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Effects of Co doping on structural, magnetic, and electrical properties of 0.6BiFeO3-0.4(Bi0.5K0.5)TiO3 solid solution
Journal of Alloys and Compounds 730 (2018) 119e126
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Effects of Co doping on structural, magnetic, and electrical properties of 0.6BiFeO3-0.4(Bi0.5K0.5)TiO3 solid solution Y.F. Qin a, b, J. Yang a, *, W.J. Huang a, b, P. Xiong a, b, J.Y. Song a, b, D. Wang a, b, L.H. Yin a, W.H. Song a, P. Tong a, X.B. Zhu a, Y.P. Sun a, c, d a
Key Laboratory of Materials Physics, Institute of Solid State Physics, Chinese Academy of Sciences, Hefei, 230031, People's Republic of China University of Science and Technology of China, Hefei, 230026, People's Republic of China High Magnetic Field Laboratory, Chinese Academy of Sciences, Hefei, 230031, People's Republic of China d Collaborative Innovation Center of Advanced Microstructures, Nanjing University, Nanjing, 210093, People's Republic of China b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 July 2017 Received in revised form 26 September 2017 Accepted 27 September 2017 Available online 28 September 2017
The effects of Co doping on structural, dielectric, and magnetic properties of multiferroic 0.6BiFe1-xCoxO30.4(Bi0.5K0.5)TiO3 (0 x 0.2) ceramics have been investigated. The samples undergo a structural transition from pseudocubic to rhombohedral phase at x ¼ 0.1. The maximum remnant polarization of 57.8 mC/cm2 and remnant magnetization of 0.107 emu/g are observed in the x ¼ 0.15 sample. The ferromagnetism is proposed to originate from the suppression of spiral spin structure with the canting of the antiferromagnetically ordered spins caused by the substitution of Co for Fe ions. Moreover, the magnetodielectric effect can be observed in the x ¼ 0.15 sample. © 2017 Elsevier B.V. All rights reserved.
Keywords: Multiferrocity Solid-solution Magnetic properties BiFeO3
1. Introduction Mutiferroic materials which simultaneously exhibit ferroelectricity, ferromagnetism or ferroelasticity have attracted much attention due to the potential applications in switches, sensors, and memory storage devices etc. [1-4]. BiFeO3 (BFO) is considered to be the most promising room-temperature (RT) single-phase mutiferroic material because BiFeO3 possesses a high ferroelectric (FE) el temperature of Curie temperature of TC ~ 1103 K, and Ne TN ~ 643 K [5,6]. However, the antiferromagnetic spin structure of BiFeO3 is G-type, and this spatially-modulated cycloidal-spin structure does not allow the appearance of net magnetization, which inhibits the observation of a notable linear magnetoelectric (ME) effect in BiFeO3 [7]. Moreover, fabrication of the bulk BiFeO3 still remain some problems of impurity phase and high leakage current, which result in a poor ferroelectric hysteresis loop [8,9]. Element substitution is a useful way to suppress the impurity phase and break up the cycloidal spin structure to induce net magnetization [10]. The formation of solid solutions was also found to be
* Corresponding author. E-mail address: [email protected] (J. Yang). https://doi.org/10.1016/j.jallcom.2017.09.295 0925-8388/© 2017 Elsevier B.V. All rights reserved.
beneficial to obtain room-temperature multiferroic properties. Pal et al. demonstrated a large enhancement of ferromagnetism in modified BiFeO3-BaSrTiO3 ceramics [11]. Mandal et al. observed a switchable polarization and magnetization at RT in BiFeO3Bi(Ti0.5Mg0.5)O3-CaTiO3 solid solution [12]. These studies indicate that the BFO-based solid solutions are very attractive to realize multiferroic properties in view of their structure-property correlations. Recently, many studies show that the solid solutions of (1-x) BFO-x(Bi0.5K0.5)TiO3 (BKT) undergo a structural transition from R3c phase to P4mm phase with an increase in the content of BKT and exhibit a coexistence of R3c phase and P4mm phase when the amount of BKT is 40% [13-15]. A relatively large remnant polarization (Pr) of 52 mC/cm2 can be obtained in 0.6BFO-0.4BKT [13]. Moreover, the dielectric constant of BFO-BKT solid solutions show a broad maximum with frequency dispersion feature, which is one of characteristic of ralaxor ferroelectrics [16-18]. Hagiwara et al. reported that CuO-doped 0.6BFO-0.4BKT exhibits well-square hysteresis loops by means of thermal quenching and there is no ferromagnetism at room temperature [19]. In summary, previous studies mostly focused on the structural, ferroelectric and dielectric properties of 0.6BFO-0.4BKT solid solutions, whereas the magnetic properties of 0.6BFO-0.4BKT solid solutions have rarely been
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investigated the magnetic, ferroelectric and dielectric properties of 0.6BF1-xCxO-0.4BKT samples. The results show that the structural transition from pseudocubic to rhombohedral phase occurs in the x ¼ 0.1 sample and a coexistence of ferroelectricity and ferromagnetism at RT can be realized in the x ¼ 0.15 sample. The results indicate that it is an effective way to explore RT multiferroic compounds through a binary solid-solution. 2. Experimental procedures 2.1. Sample preparation
Fig. 1. XRD patterns of 0.6BF1-xCxO-0.4BKT (0 x 0.2) samples. The inset shows the amplified plots of the main diffraction (110) peaks.
revealed in literatures. In this work, we prepared 0.6BiFe1-xCoxO3-0.4(Bi0.5K0.5)TiO3 (0.6BF1-xCxO-0.4BKT) (0 x 0.2) by modified Pechini method at a relatively low temperature of 800 C. And we systematically
The 0.6BF1-xCxO-0.4BKT (0 x 0.2) ceramics were prepared through modified Pechini method, which could maintain atomiclevel mixing metal cations to lower sintering temperature for sake of minimization of the volatility of Bi and K. All the chemicals used in the experimental process were analytical grade and purchased from Sinopharm Chemical Reagent Co., Ltd. Stoichiometric amounts of the starting materials, Ti[OCH(CH3)2]4, KNO3, Bi(NO3)3$5H2O, Fe(NO3)3$9H2O, and Co(CH3COO)2$4H2O were dissolved into the citric acid monohydrate C6H8O7$H2O solution with glycol and acetic acid in stoichiometric proportions. Citric acid monohydrate C6H8O7$H2O was added to the solution as a complexant. Glycol was used for adjusting the viscosity and stability of the sol. Ammonia was added into the mixture of these solutions to adjust pH to 7 and the mixture was stirred until transparent precursor sol was obtained. Then the precursor solution was heated on a hot plate at 90 C resulting in the formation of a gel. The gel was heated at 400 C to remove the organic species. The obtained
Fig. 2. Rietveld refinement results for 0.6BF1-xCxO-0.4BKT (0 x 0.15) samples with (a) x ¼ 0 (b) x ¼ 0.05 (c) x ¼ 0.1, and (d) x ¼ 0.15. The insets are their corresponding amplified plots of (111) peak with 2q around 39 .
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powders were mixed and grounded, and then calcined at 700 C for 10 h under an O2 flow. The obtained powders were again pelletized, and sintered at 800 C for 20 h under an O2 flow with two intermediate grindings, and finally the furnace was cooled slowly to room temperature.
2.2. Characterizations The x-ray powder diffraction patterns were recorded at RT on a Panalytical diffractometer (X'Pert PROMRD) with Cu-Ka radiation and a graphite monochromator in a reflection mode. Structural refinements of 0.6BF1-xCxO-0.4BKT (0 x 0.15) are carried out by using RIETICA software. The morphological and structural information of all samples were characterized with a field emission scanning electron microscopy (FE-SEM, FEI-designed Sirion 200, Hillsboro, OR). The temperature dependence of magnetization was carried out from 300 to 800 K using a Physical Property Measurement System-vibrating sample magnetometer (PPMS-VSM) with a high-temperature oven option. The magnetization hysteresis measurements were performed with a Quantum Design superconducting quantum interface device (SQUID) magnetic property
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measurement system (MPMS) system (2 T 400 K, 0 H 4.5 T). Temperature dependence of dielectric constant and loss tangent ranging from 300 to 900 K were measured using an LCR meter in the frequency range of 1 kHze1000 kHz. The ferroelectric properties were investigated using a Sawyer-Tower circuit attached to a computer-controlled standardized ferroelectric test system (Radiant Technology 609B) with probes on the electrodes of the samples, which were immersed in silicone oil to prevent arcing.
3. Results and discussion Fig. 1 shows the XRD patterns of 0.6BF1-xCxO-0.4BKT (0 x 0.2) at RT. There is no impurity phase can be observed in the 0 x 0.15 samples, whereas small impurity peaks marked by stars at around 30 corresponding to Bi12(Bi0.5Fe0.5)O19.5 are observed in the x ¼ 0.2 sample. With increasing the doping level of Co, the main diffraction peak (110) shifts toward a higher angle position, implying the decrease in lattice parameters due to the smaller ionic radius of Co3þ (0.61 Å) compared with Fe3þ (0.645 Å), which is shown in the inset of Fig. 1. We refined the structural parameters of 0.6BF1-xCxO-0.4BKT (0 x 0.15) by using RIETICA
Fig. 3. Representative FE-SEM images of 0.6BF1-xCxO-0.4BKT samples with (a) x ¼ 0, (b) x ¼ 0.05, (c) x ¼ 0.1, (d) x ¼ 0.15, and (e) x ¼ 0.2.
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Fig. 4. XPS spectra and the deconvoluted results of (a)e(d) Fe 2 p of the x ¼ 0, 0.05, 0.1 and 0.15 samples, respectively, (e) Co 2 p of the x ¼ 0.05, 0.1 and 0.15 samples, and (f) Ti 2 p of the x ¼ 0, 0.05, 0.1 and 0.15 samples. The solid curves denote the fitting results.
software, as shown in Fig. 2(a)-(d). It can be seen that the fits between the experimental and calculated XRD patterns are relatively good based on the consideration of the small RP values of 3.37%, 3.96%, 3.32%, and 3.70% for the samples with x ¼ 0, 0.05, 0.1 and 0.15, respectively. For the x ¼ 0 and 0.05 samples, a single and no splitting (111) peak at about 39 is observed (inset of Fig. 2(a), (b)), and the (002) peak remains a single and sharp shape implying the existence of pseudocubic lattice. Whereas for the x ¼ 0.1 and 0.15 samples, the splitting of (111) is obviously observed implying that a rhombohedral lattice begins to appear at this composition, as shown in the inset of Fig. 2(c) and (d). Fig. 3 illustrates the field-emission scanning electron micrograph (FE-SEM) of 0.6BF1-xCxO-0.4BKT (0 x 0.2). Small grain size and rather loose morphology are observed in the x ¼ 0 and 0.05 samples. Whereas the x ¼ 0.1 and 0.15 samples exhibit relatively uniform and dense morphology without obvious porosity. The grain size of the x ¼ 0.1 and 0.15 samples are much larger than those of the x ¼ 0 and 0.05 samples. The particles of the x ¼ 0.2 sample have near-quadrate feature, which are much different from those of the x ¼ 0.1 and 0.15 samples. The different morphology of the x ¼ 0.2 sample might be due to the existence of impurity phase. The changes in morphology of 0.6BF1-xCxO-0.4BKT (0 x 0.15) ceramics indicate that Co doping can significantly lower the sintering temperature of the ceramics and promote the growth of the grains which in favor of enhancing the grain connectivity. To study the valence states of the ions of 0.6BF1-xCxO-0.4BKT (0 x 0.15) ceramics, we performed the XPS measurements for the 0 x 0.15 samples. The binding energy of the Fe 2p3/2 peaks locates at the range of 710.5e710.6 eV. To obtain the Fe3þ/Fe2þ ratio, the Fe 2p3/2 peaks were deconvoluted and fitted, as shown in Fig. 4(a)-(d). The content of Fe3þ are 36.7%, 39.3%, 50.1% and 56.8% for the x ¼ 0, 0.05, 0.1 and 0.15 samples, respectively. It is obvious that the content of Fe3þ increases gradually with the increase of the
Co concentration which indicates Co doping in the solid solution can effectively reduce the oxygen vacancies. The core-level spectra of Co 2p3/2 and the deconvoluted results for the x ¼ 0.05, 0.1 and 0.15 samples are shown in Fig. 4(e). It is found that peaks locates at 780.4, 780.4 and 780.8 eV for the x ¼ 0.05, 0.1 and 0.15 samples, respectively, which are very close to that of Co 2p3/2 (781.30 eV) in Co2O3 implying that the doped Co ions in the samples are in þ3 valence state [20]. The Ti 2p3/2 peaks located at 457.8e457.9 eV are consistent with the binding energy of Ti 2p3/2 (458.20 eV) in TiO2 [21], as shown in Fig. 4(f), which indicates Ti ions in our samples are in þ4 valence state. In order to make clear the magnetic behaviors of 0.6BF1-xCxO0.4BKT (0 x 0.15), we performed the magnetic hysteresis measurements at 300 K, as shown in Fig. 5. The M-H curve of the x ¼ 0 sample exhibits a linear behavior implying the characteristic of paramagnetic (PM) or AFM states. The M-H curve of the x ¼ 0.05 sample shows the same behavior as x ¼ 0 sample. With further increasing the content of Co, a small Mr value can be observed in the
Fig. 5. RT ferromagnetic hysteresis loops curves of the x ¼ 0, 0.05, 0.1 and 0.15 samples. The right inset is the corresponding amplified curves.
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Fig. 6. Temperature dependence of magnetization in the ZFC and FC modes measured at H ¼ 100 Oe for the samples with (a) x ¼ 0 and (b) x ¼ 0.15. The inset of Fig. 6(b) shows the plot of dM/dT.
x ¼ 0.1 sample as shown in the inset of Fig. 5. A large hysteresis loop with a remnant magnetization Mr of 0.107 emu/g and coercive field Hc of 300 Oe can be observed in the x ¼ 0.15 sample and these values are similar to those in the other solid solutions [22]. The observed remnant magnetization Mr in the x ¼ 0.15 sample is considered to be intrinsic property of the ceramic rather than from impurity phase because the XRD result has ruled out the possible of impurity phases. Moreover, the contribution of magnetism from the impurity Bi12(Bi0.95Fe0.05)O19.5 or Bi25FeO39 can also be
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excluded, due to the PM nature of the impurity phases [23,24]. The magnetization of the x ¼ 0.15 sample does not reach saturation when the magnetic field is up to H ¼ 40 kOe implying the coexistence of ferromagnetic (FM) and AFM interactions in the x ¼ 0.15 sample. A liner extrapolation of M-H to H ¼ 0 allows us to determine the spontaneous magnetization Ms of the x ¼ 0.15 sample as 0.0168 mB/f.u. This value is much smaller than the mean spin only value of isolated Fe3þ (2.5 mB), and Co3þ (2 mB) ions implying that the FM did not originate from the Fe or Co clusters. It is known that the bulk BiFeO3 possesses a spiral spin structure below the antiferromagnetic transition temperature TN ~643 K. Many previous studies manifest that the weak ferromagnetism with M-H hysteresis loops can be generated by decreasing the grain size below 62 nm or doping at Bi or Fe site due to the suppression of the spiral structure with the canting of the antiferromagnetic ordered spins. In this work, the ferromagnetism is proposed to originate from the suppression of spiral spin structure with the canting of the antiferromagnetically ordered spins caused by the substitution of Co for Fe ions [22,25,26]. From the XPS data, we can obtain the content of Fe3þ in the samples and therefore calculate the proportion of Co3þ to Fe3þ. It is found that the ratio of Co3þ to Fe3þ are 0.13, 0.22 and 0.31 for the samples with x ¼ 0.05, 0.1 and 0.15, respectively. This value of the x ¼ 0.15 sample is the largest among the Co-doped samples. In other words, the contribution of a larger amounts of Co3þ-Fe3þ antiferromagnetic ordered spins exists in the x ¼ 0.15 sample. The canting of these antiferromagnetic ordered spins is in favor of generation of the weak ferromagnetism and the x ¼ 0.15 sample exhibits larger remnant magnetization. To further understand the magnetic properties, we performed the measurements of temperature dependence of magnetization M (T) for the selected two samples with x ¼ 0 and 0.15 in the zero-field cooled (ZFC) and field cooled (FC) modes under a magnetic field of 100 Oe. As shown in Fig. 6(a), the magnetization of both ZFC and FC curves for the x ¼ 0 sample increase monotonously with the decrease in
Fig. 7. RT ferroelectric hysteresis loops of 0.6BF1-xCxO-0.4BKT (0 x 0.15) samples with (a) x ¼ 0, (b) x ¼ 0.05, (c) x ¼ 0.1, and (d) x ¼ 0.15.
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Fig. 8. Temperature dependence of dielectric constants for the samples with (a) x ¼ 0, (b) x ¼ 0.05, (c) x ¼ 0.1, and (d) x ¼ 0.15. The insets are their corresponding Vogel-Fulcher plots of relaxation frequencies vs. temperatures.
temperature and no long-range magnetic transition is observed implying the PM nature of the x ¼ 0 sample, which is consistent with the M-H curve. For the x ¼ 0.15 samples, the ZFC and FC curves almost overlap with each other in the high temperature range. As the temperature goes down, it seems to undergo a PM-to-FM magnetic transition at 683 K (defined as the temperature corresponding to the dip of dM/dT) as shown in the inset of Fig. 6(b). In the low temperature range below 660 K, there is a deviation between the ZFC and FC magnetization curves. Fig. 7 shows the polarization-electric field (P-E) hysteresis loops of 0.6BF1-xCxO-0.4BKT (0 x 0.15) measured at frequency of 100 Hz at RT. The P-E hysteresis loop of the x ¼ 0 sample exhibits a pinched loop with a remnant polarization Pr of 27 mC/cm2 in spite of a large applied-electric field of 160 kV/cm. The P-E loop of the x ¼ 0.05 sample is not well square enough. The poor P-E loops of the x ¼ 0 and 0.05 samples might be due to the loose morphology giving rise to high leakage current densities. However, well-defined P-E loops with Pr values of 47.5 mC/cm2 and 57.8 mC/cm2 are obtained for the x ¼ 0.1 and 0.15 samples, which is in the similar magnitude of other ferroelectric solid solution [13]. Normally, the large polarization is observed in epitaxy or exceptionally high quality single-crystal BiFeO3. However, it is rather difficult to obtain the large polarization in BiFeO3 bulk. In this work, the superior polarization in the Co-doped 0.6BFO-0.4BKT are mainly due to the formation of the solid solution of BKT, giving rise to the increase of grain size and the formation of dense morphology, which are in favor of increasing the remnant polarization [27]. Moreover, the introduction of Co3þ ions in 0.6BFO-0.4BKT can effectively lower the sintering temperature of the ceramics resulting in reduction of Bi volatilization and oxygen vacancies, which is in good agreement with the variation in the Fe3þ content from the XPS results. Figs. 8 and 9 show the temperature dependence of dielectric constant ε0 and dielectric loss tan d of 0.6BF1-xCxO-0.4BKT (0 x 0.15) at different frequencies from 300 to 900 K. For the x ¼ 0.1 sample, the dielectric constant peaks are broad, and the maximum in ε0 are diffusive accompanied by a strong dispersion of dielectric maximum temperature (Tm) with the variation of frequencies. With increasing the frequency, the dielectric maximum
decreases, while the Tm increases from 690 K at 10 kHz to 740 K at 1000 kHz showing a relaxation behavior which is similar to previous report [14]. With increasing the doping level of Co, the ε0 peaks become more obvious, and the maximum values of dielectric constant ε0 are much higher than those of the x ¼ 0 sample. As shown in Fig. 9, the values of tan d of the x ¼ 0.05, 0.1 and 0.15 samples is slightly smaller than that of the x ¼ 0 sample at RT, which are in agreement with the better polarization in Co doped samples. We note that the frequency-dependent peaks turn into sharp and the dielectric relaxor behavior become weaken with increasing the content of Co, which might be due to the structural transition from pseudocubic to rhombohedral phase associated with the domain switched from nanodomains to macrodomains [28], as confirmed by the XRD results. The observed relaxor behaviors in 0.6BF1-xCxO-0.4BKT (0 x 0.15) ceramics can be ascribed to chemical disorder resulting from mixed-valence cation-site occupation, which involved Bi3þ and Kþ at A-site, Ti4þ, Co3þ and Fe3þ at B-site of ABO3 perovskite lattice. These phenomena have also been observed in other ferroelectric compounds, such as (Bi0.5K0.5)TiO3 [29], (Ba0.6Bi0.2Li0.2)TiO3 [30]. The dielectric data were analyzed in terms of the Vogel-Fulcher relationship [31,32]: f ¼ f0 exp (Ea/kB (TmTvF)),
(1)
where f ¼ 1/t is the frequency, f0 is a pre-exponential factor (1/t0), Ea is the activation energy describing the polarization fluctuations of an isolated cluster relaxation process, Tm is the dielectric maximum in absolute temperature, T VF is the Vogel-Fulcher freezing temperature, and kB is the Boltzmann constant. The insets of Fig. 8 show the plot between ln f and T (K), where the dotted points represent experimental data and the solid lines are the fitting results according to Eq. (1). The values of T VF, f0, and Ea of x 0.15 samples are presented in Table 1. Ea values of the x 0.15 samples are similar to those of weakly-coupled relaxor ferroelectrics such as (Bi0.5K0.5)TiO3-BiScO3 [33] and BaTiO3-BiScO3 [34]. The existence of a magnetodielectric (MD) effect is often used as a test for ME coupling in multiferroic material. Therefore, we
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Fig. 9. Temperature dependence of tangent loss for the samples with (a) x ¼ 0, (b) x ¼ 0.05, (c) x ¼ 0.1, and (d) x ¼ 0.15.
Table 1 Vogel Fulcher fitting parameters for 0.6BF1-xCxO-0.4BKT (0 x 0.15) (Tm at 100 kHz). Nos.
Materials
1 2 3 4
x x x x
¼ ¼ ¼ ¼
0 0.05 0.1 0.15
TvF (K)
f0 (Hz)
Ea (eV)
Tm (K)
554 560 564 540
1.3 109 3.2 109 8.5 1010 15.5 1010
0.23 0.20 0.21 0.24
716 698 686 677
performed the measurement of the field-dependent dielectric constant ε0 at a frequency of 1 kHz at RT, as shown in Fig. 10. The dielectric constant gradually increases with increasing the applied magnetic fields implying the possible existence of MD effect. The MD effect can be defined as: MD ¼ [ε0 (H) - ε0 (0)] 100%/ε0 (0),
where ε0 (0) and ε0 (H) denote the dielectric constants under zero field and magnetic field, respectively. The magnitude of MD effect under 50 kOe at RT can be calculated as 0.03% for the x ¼ 0.15 sample. 4. Conclusions In conclusion, we have successfully prepared 0.6BF1-xCxO0.4BKT (0 x 0.2) ceramics. The Co doped samples undergo a structural transition from pseudocubic to rhombohedral phase at x ¼ 0.1. The FE-SEM results manifest that the substitution of Co3þ ions can significantly promote the grain growth resulting in dense morphology. The x ¼ 0.15 sample exhibits a maximum remnant magnetization of 0.107 emu/g and a superior remnant polarization of 57.8 mC/cm2. The ferromagnetism is proposed to originate from the suppression of spiral spin structure with the canting of the antiferromagnetically ordered spins caused by structural distortion due to the substitutions of Co for Fe ions. Dielectric studies show the feature of relaxor ferroelectrics resulting from mixed valence cation-site occupation. In addition, the MD effect (0.03% at RT) can be observed in the x ¼ 0.15 sample. Acknowledgments This work was supported by the National Key Research and Development Program of China (2017YFA0403502) and Key Research Program of Frontier Sciences, CAS (QYZDB-SSW-SLH015). References
Fig. 10. RT magnetic field dependence of MD effect at a frequency of 1 kHz for the x ¼ 0.15 sample. The black solid line arrows indicate the field sweeping direction.
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Soc. 98 (2015) 469e475. [20] B.J. Tan, K.J. Klabunde, P.M.A. Sherwood, XPS studies of solvated metal atom dispersed (SMAD) catalysts. Evidence for layered cobalt-manganese particles on alumina and silica, J. Am. Chem. Soc. 113 (1991) 855. [21] A.P. Dementjev, O.P. Ivanova, L.A. Vasilyev, A.V. Naumkin, D.M. Nemirovsky, D.Y. Shalaev, Altered layer as sensitive initial chemical state indicator, J. Vac. Sci. Technol. A 12 (1994) 423. [22] T. Kanai, S. Ohkoshi, A. Nakajima, T. Watanabe, K. Hashimoto, A ferroelectric ferromagnet composed of (PLZT)(x)-(BiFeO3)(1-x) solid solution, Adv. Mater. 13 (2001) 487e490. [23] Q. Xu, H. Zai, D. Wu, T. Qiu, M.X. Xu, The magnetic properties of Bi(Fe0.95Co0.05) O3 ceramics, Appl. Phys. Lett. 95 (2009) 112510. [24] A.A. Zatsiupa, L.A. Bashkirov, I.O. Troyanchuk, G.S. Petrov, A.I. Galyas, L.S. Lobanovsky, S.V. Truhanov, Magnetization, magnetic susceptibility, effective magnetic moment of Fe3þ ions in Bi25FeO39 ferrite, J. Solid State Chem. 212 (2014) 147e150. [25] K. Singh, R.K. Kotnala, M. Singh, Study of electric and magnetic properties of (Bi0.9Pb0.1)(Fe0.9Ti0.1)O3 nanomultiferroic system, Appl. Phys. Lett. 93 (2008) 212902. [26] M.M. Kumar, S. Srinath, G.S. Kumar, S.V. Suryanarayana, Spontaneous magnetic moment in BiFeO3-BaTiO3 solid solutions at low temperatures, J. Magn. Magn Mater. 188 (1998) 203e212. [27] C.A. Randall, N. Kim, J.P. Kucera, W.W. Cao, T.R. Shrout, Intrinsic and extrinsic size effects in fine-grained morphotropic-phase-boundary lead zirconate titanate ceramics, J. Am. Ceram. Soc. 81 (1998) 677e688. [28] L. Cao, C. Zhou, Q. Fan, W. Zeng, Q. Zhou, H. Yang, Unusual relaxor-normal ferroelectric crossover in Cu-doped BiFeO3-BaTiO3 ceramics, J. Mater. Sci. Mater. Electron. 26 (2015) 3610e3614. [29] J. Yang, Y. Hou, C. Wang, M. Zhu, H. Yan, Relaxor behavior of (K0.5Bi0.5)TiO3 ceramics derived from molten salt synthesized single-crystalline nanowires, Appl. Phys. Lett. 91 (2007) 023118. [30] H. Borkar, V. Rao, S. Dutta, A. Barvat, P. Pal1, M. Tomar, V. Gupta, J.F. Scott, A. Kumar, Novel optically active lead-free relaxor ferroelectric (Ba0.6Bi0.2Li0.2) TiO3, J. Phys. Condens. Matter 28 (2016) 265901. [31] C.J. Stringer, T.R. Shrout, C.A. Randall, High-temperature perovskite relaxor ferroelectrics: a comparative study, J. Appl. Phys. 101 (2007) 054107. [32] R. Pirc, R. Blinc, Vogel-Fulcher freezing in relaxor ferroelectrics, Phys. Rev. B 76 (2007) 020101. [33] C. Kruea-In, G. Rujijanagul, F.Y. Zhu, S.J. Milne, Relaxor behaviour of K0.5Bi0.5TiO3-BiScO3 ceramics, Appl. Phys. Lett. 100 (2012) 202904. [34] H. Ogihara, C.A. Randall, S. Trolier-McKinstry, Weakly coupled relaxor behavior of BaTiO3-BiScO3 ceramics, J. Am. Ceram. Soc. 92 (2009) 110e118.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effects of Co doping on structural, magnetic, and electrical properties of 0.6BiFeO3-0.4(Bi0.5K0.5)TiO3 solid solution Y.F. Qin a, b, J. Yang a, *, W.J. Huang a, b, P. Xiong a, b, J.Y. Song a, b, D. Wang a, b, L.H. Yin a, W.H. Song a, P. Tong a, X.B. Zhu a, Y.P. Sun a, c, d a
Key Laboratory of Materials Physics, Institute of Solid State Physics, Chinese Academy of Sciences, Hefei, 230031, People's Republic of China University of Science and Technology of China, Hefei, 230026, People's Republic of China High Magnetic Field Laboratory, Chinese Academy of Sciences, Hefei, 230031, People's Republic of China d Collaborative Innovation Center of Advanced Microstructures, Nanjing University, Nanjing, 210093, People's Republic of China b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 July 2017 Received in revised form 26 September 2017 Accepted 27 September 2017 Available online 28 September 2017
The effects of Co doping on structural, dielectric, and magnetic properties of multiferroic 0.6BiFe1-xCoxO30.4(Bi0.5K0.5)TiO3 (0 x 0.2) ceramics have been investigated. The samples undergo a structural transition from pseudocubic to rhombohedral phase at x ¼ 0.1. The maximum remnant polarization of 57.8 mC/cm2 and remnant magnetization of 0.107 emu/g are observed in the x ¼ 0.15 sample. The ferromagnetism is proposed to originate from the suppression of spiral spin structure with the canting of the antiferromagnetically ordered spins caused by the substitution of Co for Fe ions. Moreover, the magnetodielectric effect can be observed in the x ¼ 0.15 sample. © 2017 Elsevier B.V. All rights reserved.
Keywords: Multiferrocity Solid-solution Magnetic properties BiFeO3
1. Introduction Mutiferroic materials which simultaneously exhibit ferroelectricity, ferromagnetism or ferroelasticity have attracted much attention due to the potential applications in switches, sensors, and memory storage devices etc. [1-4]. BiFeO3 (BFO) is considered to be the most promising room-temperature (RT) single-phase mutiferroic material because BiFeO3 possesses a high ferroelectric (FE) el temperature of Curie temperature of TC ~ 1103 K, and Ne TN ~ 643 K [5,6]. However, the antiferromagnetic spin structure of BiFeO3 is G-type, and this spatially-modulated cycloidal-spin structure does not allow the appearance of net magnetization, which inhibits the observation of a notable linear magnetoelectric (ME) effect in BiFeO3 [7]. Moreover, fabrication of the bulk BiFeO3 still remain some problems of impurity phase and high leakage current, which result in a poor ferroelectric hysteresis loop [8,9]. Element substitution is a useful way to suppress the impurity phase and break up the cycloidal spin structure to induce net magnetization [10]. The formation of solid solutions was also found to be
* Corresponding author. E-mail address: [email protected] (J. Yang). https://doi.org/10.1016/j.jallcom.2017.09.295 0925-8388/© 2017 Elsevier B.V. All rights reserved.
beneficial to obtain room-temperature multiferroic properties. Pal et al. demonstrated a large enhancement of ferromagnetism in modified BiFeO3-BaSrTiO3 ceramics [11]. Mandal et al. observed a switchable polarization and magnetization at RT in BiFeO3Bi(Ti0.5Mg0.5)O3-CaTiO3 solid solution [12]. These studies indicate that the BFO-based solid solutions are very attractive to realize multiferroic properties in view of their structure-property correlations. Recently, many studies show that the solid solutions of (1-x) BFO-x(Bi0.5K0.5)TiO3 (BKT) undergo a structural transition from R3c phase to P4mm phase with an increase in the content of BKT and exhibit a coexistence of R3c phase and P4mm phase when the amount of BKT is 40% [13-15]. A relatively large remnant polarization (Pr) of 52 mC/cm2 can be obtained in 0.6BFO-0.4BKT [13]. Moreover, the dielectric constant of BFO-BKT solid solutions show a broad maximum with frequency dispersion feature, which is one of characteristic of ralaxor ferroelectrics [16-18]. Hagiwara et al. reported that CuO-doped 0.6BFO-0.4BKT exhibits well-square hysteresis loops by means of thermal quenching and there is no ferromagnetism at room temperature [19]. In summary, previous studies mostly focused on the structural, ferroelectric and dielectric properties of 0.6BFO-0.4BKT solid solutions, whereas the magnetic properties of 0.6BFO-0.4BKT solid solutions have rarely been
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investigated the magnetic, ferroelectric and dielectric properties of 0.6BF1-xCxO-0.4BKT samples. The results show that the structural transition from pseudocubic to rhombohedral phase occurs in the x ¼ 0.1 sample and a coexistence of ferroelectricity and ferromagnetism at RT can be realized in the x ¼ 0.15 sample. The results indicate that it is an effective way to explore RT multiferroic compounds through a binary solid-solution. 2. Experimental procedures 2.1. Sample preparation
Fig. 1. XRD patterns of 0.6BF1-xCxO-0.4BKT (0 x 0.2) samples. The inset shows the amplified plots of the main diffraction (110) peaks.
revealed in literatures. In this work, we prepared 0.6BiFe1-xCoxO3-0.4(Bi0.5K0.5)TiO3 (0.6BF1-xCxO-0.4BKT) (0 x 0.2) by modified Pechini method at a relatively low temperature of 800 C. And we systematically
The 0.6BF1-xCxO-0.4BKT (0 x 0.2) ceramics were prepared through modified Pechini method, which could maintain atomiclevel mixing metal cations to lower sintering temperature for sake of minimization of the volatility of Bi and K. All the chemicals used in the experimental process were analytical grade and purchased from Sinopharm Chemical Reagent Co., Ltd. Stoichiometric amounts of the starting materials, Ti[OCH(CH3)2]4, KNO3, Bi(NO3)3$5H2O, Fe(NO3)3$9H2O, and Co(CH3COO)2$4H2O were dissolved into the citric acid monohydrate C6H8O7$H2O solution with glycol and acetic acid in stoichiometric proportions. Citric acid monohydrate C6H8O7$H2O was added to the solution as a complexant. Glycol was used for adjusting the viscosity and stability of the sol. Ammonia was added into the mixture of these solutions to adjust pH to 7 and the mixture was stirred until transparent precursor sol was obtained. Then the precursor solution was heated on a hot plate at 90 C resulting in the formation of a gel. The gel was heated at 400 C to remove the organic species. The obtained
Fig. 2. Rietveld refinement results for 0.6BF1-xCxO-0.4BKT (0 x 0.15) samples with (a) x ¼ 0 (b) x ¼ 0.05 (c) x ¼ 0.1, and (d) x ¼ 0.15. The insets are their corresponding amplified plots of (111) peak with 2q around 39 .
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powders were mixed and grounded, and then calcined at 700 C for 10 h under an O2 flow. The obtained powders were again pelletized, and sintered at 800 C for 20 h under an O2 flow with two intermediate grindings, and finally the furnace was cooled slowly to room temperature.
2.2. Characterizations The x-ray powder diffraction patterns were recorded at RT on a Panalytical diffractometer (X'Pert PROMRD) with Cu-Ka radiation and a graphite monochromator in a reflection mode. Structural refinements of 0.6BF1-xCxO-0.4BKT (0 x 0.15) are carried out by using RIETICA software. The morphological and structural information of all samples were characterized with a field emission scanning electron microscopy (FE-SEM, FEI-designed Sirion 200, Hillsboro, OR). The temperature dependence of magnetization was carried out from 300 to 800 K using a Physical Property Measurement System-vibrating sample magnetometer (PPMS-VSM) with a high-temperature oven option. The magnetization hysteresis measurements were performed with a Quantum Design superconducting quantum interface device (SQUID) magnetic property
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measurement system (MPMS) system (2 T 400 K, 0 H 4.5 T). Temperature dependence of dielectric constant and loss tangent ranging from 300 to 900 K were measured using an LCR meter in the frequency range of 1 kHze1000 kHz. The ferroelectric properties were investigated using a Sawyer-Tower circuit attached to a computer-controlled standardized ferroelectric test system (Radiant Technology 609B) with probes on the electrodes of the samples, which were immersed in silicone oil to prevent arcing.
3. Results and discussion Fig. 1 shows the XRD patterns of 0.6BF1-xCxO-0.4BKT (0 x 0.2) at RT. There is no impurity phase can be observed in the 0 x 0.15 samples, whereas small impurity peaks marked by stars at around 30 corresponding to Bi12(Bi0.5Fe0.5)O19.5 are observed in the x ¼ 0.2 sample. With increasing the doping level of Co, the main diffraction peak (110) shifts toward a higher angle position, implying the decrease in lattice parameters due to the smaller ionic radius of Co3þ (0.61 Å) compared with Fe3þ (0.645 Å), which is shown in the inset of Fig. 1. We refined the structural parameters of 0.6BF1-xCxO-0.4BKT (0 x 0.15) by using RIETICA
Fig. 3. Representative FE-SEM images of 0.6BF1-xCxO-0.4BKT samples with (a) x ¼ 0, (b) x ¼ 0.05, (c) x ¼ 0.1, (d) x ¼ 0.15, and (e) x ¼ 0.2.
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Fig. 4. XPS spectra and the deconvoluted results of (a)e(d) Fe 2 p of the x ¼ 0, 0.05, 0.1 and 0.15 samples, respectively, (e) Co 2 p of the x ¼ 0.05, 0.1 and 0.15 samples, and (f) Ti 2 p of the x ¼ 0, 0.05, 0.1 and 0.15 samples. The solid curves denote the fitting results.
software, as shown in Fig. 2(a)-(d). It can be seen that the fits between the experimental and calculated XRD patterns are relatively good based on the consideration of the small RP values of 3.37%, 3.96%, 3.32%, and 3.70% for the samples with x ¼ 0, 0.05, 0.1 and 0.15, respectively. For the x ¼ 0 and 0.05 samples, a single and no splitting (111) peak at about 39 is observed (inset of Fig. 2(a), (b)), and the (002) peak remains a single and sharp shape implying the existence of pseudocubic lattice. Whereas for the x ¼ 0.1 and 0.15 samples, the splitting of (111) is obviously observed implying that a rhombohedral lattice begins to appear at this composition, as shown in the inset of Fig. 2(c) and (d). Fig. 3 illustrates the field-emission scanning electron micrograph (FE-SEM) of 0.6BF1-xCxO-0.4BKT (0 x 0.2). Small grain size and rather loose morphology are observed in the x ¼ 0 and 0.05 samples. Whereas the x ¼ 0.1 and 0.15 samples exhibit relatively uniform and dense morphology without obvious porosity. The grain size of the x ¼ 0.1 and 0.15 samples are much larger than those of the x ¼ 0 and 0.05 samples. The particles of the x ¼ 0.2 sample have near-quadrate feature, which are much different from those of the x ¼ 0.1 and 0.15 samples. The different morphology of the x ¼ 0.2 sample might be due to the existence of impurity phase. The changes in morphology of 0.6BF1-xCxO-0.4BKT (0 x 0.15) ceramics indicate that Co doping can significantly lower the sintering temperature of the ceramics and promote the growth of the grains which in favor of enhancing the grain connectivity. To study the valence states of the ions of 0.6BF1-xCxO-0.4BKT (0 x 0.15) ceramics, we performed the XPS measurements for the 0 x 0.15 samples. The binding energy of the Fe 2p3/2 peaks locates at the range of 710.5e710.6 eV. To obtain the Fe3þ/Fe2þ ratio, the Fe 2p3/2 peaks were deconvoluted and fitted, as shown in Fig. 4(a)-(d). The content of Fe3þ are 36.7%, 39.3%, 50.1% and 56.8% for the x ¼ 0, 0.05, 0.1 and 0.15 samples, respectively. It is obvious that the content of Fe3þ increases gradually with the increase of the
Co concentration which indicates Co doping in the solid solution can effectively reduce the oxygen vacancies. The core-level spectra of Co 2p3/2 and the deconvoluted results for the x ¼ 0.05, 0.1 and 0.15 samples are shown in Fig. 4(e). It is found that peaks locates at 780.4, 780.4 and 780.8 eV for the x ¼ 0.05, 0.1 and 0.15 samples, respectively, which are very close to that of Co 2p3/2 (781.30 eV) in Co2O3 implying that the doped Co ions in the samples are in þ3 valence state [20]. The Ti 2p3/2 peaks located at 457.8e457.9 eV are consistent with the binding energy of Ti 2p3/2 (458.20 eV) in TiO2 [21], as shown in Fig. 4(f), which indicates Ti ions in our samples are in þ4 valence state. In order to make clear the magnetic behaviors of 0.6BF1-xCxO0.4BKT (0 x 0.15), we performed the magnetic hysteresis measurements at 300 K, as shown in Fig. 5. The M-H curve of the x ¼ 0 sample exhibits a linear behavior implying the characteristic of paramagnetic (PM) or AFM states. The M-H curve of the x ¼ 0.05 sample shows the same behavior as x ¼ 0 sample. With further increasing the content of Co, a small Mr value can be observed in the
Fig. 5. RT ferromagnetic hysteresis loops curves of the x ¼ 0, 0.05, 0.1 and 0.15 samples. The right inset is the corresponding amplified curves.
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Fig. 6. Temperature dependence of magnetization in the ZFC and FC modes measured at H ¼ 100 Oe for the samples with (a) x ¼ 0 and (b) x ¼ 0.15. The inset of Fig. 6(b) shows the plot of dM/dT.
x ¼ 0.1 sample as shown in the inset of Fig. 5. A large hysteresis loop with a remnant magnetization Mr of 0.107 emu/g and coercive field Hc of 300 Oe can be observed in the x ¼ 0.15 sample and these values are similar to those in the other solid solutions [22]. The observed remnant magnetization Mr in the x ¼ 0.15 sample is considered to be intrinsic property of the ceramic rather than from impurity phase because the XRD result has ruled out the possible of impurity phases. Moreover, the contribution of magnetism from the impurity Bi12(Bi0.95Fe0.05)O19.5 or Bi25FeO39 can also be
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excluded, due to the PM nature of the impurity phases [23,24]. The magnetization of the x ¼ 0.15 sample does not reach saturation when the magnetic field is up to H ¼ 40 kOe implying the coexistence of ferromagnetic (FM) and AFM interactions in the x ¼ 0.15 sample. A liner extrapolation of M-H to H ¼ 0 allows us to determine the spontaneous magnetization Ms of the x ¼ 0.15 sample as 0.0168 mB/f.u. This value is much smaller than the mean spin only value of isolated Fe3þ (2.5 mB), and Co3þ (2 mB) ions implying that the FM did not originate from the Fe or Co clusters. It is known that the bulk BiFeO3 possesses a spiral spin structure below the antiferromagnetic transition temperature TN ~643 K. Many previous studies manifest that the weak ferromagnetism with M-H hysteresis loops can be generated by decreasing the grain size below 62 nm or doping at Bi or Fe site due to the suppression of the spiral structure with the canting of the antiferromagnetic ordered spins. In this work, the ferromagnetism is proposed to originate from the suppression of spiral spin structure with the canting of the antiferromagnetically ordered spins caused by the substitution of Co for Fe ions [22,25,26]. From the XPS data, we can obtain the content of Fe3þ in the samples and therefore calculate the proportion of Co3þ to Fe3þ. It is found that the ratio of Co3þ to Fe3þ are 0.13, 0.22 and 0.31 for the samples with x ¼ 0.05, 0.1 and 0.15, respectively. This value of the x ¼ 0.15 sample is the largest among the Co-doped samples. In other words, the contribution of a larger amounts of Co3þ-Fe3þ antiferromagnetic ordered spins exists in the x ¼ 0.15 sample. The canting of these antiferromagnetic ordered spins is in favor of generation of the weak ferromagnetism and the x ¼ 0.15 sample exhibits larger remnant magnetization. To further understand the magnetic properties, we performed the measurements of temperature dependence of magnetization M (T) for the selected two samples with x ¼ 0 and 0.15 in the zero-field cooled (ZFC) and field cooled (FC) modes under a magnetic field of 100 Oe. As shown in Fig. 6(a), the magnetization of both ZFC and FC curves for the x ¼ 0 sample increase monotonously with the decrease in
Fig. 7. RT ferroelectric hysteresis loops of 0.6BF1-xCxO-0.4BKT (0 x 0.15) samples with (a) x ¼ 0, (b) x ¼ 0.05, (c) x ¼ 0.1, and (d) x ¼ 0.15.
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Fig. 8. Temperature dependence of dielectric constants for the samples with (a) x ¼ 0, (b) x ¼ 0.05, (c) x ¼ 0.1, and (d) x ¼ 0.15. The insets are their corresponding Vogel-Fulcher plots of relaxation frequencies vs. temperatures.
temperature and no long-range magnetic transition is observed implying the PM nature of the x ¼ 0 sample, which is consistent with the M-H curve. For the x ¼ 0.15 samples, the ZFC and FC curves almost overlap with each other in the high temperature range. As the temperature goes down, it seems to undergo a PM-to-FM magnetic transition at 683 K (defined as the temperature corresponding to the dip of dM/dT) as shown in the inset of Fig. 6(b). In the low temperature range below 660 K, there is a deviation between the ZFC and FC magnetization curves. Fig. 7 shows the polarization-electric field (P-E) hysteresis loops of 0.6BF1-xCxO-0.4BKT (0 x 0.15) measured at frequency of 100 Hz at RT. The P-E hysteresis loop of the x ¼ 0 sample exhibits a pinched loop with a remnant polarization Pr of 27 mC/cm2 in spite of a large applied-electric field of 160 kV/cm. The P-E loop of the x ¼ 0.05 sample is not well square enough. The poor P-E loops of the x ¼ 0 and 0.05 samples might be due to the loose morphology giving rise to high leakage current densities. However, well-defined P-E loops with Pr values of 47.5 mC/cm2 and 57.8 mC/cm2 are obtained for the x ¼ 0.1 and 0.15 samples, which is in the similar magnitude of other ferroelectric solid solution [13]. Normally, the large polarization is observed in epitaxy or exceptionally high quality single-crystal BiFeO3. However, it is rather difficult to obtain the large polarization in BiFeO3 bulk. In this work, the superior polarization in the Co-doped 0.6BFO-0.4BKT are mainly due to the formation of the solid solution of BKT, giving rise to the increase of grain size and the formation of dense morphology, which are in favor of increasing the remnant polarization [27]. Moreover, the introduction of Co3þ ions in 0.6BFO-0.4BKT can effectively lower the sintering temperature of the ceramics resulting in reduction of Bi volatilization and oxygen vacancies, which is in good agreement with the variation in the Fe3þ content from the XPS results. Figs. 8 and 9 show the temperature dependence of dielectric constant ε0 and dielectric loss tan d of 0.6BF1-xCxO-0.4BKT (0 x 0.15) at different frequencies from 300 to 900 K. For the x ¼ 0.1 sample, the dielectric constant peaks are broad, and the maximum in ε0 are diffusive accompanied by a strong dispersion of dielectric maximum temperature (Tm) with the variation of frequencies. With increasing the frequency, the dielectric maximum
decreases, while the Tm increases from 690 K at 10 kHz to 740 K at 1000 kHz showing a relaxation behavior which is similar to previous report [14]. With increasing the doping level of Co, the ε0 peaks become more obvious, and the maximum values of dielectric constant ε0 are much higher than those of the x ¼ 0 sample. As shown in Fig. 9, the values of tan d of the x ¼ 0.05, 0.1 and 0.15 samples is slightly smaller than that of the x ¼ 0 sample at RT, which are in agreement with the better polarization in Co doped samples. We note that the frequency-dependent peaks turn into sharp and the dielectric relaxor behavior become weaken with increasing the content of Co, which might be due to the structural transition from pseudocubic to rhombohedral phase associated with the domain switched from nanodomains to macrodomains [28], as confirmed by the XRD results. The observed relaxor behaviors in 0.6BF1-xCxO-0.4BKT (0 x 0.15) ceramics can be ascribed to chemical disorder resulting from mixed-valence cation-site occupation, which involved Bi3þ and Kþ at A-site, Ti4þ, Co3þ and Fe3þ at B-site of ABO3 perovskite lattice. These phenomena have also been observed in other ferroelectric compounds, such as (Bi0.5K0.5)TiO3 [29], (Ba0.6Bi0.2Li0.2)TiO3 [30]. The dielectric data were analyzed in terms of the Vogel-Fulcher relationship [31,32]: f ¼ f0 exp (Ea/kB (TmTvF)),
(1)
where f ¼ 1/t is the frequency, f0 is a pre-exponential factor (1/t0), Ea is the activation energy describing the polarization fluctuations of an isolated cluster relaxation process, Tm is the dielectric maximum in absolute temperature, T VF is the Vogel-Fulcher freezing temperature, and kB is the Boltzmann constant. The insets of Fig. 8 show the plot between ln f and T (K), where the dotted points represent experimental data and the solid lines are the fitting results according to Eq. (1). The values of T VF, f0, and Ea of x 0.15 samples are presented in Table 1. Ea values of the x 0.15 samples are similar to those of weakly-coupled relaxor ferroelectrics such as (Bi0.5K0.5)TiO3-BiScO3 [33] and BaTiO3-BiScO3 [34]. The existence of a magnetodielectric (MD) effect is often used as a test for ME coupling in multiferroic material. Therefore, we
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Fig. 9. Temperature dependence of tangent loss for the samples with (a) x ¼ 0, (b) x ¼ 0.05, (c) x ¼ 0.1, and (d) x ¼ 0.15.
Table 1 Vogel Fulcher fitting parameters for 0.6BF1-xCxO-0.4BKT (0 x 0.15) (Tm at 100 kHz). Nos.
Materials
1 2 3 4
x x x x
¼ ¼ ¼ ¼
0 0.05 0.1 0.15
TvF (K)
f0 (Hz)
Ea (eV)
Tm (K)
554 560 564 540
1.3 109 3.2 109 8.5 1010 15.5 1010
0.23 0.20 0.21 0.24
716 698 686 677
performed the measurement of the field-dependent dielectric constant ε0 at a frequency of 1 kHz at RT, as shown in Fig. 10. The dielectric constant gradually increases with increasing the applied magnetic fields implying the possible existence of MD effect. The MD effect can be defined as: MD ¼ [ε0 (H) - ε0 (0)] 100%/ε0 (0),
where ε0 (0) and ε0 (H) denote the dielectric constants under zero field and magnetic field, respectively. The magnitude of MD effect under 50 kOe at RT can be calculated as 0.03% for the x ¼ 0.15 sample. 4. Conclusions In conclusion, we have successfully prepared 0.6BF1-xCxO0.4BKT (0 x 0.2) ceramics. The Co doped samples undergo a structural transition from pseudocubic to rhombohedral phase at x ¼ 0.1. The FE-SEM results manifest that the substitution of Co3þ ions can significantly promote the grain growth resulting in dense morphology. The x ¼ 0.15 sample exhibits a maximum remnant magnetization of 0.107 emu/g and a superior remnant polarization of 57.8 mC/cm2. The ferromagnetism is proposed to originate from the suppression of spiral spin structure with the canting of the antiferromagnetically ordered spins caused by structural distortion due to the substitutions of Co for Fe ions. Dielectric studies show the feature of relaxor ferroelectrics resulting from mixed valence cation-site occupation. In addition, the MD effect (0.03% at RT) can be observed in the x ¼ 0.15 sample. Acknowledgments This work was supported by the National Key Research and Development Program of China (2017YFA0403502) and Key Research Program of Frontier Sciences, CAS (QYZDB-SSW-SLH015). References
Fig. 10. RT magnetic field dependence of MD effect at a frequency of 1 kHz for the x ¼ 0.15 sample. The black solid line arrows indicate the field sweeping direction.
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