, J. Alloys Compd. 646, 161 (2015). [24] R.N. Bhowmik, and I. Panneer Muthuselvam, J. Alloys Compd. 589, 247 (2014). [25] P.L. Meena, R. Kumar, C.L. Prajapat, K. Sreenivas, and V. Gupta, J. Appl. Phys. 106, 024105 (2009). [26] P. Chandramohan, M.P. Srinivasan, S. Velmurugan, and S.V. Narasimhan, J. Solid State Chem. 184, 89 (2011).
19
[27] Z. Ž. Lazarević, Č. Jovalekić, D. Sekulić, M. Slankamenac, M. Romčević, A. Milutinović, and N. Ž. Romčević, Science of Sintering 44, 331 (2012). [28] S. Grace Victoria, A. M. E. Raj, and C. Ravidhas, Mater. Chem. Phys. 162, 852 (2015). [29] A.C. Albrecht, The J. Chem. Phys. 34, 1476 (1961). [30] J. R.-Fuertes, T. Bernert, M. He, B. Winkler, V.L. Vinograd, and V. Milman, Appl. Phys. Lett. 105, 071911 (2014). [31] B.R. Myoung, and C.S. Kim, J. Appl. Phys. 117, 17B741 (2015). [32] S.V. Streltsov, A.I. Poteryaev, and A.N. Rubtsov, J. Phys.: Condens. Mater 27, 165601 (2015). [33] K.M. Rabe, M. Dawber, C. Lichtensteiger, C.H. Ahn, and J.-M. Triscone (Eds.): Physics of Ferroelectrics: A Modern Perspective, Topics in Applied Physics 105, 1-30 (2007) [34] A.G. Lone, and R.N. Bhowmik, J. Magn. Magn. Mater. 379, 244 (2015). [35] W.Y. An, W. Y. Bo, R. Wei, G. J. Xiong, and Z. W. Li, Chin. Phys. Lett. 29, 067701 (2012). [36] L. B. Hao, D. X. Zhou, S. P. Gong, Q. Y. Fu, W. Luo , G. Jian, F. Xue, and L. Zhou, J Mater Sci: Mater Electron 24, 2351(2013). [37] S. T. Zhang , Y. Zhang, Z. L. Luo, M. H. Lu, Z. B. Gu, and Y. F. Chen, Applied Surface Science 255, 5092 (2009). [38] J. Nie, G. Xu, Y. Yang, and C. Cheng, Mater. Chem. Phys. 115, 400 (2009). [39] L.Y. Ding, F. X. Wu, Y. B. Chen, Z. B. Gu, and S.T. Zhang, Appl. Surf. Sci. 257, 3840 (2011). [40] B. Kundys, C. Simon, and C. Martin, Phys. Rev. B 77, 172402 (2008).
20
Figure captions: Fig. 1 (Colour online) SXRD pattern and profile fitted data of the annealed samples. Fig. 2 (Colour online) (a) Fitting of the (440) peak using Gaussian curves. (b) fitting of the peak parameters according to Williamson-Hall equation. (c) variation of grain size and lattice strain (%) with annealing temperature of the sample. Fig.3 (Colour online) SEM images (a-d), EDX data- point analysis (e-f) and EDX data- line scan analysis (g-h) shown for selected samples. Fig. 4 (Colour online) Raman spectra of the samples, along with fitted data (lines). Fig. 5 (Colour online) (a-c) peak parameters from Raman spectra. (d) SXRD (331) peak, (e) (331) and (400) SXRD peak ratio, (f) A1g(1) and A1g(2) peak ratio (Raman spectra). Fig. 6 (Colour online) (a-e) P-E loop measured for different samples at 1 kHz loop frequency. (e) P-E loop measured for CF100 sample at different loop frequencies. Fig. 7 (Colour online) (a-c) frequency dependence of dielectric parameters for the samples annealed at different temperatures. (d-f) dielectric parameters of the samples at 1 kHz and 1 Hz with the variation of annealing temperature. Fig. 8 (Colour online) P-E loops of the CF90 sample annealed for 4 hrs in different modes by varying the measurement loop frequency and applied voltage. Fig. 9 (Colour online) Variation of the parameters with measurement loop frequencies and the parameters have been derived from the P-E loops in Fig. 8. Fig. 10 (Colour online) (a) Field dependence of magnetization curves. Annealing temperature dependence of magnetic parameters (b-c).
21
Highlights Ø Ø Ø Ø Ø
Nanoparticles of Co1.25Fe1.75O4 ferrite have been prepared in alkaline medium. Nanoparticles have been annealed to tailor magnetic and ferroelectric properties. Room temperature ferrimagnetism is confirmed in all samples. Some of the samples showed good signature of ferroelectric properties. Role of conductive effect and capacitive effect understood.
22
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