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Effects of trivalent gadolinium and cobalt co-substitution on the crystal structure, electronic transport, and ferromagnetic properties of bismuth ferrite
Materials Science in Semiconductor Processing 27 (2014) 899–908
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Effects of trivalent gadolinium and cobalt co-substitution on the crystal structure, electronic transport, and ferromagnetic properties of bismuth ferrite G.L. Song, J. Su, G.J. Ma, T.X. Wang, H.G. Yang, F.G. Chang n College of Physics & Information Engineering, Henan Key Laboratory of Photovoltaic Materials, Henan Normal University, Xinxiang 453007, PR China
a r t i c l e i n f o
PACS: 75. 60. Ej 77. 80. e 77. 84. s Keywords: BiFeO3 Electronic transport properties Ferromagnetic properties The magnetic hysteresis loops The magnetic phase transition temperature
abstract Multiferroic trivalent gadolinium and cobalt co-substituted bismuth ferrite Bi0.95 Gd0.05Fe1 xCoxO3(x ¼0, 0.05, 0.1, 0.15, 0.2) ceramics were prepared by rapid liquid phase sintering method. The results showed that all the peaks of X-ray Diffraction (XRD) for Bi0.95Gd0.05Fe1 xCoxO3 samples can be indexed according to the crystal structure of pure BiFeO3. XRD analysis revealed a phase transition in Gd3 þ and Co3 þ co-doped BiFeO3 when x was larger than 0.1. The Scan Electron Microscope (SEM) images indicated that Gd3 þ and Co3 þ doping significantly decreased the grain sizes of BiFeO3 ceramic. Gd3 þ and Co3 þ co-doping BiFeO3 enhanced the electrical properties with lower leakage current. The magnetic hysteresis loops and the magnetization were greatly improved in co-substituted specimens at room temperature. The Mr of Bi0.95Gd0.05Fe1 xCoxO3 (x ¼ 0, 0.05, 0.1, 0.15, 0.2) was 34, 60, 105, 103 and 180 times of that of BiFeO3 at 30 kOe, respectively. All the samples exhibited ferromagnetic behavior at 750 K and paramagnetic behavior at 900 K, indicating a high temperature magnetic phase transition of BiFeO3 at 870 K, which shifted to 780 K by Gd3 þ and Co3 þ doping. This can be attributed to the Fe3 þ –O2 –Fe3 þ superexchange strength and the relative stability of the magnetic structure. & 2014 Elsevier Ltd. All rights reserved.
1. Introduction Multiferroic materials with simultaneous anti/ferromagnetic and ferroelectric orderings have been extensively studied for their potential applications in functional sensors and spintronic devices [1–3]. One of the most widely investigated multiferroics is BiFeO3 (BFO) due to its ferroelectric and magnetic transition temperatures located well above room temperature (RT), giving rise to possibilities of RT multiferroic devices [4]. BiFeO3 has a rhombohedral distorted perovskite structure in the space group R3c with
n
Corresponding author. E-mail address: [email protected] (F.G. Chang).
http://dx.doi.org/10.1016/j.mssp.2014.09.004 1369-8001/& 2014 Elsevier Ltd. All rights reserved.
a Neel temperature of 640 K and Curie temperature of 1100 K [2]. However, the high leakage current and weak macroscopic magnetism of BiFeO3 are the main barriers to its practical applications. It is essential in device applications to reduce leakage current and improve the magnetic behavior without disturbing the ferroelectric properties. One possible strategy is partial ionic substitution to obtain spontaneous magnetization in BiFeO3. Recently, great efforts have been focused on ferroelectric property, ferromagnetic property and magnetoelectric effect of the multiferroic material BiFeO3 [5–14]. Theoretically, Kornev et al. investigated high temperature magnetic behavior of BiFeO3 using Monte Carlo (MC) simulations methods. It is found that the aniferromagnetic phase transition (TN) of BiFeO3 is around 625–645 K and
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ferroelectric phase transition (TC) of BiFeO3 is around 1100 K [5]. Liu et al. investigated the magnetic structure, energy band structure and electronic structure of BiFeO3 and BiFe0.75Co0.25O3 by using density functional theory. The results showed that the ferromagnetism can be significantly improved by the doping of Co which changed the G-type antiferromagnetic order into ferromagnetism [6]. Experimentally, the magnetic properties of BiFeO3 were further modified by substitution of either A sites (Bi) with rare earth (Gd, Sm, Ho, Eu) [7–10] or B sites (Fe) with transition-metal (Co, Mn, Cr) [11–14] or A–B sites with rare earth and transition-metal at the same time, leading to improved ferroelectric and ferromagnetic properties [15–18]. Uniyal et al. reported that Gd3 þ doping in BiFeO3 suppressed secondary phases and improved the electrical and magnetic properties with Gd3 þ concentration from 0.05 to 0.1 [7]. Large magneto-electric effect was observed in Gd3 þ doped BiFeO3 at room temperature [19]. Naganuma et al. reported both the ferroelectric and magnetic properties of BiFeO3 films were enhanced by Co3 þ doping. The saturation magnetization was drastically enhanced by doping Co3 þ up to 12% [12]. The magnetic properties were greatly improved and the TN was obviously increased in Ba2 þ and Ca2 þ co-doped BiFeO3 powders [20]. Yang et al. observed saturation magnetization in BiFeO3 thin films could be enhanced by La–Co co-doping. However, they did not measure the temperature dependence of magnetic moment for Bi1 xLaxFe1 yCoyO3 specimen [16]. Up to now, there have been fewer reports on the influence of different element dopings, both in the A and B-sites of the BiFeO3 ceramics, on the high temperature magnetic characteristics and ferromagnetic phase transition temperature (TM). Thus, Gd3 þ and Co3 þ co-doped BiFeO3 was expected to have improved magnetic properties. In this work, polycrystalline Bi0.95Gd0.05Fe1 xCoxO3 ceramics were prepared by rapid liquid phase sintering method. Particularly, we focus on the structure and high temperature magnetic characteristics of Bi0.95Gd0.05Fe1 xCoxO3 compound, and for the first time the high temperature magnetic transition is studied.
characterized by X-ray Diffraction (XRD) using a diffractometer with Cu Kα1 radiation. The microstructure of Bi0.95Gd0.05Fe1 xCoxO3 was imaged with Scan Electron Microscope (SEM, CARIZEISS German) under operating voltage of 25 kV. The J–E curves of all samples were measured with Multiferroic Ferroelectric Test System. The magnetic properties of Bi0.95Gd0.05Fe1 xCoxO3 were measured with PPMS (Physical Property Measurement System, Versa Lab, Quantum Design).
3. Results and discussions Fig. 1 presents the x-ray diffraction (XRD) patterns (2θ from 201 to 801) of BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 (x¼0, 0.05, 0.1, 0.15, 0.2) ceramic samples. From XRD analysis (Fig. 1), we observe that the Gd3 þ and Co3 þ codoped BiFeO3 compounds are single-phase materials crystallizing in the same structure (rhombohedra with space group R3c) of the parent BiFeO3 compound.
2. Experimental details The BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 (x¼0, 0.05, 0.1, 0.15, 0.2) namely, BiFeO3 (BFO), Bi0.95Gd0.05FeO3 (BGFO), Bi0.95Gd0.05Fe0.95Co0.05O3 (BGFO–Co 5%), Bi0.95Gd0.05Fe0.1 Co0.1O3 (BGFO–Co 10%), Bi0.95Gd0.05Fe0.85Co0.15O3 (BGFO– Co 15%), Bi0.95Gd0.05Fe0.8Co0.2O3 (BGFO–Co 20%) ceramics were prepared by rapid liquid phase sintering method. The dried materials of Bi2O3, Fe2O3, Gd2O3 and Co2O3 (purityZ99.99%) were carefully weighed in stoichiometric proportion and thoroughly mixed in an agate mortar ballmilling machine for about 24 h (320 round/min) using high purity isopropyl alcohol as a medium. Then the mixture was dried and pressed into disks with diameter of about 13 mm and thickness 1 mm. Then, the disks were sintered at 870–880 1C for 450 s in air and subsequently quenched to room temperature naturally. All samples were carefully polished before the addiction of Ag electrodes on both surfaces to form metal insulator metal capacitors. The crystalline structure of the Bi0.95Gd0.05Fe1 xCoxO3 samples was
Fig. 1. X-ray diffraction (XRD) patterns of BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 samples (*: Bi2Fe4O9; a: XRD patterns of Bi0.95Gd0.05Fe1 xCoxO3; b: XRD patterns Bi0.95Gd0.05Fe1 xCoxO3 in the range of 2θ from 301 to 341).
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Table 1 The unit the crystal cell parameters (a, c, V), the remnant magnetization (Mr), the saturation magnetization(Ms), the Neel temperature (TN) and the magnetic phase transition temperature (TM) for BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 ceramics. Samples BFO BGFO BGFO-5% BGFO-10% BGFO-15% BGFO-20%
a ¼b (nm) 0.5601 0.5581 0.5572 0.5566 0.5559 0.5553
c (nm) 1.3921 1.3812 1.3771 1.3632 0.6883 0.6859
V (nm3) 0.378.26 0.372.35 0.36837 0.36517 0.18368 0.18352
It can be seen from Fig. 1 that impurity phase of Bi2Fe4O9 is present in BiFeO3 sample. The peaks with a 2θ 29 301 (marked with n) correspond to impurity phases of Bi2Fe4O9 [21]. It should be noted that the diffraction intensity of impurity phase Bi2Fe4O9 in Bi0.95Gd0.05Fe1 xCoxO3 are completely removed from the XRD pattern with Gd3 þ doping. The result is consistent with the literature, which reported that doping with rare earth La3 þ can eliminate the impurity phase Bi2Fe4O9 in BiFeO3 [7,9]. No Gd2O3 or Co2O3 diffraction peaks are detected in Bi0.95Gd0.05Fe1 xCoxO3 ceramics from Fig. 1, indicating Gd3 þ or Co3 þ ions have entered the BiFeO3 lattice. It disagreed with that reported by Yang et al. [16]. Yang et al. reported that the XRD of BLFO–Co 15% showed two additional peaks suggesting the possible existence of Bi25FeO40 [16]. All the peaks of Bi0.95Gd0.05Fe1 xCoxO3 (x ¼0, 0.05, 0.1, 0.15, 0.2) compounds agree well with the crystal structure of pure BiFeO3, exhibiting the rhombohedra lattice type with R3c space group. And also, the characteristic diffraction peaks of samples became gradually wider and shift right with Gd3 þ and Co3 þ doping. From these characters, we conclude as below: i). The samples prepared by rapid liquid phase sintering method are better than that prepared by solid state reaction method. The liquid phase accelerates the synthesizing reaction and probably prevents the formation of the second phase (for example: Bi2Fe4O9). Thus pure BiFeO3 phase is obtained by reducing the volatilization of Bi3 þ [21]. ii). Due to the substitution of smaller Gd3 þ (RGd3 þ ¼ 0.0938 nm and RBi3 þ ¼0.103 nm) and Co3 þ (RCo3 þ ¼ 0.0545 nm and RFe3 þ ¼0.0645 nm), the peaks shift slightly to larger θ values as the Gd3 þ and Co3 þ content increases. This shift in the diffraction angle might be ascribed to the unit cell contraction or the decrease in lattice constants.
In order to further analyze such transformation, the measured XRD patterns of Bi0.95Gd0.05Fe1 xCoxO3 samples were analyzed with Rietveld Refinement Program. The best fits to the measured data are observed using rhombohedral lattice type with R3c space group for x r0.10 samples (JCPDS20-10169) and with orthorhombic lattice type with Pnma space group for x Z0.15 samples (JCPDS14-0181) [10]. Details of the hkl index with refined unit cell parameters are shown in Table 1. It is noted from
Mr (emu/g)
Ms (emu/g)
TN (K)
TC (K)
0.1618 1.3237 2.2958 2.0698 3.6004
644 655 655 655 655 655
870 870 780 780 780 780
4
1 10 3.4 10 3 0.598 1.0451 1.0282 1.802
Table 1 that with the increasing concentration of Co3 þ in Bi0.95Gd0.05Fe1 xCoxO3, i.e. for x ¼0–0.10, the lattice constants (ar and cr) of Bi0.95Gd0.05Fe1 xCoxO3 are marginally reduced, as a result of unit cell volume contraction. Fig. 1(b) shows the enlarged version of the XRD patterns in the range of 2θ from 301 to 341. Fig. 1(b) reveals that the (104) and (110) peaks of BiFeO3 merged partially to formed a broadened peak (110) at xZ0.15. Such behavior indicates the propensity of the structure undergoing phase transformation from rhombohedral to orthorhombic structure with Co3 þ substitution [10,16,19]. Bi0.95Gd0.05Fe1 xCoxO3 (x¼0.15, 0.2) has orthorhombic symmetry with lattice parameters a¼0.5559 Å nm, c¼0.6883 nm for x¼0.15 and a¼0. 5553 nm, c¼0. 6859 nm for x¼0.2. The similar behavior has also been reported in Gd3 þ and Co3 þ doped BiFeO3 [10,16,19]. It can be inferred that the crystal structure of BiFeO3 sample doped with proper concentration of rare earth element doping can effectively modulate the structure of BiFeO3 [9,10,16]. The morphology of the surface for BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 specimens is shown in Fig. 2. The SEM image revealed microstructure, which comprised of grains of varying sizes with well-dented boundaries indicating polycrystalline nature of the material. The images indicated that Gd3 þ and Co3 þ doping significantly decreased the grain sizes of BiFeO3 ceramics. The grain size of unsubstituted BiFeO3 was larger than 8 μm with spherical pores inside the grains, which appeared to grow normally or continuously. Gd3 þ doping did not exhibit systematic influence on the grain size. The Gd3 þ substitution A site of Bi3 þ decreased the grain size and resulted in the relatively homogeneous microstructure of 2–3 μm sized grains at 5 mol% substitution (Fig. 2(b)). The decrease of grain size of Gd3 þ doped BiFeO3 can be interpreted by the suppression of oxygen vacancy concentration, which resulted in slower oxygen ion motion and lower grain growth rate [18]. When Gd3 þ substitution was constant, the effect of Co3 þ content on grain size was significant and Co3 þ substitution played a dominant role in decreasing the grain size of BiFeO3 by inhibiting the grain growth [22]. Thus, Gd3 þ and Co3 þ cosubstitution reduced the grain size by approximately 10 times compared to the un-substituted specimen, which is identified by the presence of diffraction peaks in Fig. 1. The current density of BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 ceramics was measured as a function of applied electric field (J–E) shown in Fig. 3. The current density curves of all the samples measured in this study are symmetrical. Therefore, the leakage behavior of all the samples is discussed only in a positively applied electric field. The current
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Fig. 2. SEM micrographs of BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 samples. (a) BiFeO3. (b) Bi0.95Gd0.05FeO3. (c) Bi0.95Gd0.05Fe0.95Co0.05O3. (d) Bi0.95Gd0.05Fe0.9 Co0.1O3. (e) Bi0.95Gd0.05Fe0.85Co0.15O3. (f) Bi0.95Gd0.05Fe0.8Co0.2O3.
densities of BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 (x¼0, 0.05, 0.1, 0.15, 0.2) were 3.65 10 3 A/cm2, 4.48 10 5 A/cm2, 1.05 10 4 A/cm2, 1.15 10 4 A/cm2, 9.47 10 5 A/cm2, 2.52 10 4 A/cm2 at electric field of 10 kV/cm, respectively. We can see that the current density decreased and then increased with doping of Gd3 þ and Co3 þ from Fig. 3(a). Pure BiFeO3 samples shows highest current density and Bi0.95 Gd0.05FeO3 samples shows lowest current density compared with other samples. It is obvious from Fig. 3(a) that even a small quantity of Gd3 þ doping can change the current density of BiFeO3 dramatically. In order to study further the effects of Gd3 þ and Co3 þ co-doping on the conduction mechanism of BiFeO3 samples, the current density (lg J) versus electric field (lg E) for Bi0.95Gd0.05Fe1 xCoxO3 ceramics at room temperature is shown in Fig. 3(b). It can be seen that lg J increases linearly with lg E below a critical voltage. Namely: lg J ¼ a lg E þ b
ð1Þ
Here a is the slope of linear fitting; b is logarithm of electric current density (J) when electric field (E) is equal to 1 V; Fig. 3(b) gives the least squares fitting of the corresponding data with fitting parameters listed in Table 2. For the samples of BiFeO3, Bi0.95Gd0.05Fe1 xCoxO3, Bi0.95Sm0.05Fe0.9Co0.1O3, the slope of the curve is 1.54, 1.38, 1.38, 1.22, 1.36 rather than 1, respectively. The slope of the curve of Bi0.95Gd0.05FeO3 is 1.08, which approximately equal to 1. It may be an indication that the charge density in sample is made of a small amount of free charges and space charges [16,18]. The former transfer in line with Ohm's law with a slope 1(J pEa a 1) [16], the latter transfer in line with the Mott–Gurney law (Jp Ea a 2) with a corresponding slope 2, in consistent with that
reported in references [16,18], which showed that the interaction of the two fitting slopes was between 1 and 2. For BiFeO3, It might be due to the appearance of charged defects governed by Fe2 þ ions, oxygen vacancies Vo and/or bismuth vacancies VBi in BiFeO3 sample [23–24]. According to the following reaction mechanisms [18,25–26]: 2Bi3 þ þ 3O2 -Bi2 O3 ↑ þ 2V3 Bi þ 3V2 þ O
ð2Þ
2Fe3 þ þ O2 -2Fe2 þ þ0:5O2 ↑ þ 2VO 2 þ 2þ Vo
ð3Þ V3 Bi
vacancy and as Importantly, the presence of stated in Eqs. (2) and (3) has the predominant effect on the reduction of the electrical resistivity of the bulk samples, giving rise to high leakage currents in the samples [23–24,26]. When Bi3 þ is partially replaced by Gd3 þ , which will suppress Bi3 þ volatile because the melting point of Gd2O3 (2330 1C) is much higher than Bi2O3 (820 1C). Thus, oxygen vacancies of Bi0.95Gd0.05FeO3 become smaller than that of BiFeO3. The best leakage characteristics of Bi0.95Gd0.05FeO3 samples are obtained. It is inferred that the origin of reduced leakage current in the Bi0.95Gd0.05FeO3 is a synergetic effect of “suppression of oxygen vacancies VO and/or bismuth vacancies VBi” leading to reduced Fe2 þ ions [17,29]. It shows that Gd3 þ doping BiFeO3 enhanced the electrical properties with lower leakage current density. The result is consistent with the results reported by Singh et al. [8–10,25,28,29]. With Co3 þ introduced once again, the current density of Bi0.95Gd0.05Fe1 xCoxO3 (x¼0.05–0.2) is slightly higher than that of Bi0.95Gd0.05FeO3 from Fig. 3(a). It can be seen from Fig. 3(b) that the slope of linear fitting of Bi0.95Gd0.05 Fe1 xCoxO3 (x¼0.05–0.2) is larger than that of Bi0.95Gd0.05 FeO3 and smaller than that of BiFeO3. It shows that space
G.L. Song et al. / Materials Science in Semiconductor Processing 27 (2014) 899–908
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Fig. 3. Current density (J) vs. applied electric field (E) characteristics for all samples measured at room temperature. (a) J–E plot of all samples. (b) lg J–lg E plots of all samples. Table 2 Fitting parameters a and b of electric current density (lg J) vs. voltage (lg E) date for all samples: a is the slope of fitting linear; b is logarithm of electric current density (J) when electric field (E) is equal to 1 v/cm; R2 is the linear regression coefficient. Samples
a
b
R2
BiFeO3 Bi0.95Gd0.05FeO3 Bi0.95Gd0.05Fe0.95Co0.0.5O3 Bi0.95Gd0.05Fe0.9Co0.1O3 Bi0.95Gd0.05Fe0.85Co0.15O3 Bi0.95Gd0.05Fe0.8Co0.2O3
1.54 1.08 1.38 1.38 1.22 1.36
7.21 7. 61 8.134 8.12 7.74 7.67
0.997 0.996 0.997 0.997 0.996 0.993
charge exists in Bi0.95Gd0.05Fe1 xCoxO3 (x¼0.05–0.2) samples. The theory of space charge formation can be found in Refs. [18,24]. The space charge takes part in long-term migration and formed space-charge-limited current resulting in the current density increases under the effect of an external electric field. Similar phenomena have been reported
Fig. 4. Magnetization vs. magnetic field curves for samples Bi0.95Gd0.05 Fe1 xCoxO3 at 50 K, 100 K and 300 K. (a) M–H for Bi0.95Gd0.05Fe1 xCoxO3 at 300 K. (b) M-H for Bi0.95Gd0.05Fe1 xCoxO3 at 100 K. (c) M–H for Bi0.95Gd0.05Fe1 xCoxO3 at 50 K.
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in the literature [16,18,27]. Panwar reported that Pr–Co codoped BiFeO3 thin films exhibited the lower leakage current than BiFeO3, due to Co ion substitution at Fe-site in BiFeO3 and the Fe ions concentration will be reduced, resulting in the decrease of the film conductivity [27]. The research works showed the conduction mechanism of Co doped and La–Co co-doped BiFeO3 thin films were found to be dominantly Ohmic in the low applied filed and Schottky emission in the high applied filed by Yang et al. [16]. Fig. 4 shows that the magnetic hysteresis loops for BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 ceramics with an applied magnetic filed up to 3 T at 300 K, 100 K and 50 K, respectively. All the samples have a complete magnetic hysteresis loops at room temperature as shown in Fig. 4(a). The saturated magnetic hysteresis loops of Bi0.95Gd0.05 Fe1 xCoxO3 are obtained with the introducing of Co3 þ . For pure BiFeO3, M–H curves exhibit roughly linear filed dependence at room temperature. But a nonlinear behavior of the M–H curve can be seen with enlarged scales as shown in the inset of Fig. 4(a). The remnant magnetization of pure BiFeO3 (Mr) is 1 10 4 emu/g, indicating that Pure BiFeO3 ceramic is weakly ferromagnetic. This result agreed very well with that reported by Du et al. [23–25]. Unexpectedly, the Gd-substituted specimen (Bi0.95Gd0.05FeO3) shows a linear dependence of magnetization on magnetic filed without hysteresis, and the curve is almost indistinguishable from that of pure BiFeO3. However, opening-up of the hysteresis loop can be seen in the inset of Fig. 4(a). The magnetization of Bi0.95Gd0.05FeO3 specimen is larger than that of pure BiFeO3 ceramic and also the Mr, Hc of Bi0.95Gd0.05FeO3 is 3.4 10 3 emu/g, 714.2 Oe, respectively. Clearly, doping with Gd3þ can improve ferromagnetic properties of BiFeO3 sample. The Gd3 þ and Co3 þ co-substituted BiFeO3 samples exhibit saturated and symmetric magnetic hysteresis loops, indicating a ferromagnetic behavior. This is consistent with zheng's results [15–16,28,29]. Zheng et al. prepared Bi0.9La0.1Fe0.95R0.05O3 (R¼Mn, Co) by sol–gel method with rapid sintering process and investigated magnetic properties of samples. They found that B-site doping of Mn or Co enhanced the magnetization of BiFeO3, but its Ms and Mr values were relatively small. Furthermore, the detailed changes in magnetization with Mn or Co content were not given [29]. When Co3 þ concentration is increased from 0.05 to 0.2, the remanent magnetization (Mr) increases from 0.33 to 1.01 emu/g and the saturation magnetization (Ms) increases from 1.05 to 0.267 emu/g. But the coercive field for all samples remains at 800 Oe. the Mr and Ms of the Bi0.95Gd0.05Fe1 xCoxO3 samples are re-plotted in Fig. 4(c) inset respectively as a function of Co3 þ concentrations x. The Ms, Mr, TN and TC parameters for BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 ceramics are summarized in Table 1. Fig. 4(b) inset shows the magnetization for Bi0.95Gd0.05 Fe1 xCoxO3 (x¼0.05, 0.1, 0.15, 0.2) samples as a function of temperature ranging from 50 to 400 K in a field of 10 kOe. For Bi0.95Gd0.05Fe1 xCoxO3, the magnetization gradually decreases with increasing temperature from 50 to 400 K under 10 kOe magnetic fields. It can be seen that Bi0.95Gd0.05 Fe1 xCoxO3 ceramics exhibit strong ferromagnetic properties, in agreement with the result of Kim [22,30].
Four possible explanations may account for the strong ferromagnetic properties of Bi0.95Gd0.05Fe1 xCoxO3 ceramics [31]: i). The oxygen vacancies existence. BiFeO3 ceramic exhibits weak ferromagnetic because oxygen vacancies of BiFeO3 sample lead to the presence of Fe2 þ . A lot of authoritative literatures have proved that the Fe2 þ exists in the BiFeO3 by XPS experiments [21,32,33]. Hu et al. reported the asymmetric broad peak at 710.0 eV can be observed for all the prepared samples by XPS indicating the coexistence of Fe3 þ and Fe2 þ [29]. According to the report of Jun et al. [13], the presence of Fe2 þ can lead to the enhancement of magnetization. Ionic radius of Fe2 þ is almost 15% larger than that of Fe3 þ , which can cause distortion in the crystal structure, resulting in the collinear spin arrangement in [111] planes and a non-zero net magnetic moment [13]. Fe3 þ –O2 –Fe2 þ magnetic exchange should have important impact on the magnetic properties of these samples. ii). The structure distortion. With Gd3 þ and Co3 þ co-doping, the unit cell contract or decrease in lattice constants of BiFeO3 due to the substitution of smaller Gd3 þ and Co3 þ than Bi3 þ and Fe3 þ . The reduction in unit cell volume and phase transition with the increase of Co3 þ is observed in Fig. 1 and Table 1. Since the superexchange interaction is sensitive to bond length and bond angle, the structure distortion can change Fe–O bond angle or suppress the spin spiral [30]. It releases potentially weak ferromagnetic order, which leads to enhancement of remnant magnetization in BiFeO3 system by doping of Co3 þ [34,35]. When xZ0.15, the Co3 þ substitution results in structural phase transition, wherein the spin cycloid might be destructed and homogeneous spin structure formed. It is observed that the latent magnetization locked within the cycloid might be released and the significant Mr value increased. Similar phenomena have been also reported in Co3 þ and Mn4 þ doped BiFeO3 [11,12]. iii). In Bi0.95Gd0.05Fe1 xCoxO3. The space modulated spin structure collapsed on Gd3 þ and Co3 þ co-doping. And the interaction between 4f electron of Gd3 þ and 3d electron of Fe3 þ or Co3 þ leads to parallel distribution of Fe3 þ spins when a magnetic field applied. Substitution of Co3 þ with smaller magnetic moment for Fe3 þ enhances the magnetic properties of BiFeO3 [22]. This can be ascribed to two aspects. Firstly, there is a local Fe3 þ –O2–Co3 þ magnetic coupling in the samples. This coupling provides ferromagnetic interaction between Fe3 þ and Co3 þ in macroscopic scale. But its microscopic mechanism is still to be explored. Huang et al. was reported similar results [35]. Secondly, the enhancement of magnetization may be attributed to the collapse of the spiral spin anti-ferromagnetic magnetic structure of BiFeO3 to form a new ferrous magnetic structure in BiFeO3 with Co3 þ substitute Fe3 þ [12].
Therefore, Gd3 þ and Co3 þ co-doping in BiFeO3 ceramic is proved to be an effective way to modulate the structure
G.L. Song et al. / Materials Science in Semiconductor Processing 27 (2014) 899–908
and magnetic properties, which are essential for practical applications. Fig. 5 shows the variation of the magnetic moment of BiFeO3 and B0.95Gd0.05Fe1 xCoxO3 samples with temperature from 300 to 900 K at an external magnetic field of 5 kOe. For BiFeO3 and B0.95Gd0.05Fe1 xCoxO3 samples, an antiferromagnetic transition is observed, with a ferromagnetic order below 640 K, 655 K (TN), and another magnetic transition is also observed above 870 K, 780 K (TM) respectively. From Fig. 5(a), the magnetic moment of BiFeO3 and Bi0.95Gd0.05FeO3 samples decreases with increasing temperature. The magnetic moment (M) of BiFeO3 changes abruptly at 644 K showing the anti-ferromagnetic phase transition of BiFeO3 sample, which is in agreement with the reported Neel temperature (TN ¼640 K) [1,2]. The magnetic moment of Bi0.95Gd0.05FeO3 shows transformation at 655 K (TN ¼655 K) increased slightly from 644 to 655 K when doped with Gd3 þ . The increase of TN of BiFeO3 from 644 to 655 K with Gd3 þ doping can be contributed to following two factors: Firstly, in BiFeO3, the magnetic order of anti-ferromagnetic arrangement between adjacent Fe3þ planes is known as G-type anti-ferromagnetic structure which is relatively stable [5]. Therefore, the TN, TM of BiFeO3 is equal to 644 K, 870 K, respectively. With the substitution of smaller Gd3þ for the larger Bi3þ , the structure of BiFeO3 will be distorted, resulting in unit cell contraction or decrease in lattice constants. The structure distortion can change Fe–O bond angle or spin order and enhance the Fe3þ –O2 –Fe3 þ double exchange [12]. Secondly, Gd3 þ has a large magnetic moment arising from its 4f electrons. This magnetic moment will interact with the 3d electrons of Fe3 þ (f–d exchange interaction) resulting in a distortion in the ant-ferromagnetic structure of the parent material. As a result, an enhanced ferromagnetic behavior and increased Neel temperature are observed [12,19]. The second magnetic phase transition happened at 870 K for BiFeO3 and Bi0.95Gd0.05FeO3, at 780 K for Bi0.95Gd0.05 Fe1 xCoxO3 showing a possible ferromagnetic or ferrimagnetic transition. The result is consistent with Das's report. Das et al. observed an antiferromagnetic transition around 370 1C at the Neel temperature (TN ¼640 K) and another magnetic transition above 600 1C for all Bi1 xBaxFeO3 samples [36]. The magnetic phase transition temperature is defined as TM (TM ¼ 870 K) in this paper, which is different from the Curie temperature of BiFeO3 (TC ¼1100 K). At present, there are few reports about the high temperature magnetic measurements of doped BiFeO3 and no one has mentioned the magnetic phase transition at 870 K. It could not be determined whether the magnetic transitions above TN for all samples were intrinsic or from impurities. If the transition were caused by very small amount of magnetic impurities which could not be detected by XRD method, the possible candidates for the impurities would most likely be Bi2O3, Fe2O3, Fe3O4, Co2O3, Gd2O3, Bi25FeO40 and Bi2Fe4O9. But Bi2O3 is non-magnetic and unlikely to cause such a magnetic transition. Impurities such as Bi25FeO40 and Bi2Fe4O9 are often present in BiFeO3 compound. Neutron diffraction measurements showed that Bi2Fe4O9 is paramagnetic at room temperature
905
Fig. 5. Magnetization as a function of temperature for all samples from 300 K to 900 K in 5000 Oe field. (a) BiFeO3 and Bi0.95Gd0.05FeO3. (b) Bi0.95Gd0.05Fe1 xCoxO3.
and undergoes a transition to an antiferromagnetic state at TN ¼(26473) K [37] and it can be effectively eliminated by doping with rare earth elements. Magnetic material Fe2O3 can exist in two forms: α-Fe2O3 and γ-Fe2O3. The former has a Curie temperature of 1013 K and the other has an anti-ferromagnetic transition at 952 K; neither of them can possibly lead to a transition at 780 K or 870 K. The Curie temperature of Co2O3 is 1388 K while BGFC Co is paramagnetic above 780 K. So Co2O3 is also unlikely responsible for the transition. Introducing Gd3 þ does not change the magnetic phase transition temperature (TM) of BiFeO3. In addition, Gd2O3 has a Tc of 292 K. Above that the material is paramagnetic. Therefore, Gd2O3 can be excluded from candidate list. If impure phase is trace of Fe3O4, the Curie temperature is 858 K. The value is well consistent with the measured magnetic phase transition temperatures of 870 K for BiFeO3 and BGFO samples, and the magnetic phase transition temperature (TM) of the BGFO3 sample will not
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change (TM ¼870 K) with doping Co3 þ in BGFO samples. The BGFO Co samples exhibited ferromagnetic behavior at 780 K. However, we found the phase transition temperature of BGFO–Co samples moving to lower temperature of 780 K from the experimental data in Fig. 5(b). It has been shown experimentally that the strength and position of the transition at 870 K for BiFeO3 are not affected by these impurities. To sum up, there is no evidence indicating that the high temperature magnetic transition in BFO series arises from impurities. The relatively high magnetic transition temperature (TM ¼870 K) of BiFeO3 remains the same with Gd3 þ doping. Although the physical origin of this ferromagnetic/frrimagnetic transition is not clear at the moment. It seems likely that the interaction between the Fe3 þ ions in the (111) planes is stronger and reaches ferromagnetic order before the Fe3 þ –O2-–Fe3 þ antimagnetic order as the temperature decreases [13]. Substitution of Gd3 þ for Bi3 þ has little effects on the ferromagnetic order of Fe3 þ and the TM of BiFeO3. Another possibility may be that the Fe3 þ –O2–Fe2 þ interaction is stronger than the Fe3 þ –O2 –Fe3 þ interaction [9,16,20,23]. Therefore, BiFeO3 gives a higher ferrimagnetic phase transition temperature of 870 K. When Co3 þ was further introduced and increased from 0.05 to 0.2, the Gd3 þ and Co3 þ co-substituted specimens exhibited a strong ferromagnetic behavior. The magnetic moments of Bi0.95Gd0.05Fe1 xCoxO3 decreased with increasing temperature. We found that the magnetic moments of Bi0.95Gd0.05Fe1 xCoxO3 (x¼0.05, 0.1) also changed abruptly at 655 K, indicating the anti-ferromagnetic phase transition.
The magnetic moments of Bi0.95Gd0.05Fe1 xCoxO3 changed from ferromagnetic to paramagnetic at 780 K. The magnetic phase transition temperature (TM) moving to lower temperature from 870 K to 780 K demonstrates that Co3þ doping reduces the magnetic phase transition temperature (TM) compared with pure BiFeO3 samples. This enhancement in ferrromagnetic ordering and decrease of TM is not due to the impurity of Co2O3 and Fe3O4 which have rather high Curie temperature of TC ¼1388 K, 857 K, respectively. The phase transition cannot occur at 780 K if Co2O3 and Fe3O4 exist in the samples, which are in consistent with the result given in Fig. 1. Three possible explanations may account for the magnetic phase transition temperature of Bi0.95Gd0.05Fe1 xCoxO3 ceramics moving to lower temperature from 870 K to 780 K. i). As the electronic configuration of Co3þ is exactly the same as that of Fe3þ , Fe3þ –O2 –Co3 þ coupling would enhance the ferrimagnetic behavior of the material just as the Fe3 þ –O2–Fe2þ coupling. Naturally, substitution of Co3þ for Fe3 þ will increase the number of Fe3þ –O2 –Co3 þ coupling leading to much stronger magnetization. The bond angle of Fe3þ –O2 –Fe3 þ might become larger and the bond length of Fe3þ –O2 –Co3þ is greater than that of Fe3þ –O2 –Fe3 þ [22,28], resulting in a relatively weak Fe–O–Co superexchange interaction with co-doping of Gd3 þ and Co3þ . It can be a major reason that weak Fe–O–Co superexchange interaction and instable ferrous magnetic structure of Bi0.95Gd0.05Fe1 xCoxO3 lead to the magnetic
Fig. 6. Magnetization vs. magnetic field curves for samples Bi0.95Gd0.05Fe1 x CoxO3 at 700 K, 750 K, 900 K, respectively. (a) BiFeO3. (b) Bi0.95Gd0.05FeO3. (c) Bi0.95Gd0.05Fe0.95Co0.05O3. (d) Bi0.95Gd0.05Fe0.9Co0.1O3. (e) Bi0.95Gd0.05Fe0.85Co0.15O3. (f) Bi0.95Gd0.05Fe0.8Co0.2O3.
G.L. Song et al. / Materials Science in Semiconductor Processing 27 (2014) 899–908
transition temperature shifting toward lower temperature at 780 K. The enhancement of magnetization may also be attributed to the collapse of the spiral spin anti-ferromagnetic magnetic structure to form a new ferrous magnetic structure in BiFeO3 in spite of the existence of magnetic coupling between the local Fe–O–Co with Co3þ substitute Fe3þ . Since the super-exchange interaction is sensitive to bond length and bond angle, the spiral spin structure might be destroyed [30]. ii). With Co3 þ doped, the main diffraction peaks between (104) and (110) of BiFeO3 sample are overlapped (110) to single peak gradually, rhombohedra symmetry of BiFeO3 has transformed to tetragonal perovskite structure from Fig. 1. The structure phase transition was one of the reasons that could lead to the magnetic transition temperature (TM) of BiFeO3 shift from 870 K to 780 K. iii). 4f of Gd3 þ and 3d of Fe3 þ or Co3 þ form a weak exchange interaction (f–d exchange interaction) in Bi0.95Gd0.05Fe1 xCoxO3. But the f–d exchange interaction is relatively weak compared with d–d, a difference of one order of magnitude [38]. Therefore, the f–d exchange interaction does not affect the magnetic transition temperature (TM) of BiFeO3. Considering the three reasons, we may draw a conclusion that the change in magnetic transition temperature (TM) of Bi0.95Gd0.05Fe1 xCoxO3 depends mainly on the Fe3 þ –O2 –Fe3 þ super-exchange strength and the relative stability of the magnetic structure. In order to verify the fact that the ferromagnetic phase transition of BiFeO3 and Bi0.95Gd0.05Fe1 x CoxO3 samples occurred at 870 K and 780 K. We measured the magnetic hysteresis loops for all samples with an applied field up to 3 T at 700 K, 750 K and 900 K, respectively, as shown in Fig. 6. All the samples exhibit symmetric magnetic hysteresis loops at 700 K and 750 K, indicating a ferromagnetic behavior. However, all the samples exhibit linear magnetic hysteresis loops at 900 K, indicating the paramagnetic behavior. This is in consistent with Das's results [36]. Das reported that a large change in the magnetization was observed around 370 1C, which was close to the Neel temperature (TN) of Ba2 þ doped BiFeO3, and another magnetic transition which exhibited the paramagnetism was observed upon 600 1C [36]. It evidences that the ferromagnetic phase transition of Bi0.95Gd0.05Fe1 xCoxO3 samples occur at 780 K in Fig. 5 and no Co2O3 impurity exist in all the samples. Here for the first time the transition temperature has been determined. Naturally, more experimental and theoretical works are needed to clarify the exact origin of this high temperature ferrimagnetic transition. To sum up, we think that Gd3 þ and Co3 þ co-doping in BiFeO3 ceramics is proved to be an effective way to modulate the structure and improve magnetic properties, which is essential for practical applications. 4. Conclusion In conclusion, we have studied the effect of Gd3 þ and Co3 þ doping in the BiFeO3 compound on its crystalline
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structure, electronic transport, ferromagnetism properties and the high temperature magnetic phase transition. The following conclusions have been reached: The main diffraction peak between (104) and (110) of BiFeO3 samples gradually overlapped (110) to single peak with increasing x, implying that rhombohedral symmetry has transformed to orthorhombic perovskite structure. The SEM images indicated that Gd3 þ and Co3 þ doping significantly decreased the grain sizes of BiFeO3 ceramics. Gd3 þ and Co3 þ co-substitution reduced the grain size by approximately 10 times compared to the un-substituted specimen. Gd3 þ and Co3 þ co-doping BiFeO3 enhanced the electrical properties with lower leakage current. The current densities of Bi0.95Gd0.05FeO3 measured at 10 kV are approximately two orders of magnitude lower than that of pure BiFeO3. All samples possess strong ferromagnetism at room temperature except BiFeO3 and Bi0.95Gd0.05FeO3, which show rather weak ferromagnetism. The TN of Bi0.95Gd0.05FeO3 increases from 644 to 655 K, which can be ascribed to the crystal structure change and the exchange interaction between 4f electrons of Gd3 þ and 3d electrons of Fe3 þ . The change in TM of Bi0.95Gd0.05Fe1 xCoxO3 depends mainly on the Fe–O–Fe super-exchange strength and the relative stability of the magnetic structure. All the samples exhibit ferromagnetic behavior below 750 K and paramagnetic behavior above 900 K. It is evidenced that the ferromagnetic phase transition of Bi0.95Gd0.05Fe1 xCoxO3 samples occur at 780 K.
Acknowledgments This work was supported by the National Natural Science Foundation of China (Project no: U120411); Basic and Advanced Technology Research Projects in Henan Province, China (Project no: 122300410203, 122300413208); and Basic Research Program of Education Bureau of Henan Province, China (2011A140014). References [1] W. Eerenstein, N.D. Mathur, J.F. Scott, Nature 442 (2006) 759–765. [2] C.T. Nelson, P. Gao, J.R. Jokisaari, C.M. Folkman, C.B. Eom, D.G. Schlom, Science 334 (2011) 968–971. [3] L.H. Yin, W.H. Song, X.L. Jiao, W.B. Wu, X.B. Zhu, Z.R. Yang, J.M. Dai, R.L. Zhang, Y.P. Sun, J. Phys. D: Appl. Phys. 42 (2009) 205402. [4] Z. Wen, X. Shen, J.X. Wu, D. Wu, A.D. Li., J.L. Wang, Appl Phys Lett. 96 (2010) 202904-1–202904-4. [5] Igor A. Kornev, S. Lisenkov, R. Haumont, B. Dkhil, L. Bellaiche, Phys. Rev. Lett. 99 (2007) 227602-1–227602-4. [6] H. Zhang, Y.J. Liu, L.H. Pan, Acta Phys. Sin. 58 (2009) 7141. [7] P. Uniyal, K.L. Yadav, Mater. Lett. 62 (2008) 2858. (4). [8] K.S. Nalwa, A. Garg, A. Upadhyaya, Mater. Lett. 62 (2008) 878. (4). [9] G.L. Song, G.J. Ma, J. Su, T.X. Wang, H.Y. Yang, F.G. Chang, Ceram. Int. 40 (2014) 3579–3587. [10] X.Q. Zhang, Y. Sui, X.J. Wang, Y. Wang, Z. Wang, J. Alloys Compd. 507 (2010) 157–161. [11] E.M. Choi, S. Patnaik, E. Weal, S.L. Sahonta, H. Wang, Z. Bi, J. Xiong, M.G. Blamire, Q.X. Jia, J.L. Mac Manus Driscoll, Appl. Phys. Lett. 98 (2011) 012509–012509-3. [12] H. Naganum, N. Shimur, J. Miur, H. Shim, S. Yasui, K. Nishida, S. Okamur, J. Appl. Phys. 103 (2008) 07E314. [13] Y.K. Jun, S.H. Hong, Solid State Commun. 144 (2007) 329–333. [14] F.G. Chang, N. Zhang, G.L. Song, J. Phys. D: Appl. Phys. 40 (2007) 7799. [15] F.Z. Qian, J.S. Jiang, D.M. Jiang, C.M. Wang, W.G. Zhang, J. Magn. Magn. Mater. 322 (2010) 3127–3130. [16] K.G. Yang, Y.L. Zhang, S.H. Yang, B. Wang, J. Appl. Phys. 107 (2010) 124109.
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Materials Science in Semiconductor Processing journal homepage: www.elsevier.com/locate/mssp
Effects of trivalent gadolinium and cobalt co-substitution on the crystal structure, electronic transport, and ferromagnetic properties of bismuth ferrite G.L. Song, J. Su, G.J. Ma, T.X. Wang, H.G. Yang, F.G. Chang n College of Physics & Information Engineering, Henan Key Laboratory of Photovoltaic Materials, Henan Normal University, Xinxiang 453007, PR China
a r t i c l e i n f o
PACS: 75. 60. Ej 77. 80. e 77. 84. s Keywords: BiFeO3 Electronic transport properties Ferromagnetic properties The magnetic hysteresis loops The magnetic phase transition temperature
abstract Multiferroic trivalent gadolinium and cobalt co-substituted bismuth ferrite Bi0.95 Gd0.05Fe1 xCoxO3(x ¼0, 0.05, 0.1, 0.15, 0.2) ceramics were prepared by rapid liquid phase sintering method. The results showed that all the peaks of X-ray Diffraction (XRD) for Bi0.95Gd0.05Fe1 xCoxO3 samples can be indexed according to the crystal structure of pure BiFeO3. XRD analysis revealed a phase transition in Gd3 þ and Co3 þ co-doped BiFeO3 when x was larger than 0.1. The Scan Electron Microscope (SEM) images indicated that Gd3 þ and Co3 þ doping significantly decreased the grain sizes of BiFeO3 ceramic. Gd3 þ and Co3 þ co-doping BiFeO3 enhanced the electrical properties with lower leakage current. The magnetic hysteresis loops and the magnetization were greatly improved in co-substituted specimens at room temperature. The Mr of Bi0.95Gd0.05Fe1 xCoxO3 (x ¼ 0, 0.05, 0.1, 0.15, 0.2) was 34, 60, 105, 103 and 180 times of that of BiFeO3 at 30 kOe, respectively. All the samples exhibited ferromagnetic behavior at 750 K and paramagnetic behavior at 900 K, indicating a high temperature magnetic phase transition of BiFeO3 at 870 K, which shifted to 780 K by Gd3 þ and Co3 þ doping. This can be attributed to the Fe3 þ –O2 –Fe3 þ superexchange strength and the relative stability of the magnetic structure. & 2014 Elsevier Ltd. All rights reserved.
1. Introduction Multiferroic materials with simultaneous anti/ferromagnetic and ferroelectric orderings have been extensively studied for their potential applications in functional sensors and spintronic devices [1–3]. One of the most widely investigated multiferroics is BiFeO3 (BFO) due to its ferroelectric and magnetic transition temperatures located well above room temperature (RT), giving rise to possibilities of RT multiferroic devices [4]. BiFeO3 has a rhombohedral distorted perovskite structure in the space group R3c with
n
Corresponding author. E-mail address: [email protected] (F.G. Chang).
http://dx.doi.org/10.1016/j.mssp.2014.09.004 1369-8001/& 2014 Elsevier Ltd. All rights reserved.
a Neel temperature of 640 K and Curie temperature of 1100 K [2]. However, the high leakage current and weak macroscopic magnetism of BiFeO3 are the main barriers to its practical applications. It is essential in device applications to reduce leakage current and improve the magnetic behavior without disturbing the ferroelectric properties. One possible strategy is partial ionic substitution to obtain spontaneous magnetization in BiFeO3. Recently, great efforts have been focused on ferroelectric property, ferromagnetic property and magnetoelectric effect of the multiferroic material BiFeO3 [5–14]. Theoretically, Kornev et al. investigated high temperature magnetic behavior of BiFeO3 using Monte Carlo (MC) simulations methods. It is found that the aniferromagnetic phase transition (TN) of BiFeO3 is around 625–645 K and
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G.L. Song et al. / Materials Science in Semiconductor Processing 27 (2014) 899–908
ferroelectric phase transition (TC) of BiFeO3 is around 1100 K [5]. Liu et al. investigated the magnetic structure, energy band structure and electronic structure of BiFeO3 and BiFe0.75Co0.25O3 by using density functional theory. The results showed that the ferromagnetism can be significantly improved by the doping of Co which changed the G-type antiferromagnetic order into ferromagnetism [6]. Experimentally, the magnetic properties of BiFeO3 were further modified by substitution of either A sites (Bi) with rare earth (Gd, Sm, Ho, Eu) [7–10] or B sites (Fe) with transition-metal (Co, Mn, Cr) [11–14] or A–B sites with rare earth and transition-metal at the same time, leading to improved ferroelectric and ferromagnetic properties [15–18]. Uniyal et al. reported that Gd3 þ doping in BiFeO3 suppressed secondary phases and improved the electrical and magnetic properties with Gd3 þ concentration from 0.05 to 0.1 [7]. Large magneto-electric effect was observed in Gd3 þ doped BiFeO3 at room temperature [19]. Naganuma et al. reported both the ferroelectric and magnetic properties of BiFeO3 films were enhanced by Co3 þ doping. The saturation magnetization was drastically enhanced by doping Co3 þ up to 12% [12]. The magnetic properties were greatly improved and the TN was obviously increased in Ba2 þ and Ca2 þ co-doped BiFeO3 powders [20]. Yang et al. observed saturation magnetization in BiFeO3 thin films could be enhanced by La–Co co-doping. However, they did not measure the temperature dependence of magnetic moment for Bi1 xLaxFe1 yCoyO3 specimen [16]. Up to now, there have been fewer reports on the influence of different element dopings, both in the A and B-sites of the BiFeO3 ceramics, on the high temperature magnetic characteristics and ferromagnetic phase transition temperature (TM). Thus, Gd3 þ and Co3 þ co-doped BiFeO3 was expected to have improved magnetic properties. In this work, polycrystalline Bi0.95Gd0.05Fe1 xCoxO3 ceramics were prepared by rapid liquid phase sintering method. Particularly, we focus on the structure and high temperature magnetic characteristics of Bi0.95Gd0.05Fe1 xCoxO3 compound, and for the first time the high temperature magnetic transition is studied.
characterized by X-ray Diffraction (XRD) using a diffractometer with Cu Kα1 radiation. The microstructure of Bi0.95Gd0.05Fe1 xCoxO3 was imaged with Scan Electron Microscope (SEM, CARIZEISS German) under operating voltage of 25 kV. The J–E curves of all samples were measured with Multiferroic Ferroelectric Test System. The magnetic properties of Bi0.95Gd0.05Fe1 xCoxO3 were measured with PPMS (Physical Property Measurement System, Versa Lab, Quantum Design).
3. Results and discussions Fig. 1 presents the x-ray diffraction (XRD) patterns (2θ from 201 to 801) of BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 (x¼0, 0.05, 0.1, 0.15, 0.2) ceramic samples. From XRD analysis (Fig. 1), we observe that the Gd3 þ and Co3 þ codoped BiFeO3 compounds are single-phase materials crystallizing in the same structure (rhombohedra with space group R3c) of the parent BiFeO3 compound.
2. Experimental details The BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 (x¼0, 0.05, 0.1, 0.15, 0.2) namely, BiFeO3 (BFO), Bi0.95Gd0.05FeO3 (BGFO), Bi0.95Gd0.05Fe0.95Co0.05O3 (BGFO–Co 5%), Bi0.95Gd0.05Fe0.1 Co0.1O3 (BGFO–Co 10%), Bi0.95Gd0.05Fe0.85Co0.15O3 (BGFO– Co 15%), Bi0.95Gd0.05Fe0.8Co0.2O3 (BGFO–Co 20%) ceramics were prepared by rapid liquid phase sintering method. The dried materials of Bi2O3, Fe2O3, Gd2O3 and Co2O3 (purityZ99.99%) were carefully weighed in stoichiometric proportion and thoroughly mixed in an agate mortar ballmilling machine for about 24 h (320 round/min) using high purity isopropyl alcohol as a medium. Then the mixture was dried and pressed into disks with diameter of about 13 mm and thickness 1 mm. Then, the disks were sintered at 870–880 1C for 450 s in air and subsequently quenched to room temperature naturally. All samples were carefully polished before the addiction of Ag electrodes on both surfaces to form metal insulator metal capacitors. The crystalline structure of the Bi0.95Gd0.05Fe1 xCoxO3 samples was
Fig. 1. X-ray diffraction (XRD) patterns of BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 samples (*: Bi2Fe4O9; a: XRD patterns of Bi0.95Gd0.05Fe1 xCoxO3; b: XRD patterns Bi0.95Gd0.05Fe1 xCoxO3 in the range of 2θ from 301 to 341).
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Table 1 The unit the crystal cell parameters (a, c, V), the remnant magnetization (Mr), the saturation magnetization(Ms), the Neel temperature (TN) and the magnetic phase transition temperature (TM) for BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 ceramics. Samples BFO BGFO BGFO-5% BGFO-10% BGFO-15% BGFO-20%
a ¼b (nm) 0.5601 0.5581 0.5572 0.5566 0.5559 0.5553
c (nm) 1.3921 1.3812 1.3771 1.3632 0.6883 0.6859
V (nm3) 0.378.26 0.372.35 0.36837 0.36517 0.18368 0.18352
It can be seen from Fig. 1 that impurity phase of Bi2Fe4O9 is present in BiFeO3 sample. The peaks with a 2θ 29 301 (marked with n) correspond to impurity phases of Bi2Fe4O9 [21]. It should be noted that the diffraction intensity of impurity phase Bi2Fe4O9 in Bi0.95Gd0.05Fe1 xCoxO3 are completely removed from the XRD pattern with Gd3 þ doping. The result is consistent with the literature, which reported that doping with rare earth La3 þ can eliminate the impurity phase Bi2Fe4O9 in BiFeO3 [7,9]. No Gd2O3 or Co2O3 diffraction peaks are detected in Bi0.95Gd0.05Fe1 xCoxO3 ceramics from Fig. 1, indicating Gd3 þ or Co3 þ ions have entered the BiFeO3 lattice. It disagreed with that reported by Yang et al. [16]. Yang et al. reported that the XRD of BLFO–Co 15% showed two additional peaks suggesting the possible existence of Bi25FeO40 [16]. All the peaks of Bi0.95Gd0.05Fe1 xCoxO3 (x ¼0, 0.05, 0.1, 0.15, 0.2) compounds agree well with the crystal structure of pure BiFeO3, exhibiting the rhombohedra lattice type with R3c space group. And also, the characteristic diffraction peaks of samples became gradually wider and shift right with Gd3 þ and Co3 þ doping. From these characters, we conclude as below: i). The samples prepared by rapid liquid phase sintering method are better than that prepared by solid state reaction method. The liquid phase accelerates the synthesizing reaction and probably prevents the formation of the second phase (for example: Bi2Fe4O9). Thus pure BiFeO3 phase is obtained by reducing the volatilization of Bi3 þ [21]. ii). Due to the substitution of smaller Gd3 þ (RGd3 þ ¼ 0.0938 nm and RBi3 þ ¼0.103 nm) and Co3 þ (RCo3 þ ¼ 0.0545 nm and RFe3 þ ¼0.0645 nm), the peaks shift slightly to larger θ values as the Gd3 þ and Co3 þ content increases. This shift in the diffraction angle might be ascribed to the unit cell contraction or the decrease in lattice constants.
In order to further analyze such transformation, the measured XRD patterns of Bi0.95Gd0.05Fe1 xCoxO3 samples were analyzed with Rietveld Refinement Program. The best fits to the measured data are observed using rhombohedral lattice type with R3c space group for x r0.10 samples (JCPDS20-10169) and with orthorhombic lattice type with Pnma space group for x Z0.15 samples (JCPDS14-0181) [10]. Details of the hkl index with refined unit cell parameters are shown in Table 1. It is noted from
Mr (emu/g)
Ms (emu/g)
TN (K)
TC (K)
0.1618 1.3237 2.2958 2.0698 3.6004
644 655 655 655 655 655
870 870 780 780 780 780
4
1 10 3.4 10 3 0.598 1.0451 1.0282 1.802
Table 1 that with the increasing concentration of Co3 þ in Bi0.95Gd0.05Fe1 xCoxO3, i.e. for x ¼0–0.10, the lattice constants (ar and cr) of Bi0.95Gd0.05Fe1 xCoxO3 are marginally reduced, as a result of unit cell volume contraction. Fig. 1(b) shows the enlarged version of the XRD patterns in the range of 2θ from 301 to 341. Fig. 1(b) reveals that the (104) and (110) peaks of BiFeO3 merged partially to formed a broadened peak (110) at xZ0.15. Such behavior indicates the propensity of the structure undergoing phase transformation from rhombohedral to orthorhombic structure with Co3 þ substitution [10,16,19]. Bi0.95Gd0.05Fe1 xCoxO3 (x¼0.15, 0.2) has orthorhombic symmetry with lattice parameters a¼0.5559 Å nm, c¼0.6883 nm for x¼0.15 and a¼0. 5553 nm, c¼0. 6859 nm for x¼0.2. The similar behavior has also been reported in Gd3 þ and Co3 þ doped BiFeO3 [10,16,19]. It can be inferred that the crystal structure of BiFeO3 sample doped with proper concentration of rare earth element doping can effectively modulate the structure of BiFeO3 [9,10,16]. The morphology of the surface for BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 specimens is shown in Fig. 2. The SEM image revealed microstructure, which comprised of grains of varying sizes with well-dented boundaries indicating polycrystalline nature of the material. The images indicated that Gd3 þ and Co3 þ doping significantly decreased the grain sizes of BiFeO3 ceramics. The grain size of unsubstituted BiFeO3 was larger than 8 μm with spherical pores inside the grains, which appeared to grow normally or continuously. Gd3 þ doping did not exhibit systematic influence on the grain size. The Gd3 þ substitution A site of Bi3 þ decreased the grain size and resulted in the relatively homogeneous microstructure of 2–3 μm sized grains at 5 mol% substitution (Fig. 2(b)). The decrease of grain size of Gd3 þ doped BiFeO3 can be interpreted by the suppression of oxygen vacancy concentration, which resulted in slower oxygen ion motion and lower grain growth rate [18]. When Gd3 þ substitution was constant, the effect of Co3 þ content on grain size was significant and Co3 þ substitution played a dominant role in decreasing the grain size of BiFeO3 by inhibiting the grain growth [22]. Thus, Gd3 þ and Co3 þ cosubstitution reduced the grain size by approximately 10 times compared to the un-substituted specimen, which is identified by the presence of diffraction peaks in Fig. 1. The current density of BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 ceramics was measured as a function of applied electric field (J–E) shown in Fig. 3. The current density curves of all the samples measured in this study are symmetrical. Therefore, the leakage behavior of all the samples is discussed only in a positively applied electric field. The current
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Fig. 2. SEM micrographs of BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 samples. (a) BiFeO3. (b) Bi0.95Gd0.05FeO3. (c) Bi0.95Gd0.05Fe0.95Co0.05O3. (d) Bi0.95Gd0.05Fe0.9 Co0.1O3. (e) Bi0.95Gd0.05Fe0.85Co0.15O3. (f) Bi0.95Gd0.05Fe0.8Co0.2O3.
densities of BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 (x¼0, 0.05, 0.1, 0.15, 0.2) were 3.65 10 3 A/cm2, 4.48 10 5 A/cm2, 1.05 10 4 A/cm2, 1.15 10 4 A/cm2, 9.47 10 5 A/cm2, 2.52 10 4 A/cm2 at electric field of 10 kV/cm, respectively. We can see that the current density decreased and then increased with doping of Gd3 þ and Co3 þ from Fig. 3(a). Pure BiFeO3 samples shows highest current density and Bi0.95 Gd0.05FeO3 samples shows lowest current density compared with other samples. It is obvious from Fig. 3(a) that even a small quantity of Gd3 þ doping can change the current density of BiFeO3 dramatically. In order to study further the effects of Gd3 þ and Co3 þ co-doping on the conduction mechanism of BiFeO3 samples, the current density (lg J) versus electric field (lg E) for Bi0.95Gd0.05Fe1 xCoxO3 ceramics at room temperature is shown in Fig. 3(b). It can be seen that lg J increases linearly with lg E below a critical voltage. Namely: lg J ¼ a lg E þ b
ð1Þ
Here a is the slope of linear fitting; b is logarithm of electric current density (J) when electric field (E) is equal to 1 V; Fig. 3(b) gives the least squares fitting of the corresponding data with fitting parameters listed in Table 2. For the samples of BiFeO3, Bi0.95Gd0.05Fe1 xCoxO3, Bi0.95Sm0.05Fe0.9Co0.1O3, the slope of the curve is 1.54, 1.38, 1.38, 1.22, 1.36 rather than 1, respectively. The slope of the curve of Bi0.95Gd0.05FeO3 is 1.08, which approximately equal to 1. It may be an indication that the charge density in sample is made of a small amount of free charges and space charges [16,18]. The former transfer in line with Ohm's law with a slope 1(J pEa a 1) [16], the latter transfer in line with the Mott–Gurney law (Jp Ea a 2) with a corresponding slope 2, in consistent with that
reported in references [16,18], which showed that the interaction of the two fitting slopes was between 1 and 2. For BiFeO3, It might be due to the appearance of charged defects governed by Fe2 þ ions, oxygen vacancies Vo and/or bismuth vacancies VBi in BiFeO3 sample [23–24]. According to the following reaction mechanisms [18,25–26]: 2Bi3 þ þ 3O2 -Bi2 O3 ↑ þ 2V3 Bi þ 3V2 þ O
ð2Þ
2Fe3 þ þ O2 -2Fe2 þ þ0:5O2 ↑ þ 2VO 2 þ 2þ Vo
ð3Þ V3 Bi
vacancy and as Importantly, the presence of stated in Eqs. (2) and (3) has the predominant effect on the reduction of the electrical resistivity of the bulk samples, giving rise to high leakage currents in the samples [23–24,26]. When Bi3 þ is partially replaced by Gd3 þ , which will suppress Bi3 þ volatile because the melting point of Gd2O3 (2330 1C) is much higher than Bi2O3 (820 1C). Thus, oxygen vacancies of Bi0.95Gd0.05FeO3 become smaller than that of BiFeO3. The best leakage characteristics of Bi0.95Gd0.05FeO3 samples are obtained. It is inferred that the origin of reduced leakage current in the Bi0.95Gd0.05FeO3 is a synergetic effect of “suppression of oxygen vacancies VO and/or bismuth vacancies VBi” leading to reduced Fe2 þ ions [17,29]. It shows that Gd3 þ doping BiFeO3 enhanced the electrical properties with lower leakage current density. The result is consistent with the results reported by Singh et al. [8–10,25,28,29]. With Co3 þ introduced once again, the current density of Bi0.95Gd0.05Fe1 xCoxO3 (x¼0.05–0.2) is slightly higher than that of Bi0.95Gd0.05FeO3 from Fig. 3(a). It can be seen from Fig. 3(b) that the slope of linear fitting of Bi0.95Gd0.05 Fe1 xCoxO3 (x¼0.05–0.2) is larger than that of Bi0.95Gd0.05 FeO3 and smaller than that of BiFeO3. It shows that space
G.L. Song et al. / Materials Science in Semiconductor Processing 27 (2014) 899–908
903
Fig. 3. Current density (J) vs. applied electric field (E) characteristics for all samples measured at room temperature. (a) J–E plot of all samples. (b) lg J–lg E plots of all samples. Table 2 Fitting parameters a and b of electric current density (lg J) vs. voltage (lg E) date for all samples: a is the slope of fitting linear; b is logarithm of electric current density (J) when electric field (E) is equal to 1 v/cm; R2 is the linear regression coefficient. Samples
a
b
R2
BiFeO3 Bi0.95Gd0.05FeO3 Bi0.95Gd0.05Fe0.95Co0.0.5O3 Bi0.95Gd0.05Fe0.9Co0.1O3 Bi0.95Gd0.05Fe0.85Co0.15O3 Bi0.95Gd0.05Fe0.8Co0.2O3
1.54 1.08 1.38 1.38 1.22 1.36
7.21 7. 61 8.134 8.12 7.74 7.67
0.997 0.996 0.997 0.997 0.996 0.993
charge exists in Bi0.95Gd0.05Fe1 xCoxO3 (x¼0.05–0.2) samples. The theory of space charge formation can be found in Refs. [18,24]. The space charge takes part in long-term migration and formed space-charge-limited current resulting in the current density increases under the effect of an external electric field. Similar phenomena have been reported
Fig. 4. Magnetization vs. magnetic field curves for samples Bi0.95Gd0.05 Fe1 xCoxO3 at 50 K, 100 K and 300 K. (a) M–H for Bi0.95Gd0.05Fe1 xCoxO3 at 300 K. (b) M-H for Bi0.95Gd0.05Fe1 xCoxO3 at 100 K. (c) M–H for Bi0.95Gd0.05Fe1 xCoxO3 at 50 K.
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in the literature [16,18,27]. Panwar reported that Pr–Co codoped BiFeO3 thin films exhibited the lower leakage current than BiFeO3, due to Co ion substitution at Fe-site in BiFeO3 and the Fe ions concentration will be reduced, resulting in the decrease of the film conductivity [27]. The research works showed the conduction mechanism of Co doped and La–Co co-doped BiFeO3 thin films were found to be dominantly Ohmic in the low applied filed and Schottky emission in the high applied filed by Yang et al. [16]. Fig. 4 shows that the magnetic hysteresis loops for BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 ceramics with an applied magnetic filed up to 3 T at 300 K, 100 K and 50 K, respectively. All the samples have a complete magnetic hysteresis loops at room temperature as shown in Fig. 4(a). The saturated magnetic hysteresis loops of Bi0.95Gd0.05 Fe1 xCoxO3 are obtained with the introducing of Co3 þ . For pure BiFeO3, M–H curves exhibit roughly linear filed dependence at room temperature. But a nonlinear behavior of the M–H curve can be seen with enlarged scales as shown in the inset of Fig. 4(a). The remnant magnetization of pure BiFeO3 (Mr) is 1 10 4 emu/g, indicating that Pure BiFeO3 ceramic is weakly ferromagnetic. This result agreed very well with that reported by Du et al. [23–25]. Unexpectedly, the Gd-substituted specimen (Bi0.95Gd0.05FeO3) shows a linear dependence of magnetization on magnetic filed without hysteresis, and the curve is almost indistinguishable from that of pure BiFeO3. However, opening-up of the hysteresis loop can be seen in the inset of Fig. 4(a). The magnetization of Bi0.95Gd0.05FeO3 specimen is larger than that of pure BiFeO3 ceramic and also the Mr, Hc of Bi0.95Gd0.05FeO3 is 3.4 10 3 emu/g, 714.2 Oe, respectively. Clearly, doping with Gd3þ can improve ferromagnetic properties of BiFeO3 sample. The Gd3 þ and Co3 þ co-substituted BiFeO3 samples exhibit saturated and symmetric magnetic hysteresis loops, indicating a ferromagnetic behavior. This is consistent with zheng's results [15–16,28,29]. Zheng et al. prepared Bi0.9La0.1Fe0.95R0.05O3 (R¼Mn, Co) by sol–gel method with rapid sintering process and investigated magnetic properties of samples. They found that B-site doping of Mn or Co enhanced the magnetization of BiFeO3, but its Ms and Mr values were relatively small. Furthermore, the detailed changes in magnetization with Mn or Co content were not given [29]. When Co3 þ concentration is increased from 0.05 to 0.2, the remanent magnetization (Mr) increases from 0.33 to 1.01 emu/g and the saturation magnetization (Ms) increases from 1.05 to 0.267 emu/g. But the coercive field for all samples remains at 800 Oe. the Mr and Ms of the Bi0.95Gd0.05Fe1 xCoxO3 samples are re-plotted in Fig. 4(c) inset respectively as a function of Co3 þ concentrations x. The Ms, Mr, TN and TC parameters for BiFeO3 and Bi0.95Gd0.05Fe1 xCoxO3 ceramics are summarized in Table 1. Fig. 4(b) inset shows the magnetization for Bi0.95Gd0.05 Fe1 xCoxO3 (x¼0.05, 0.1, 0.15, 0.2) samples as a function of temperature ranging from 50 to 400 K in a field of 10 kOe. For Bi0.95Gd0.05Fe1 xCoxO3, the magnetization gradually decreases with increasing temperature from 50 to 400 K under 10 kOe magnetic fields. It can be seen that Bi0.95Gd0.05 Fe1 xCoxO3 ceramics exhibit strong ferromagnetic properties, in agreement with the result of Kim [22,30].
Four possible explanations may account for the strong ferromagnetic properties of Bi0.95Gd0.05Fe1 xCoxO3 ceramics [31]: i). The oxygen vacancies existence. BiFeO3 ceramic exhibits weak ferromagnetic because oxygen vacancies of BiFeO3 sample lead to the presence of Fe2 þ . A lot of authoritative literatures have proved that the Fe2 þ exists in the BiFeO3 by XPS experiments [21,32,33]. Hu et al. reported the asymmetric broad peak at 710.0 eV can be observed for all the prepared samples by XPS indicating the coexistence of Fe3 þ and Fe2 þ [29]. According to the report of Jun et al. [13], the presence of Fe2 þ can lead to the enhancement of magnetization. Ionic radius of Fe2 þ is almost 15% larger than that of Fe3 þ , which can cause distortion in the crystal structure, resulting in the collinear spin arrangement in [111] planes and a non-zero net magnetic moment [13]. Fe3 þ –O2 –Fe2 þ magnetic exchange should have important impact on the magnetic properties of these samples. ii). The structure distortion. With Gd3 þ and Co3 þ co-doping, the unit cell contract or decrease in lattice constants of BiFeO3 due to the substitution of smaller Gd3 þ and Co3 þ than Bi3 þ and Fe3 þ . The reduction in unit cell volume and phase transition with the increase of Co3 þ is observed in Fig. 1 and Table 1. Since the superexchange interaction is sensitive to bond length and bond angle, the structure distortion can change Fe–O bond angle or suppress the spin spiral [30]. It releases potentially weak ferromagnetic order, which leads to enhancement of remnant magnetization in BiFeO3 system by doping of Co3 þ [34,35]. When xZ0.15, the Co3 þ substitution results in structural phase transition, wherein the spin cycloid might be destructed and homogeneous spin structure formed. It is observed that the latent magnetization locked within the cycloid might be released and the significant Mr value increased. Similar phenomena have been also reported in Co3 þ and Mn4 þ doped BiFeO3 [11,12]. iii). In Bi0.95Gd0.05Fe1 xCoxO3. The space modulated spin structure collapsed on Gd3 þ and Co3 þ co-doping. And the interaction between 4f electron of Gd3 þ and 3d electron of Fe3 þ or Co3 þ leads to parallel distribution of Fe3 þ spins when a magnetic field applied. Substitution of Co3 þ with smaller magnetic moment for Fe3 þ enhances the magnetic properties of BiFeO3 [22]. This can be ascribed to two aspects. Firstly, there is a local Fe3 þ –O2–Co3 þ magnetic coupling in the samples. This coupling provides ferromagnetic interaction between Fe3 þ and Co3 þ in macroscopic scale. But its microscopic mechanism is still to be explored. Huang et al. was reported similar results [35]. Secondly, the enhancement of magnetization may be attributed to the collapse of the spiral spin anti-ferromagnetic magnetic structure of BiFeO3 to form a new ferrous magnetic structure in BiFeO3 with Co3 þ substitute Fe3 þ [12].
Therefore, Gd3 þ and Co3 þ co-doping in BiFeO3 ceramic is proved to be an effective way to modulate the structure
G.L. Song et al. / Materials Science in Semiconductor Processing 27 (2014) 899–908
and magnetic properties, which are essential for practical applications. Fig. 5 shows the variation of the magnetic moment of BiFeO3 and B0.95Gd0.05Fe1 xCoxO3 samples with temperature from 300 to 900 K at an external magnetic field of 5 kOe. For BiFeO3 and B0.95Gd0.05Fe1 xCoxO3 samples, an antiferromagnetic transition is observed, with a ferromagnetic order below 640 K, 655 K (TN), and another magnetic transition is also observed above 870 K, 780 K (TM) respectively. From Fig. 5(a), the magnetic moment of BiFeO3 and Bi0.95Gd0.05FeO3 samples decreases with increasing temperature. The magnetic moment (M) of BiFeO3 changes abruptly at 644 K showing the anti-ferromagnetic phase transition of BiFeO3 sample, which is in agreement with the reported Neel temperature (TN ¼640 K) [1,2]. The magnetic moment of Bi0.95Gd0.05FeO3 shows transformation at 655 K (TN ¼655 K) increased slightly from 644 to 655 K when doped with Gd3 þ . The increase of TN of BiFeO3 from 644 to 655 K with Gd3 þ doping can be contributed to following two factors: Firstly, in BiFeO3, the magnetic order of anti-ferromagnetic arrangement between adjacent Fe3þ planes is known as G-type anti-ferromagnetic structure which is relatively stable [5]. Therefore, the TN, TM of BiFeO3 is equal to 644 K, 870 K, respectively. With the substitution of smaller Gd3þ for the larger Bi3þ , the structure of BiFeO3 will be distorted, resulting in unit cell contraction or decrease in lattice constants. The structure distortion can change Fe–O bond angle or spin order and enhance the Fe3þ –O2 –Fe3 þ double exchange [12]. Secondly, Gd3 þ has a large magnetic moment arising from its 4f electrons. This magnetic moment will interact with the 3d electrons of Fe3 þ (f–d exchange interaction) resulting in a distortion in the ant-ferromagnetic structure of the parent material. As a result, an enhanced ferromagnetic behavior and increased Neel temperature are observed [12,19]. The second magnetic phase transition happened at 870 K for BiFeO3 and Bi0.95Gd0.05FeO3, at 780 K for Bi0.95Gd0.05 Fe1 xCoxO3 showing a possible ferromagnetic or ferrimagnetic transition. The result is consistent with Das's report. Das et al. observed an antiferromagnetic transition around 370 1C at the Neel temperature (TN ¼640 K) and another magnetic transition above 600 1C for all Bi1 xBaxFeO3 samples [36]. The magnetic phase transition temperature is defined as TM (TM ¼ 870 K) in this paper, which is different from the Curie temperature of BiFeO3 (TC ¼1100 K). At present, there are few reports about the high temperature magnetic measurements of doped BiFeO3 and no one has mentioned the magnetic phase transition at 870 K. It could not be determined whether the magnetic transitions above TN for all samples were intrinsic or from impurities. If the transition were caused by very small amount of magnetic impurities which could not be detected by XRD method, the possible candidates for the impurities would most likely be Bi2O3, Fe2O3, Fe3O4, Co2O3, Gd2O3, Bi25FeO40 and Bi2Fe4O9. But Bi2O3 is non-magnetic and unlikely to cause such a magnetic transition. Impurities such as Bi25FeO40 and Bi2Fe4O9 are often present in BiFeO3 compound. Neutron diffraction measurements showed that Bi2Fe4O9 is paramagnetic at room temperature
905
Fig. 5. Magnetization as a function of temperature for all samples from 300 K to 900 K in 5000 Oe field. (a) BiFeO3 and Bi0.95Gd0.05FeO3. (b) Bi0.95Gd0.05Fe1 xCoxO3.
and undergoes a transition to an antiferromagnetic state at TN ¼(26473) K [37] and it can be effectively eliminated by doping with rare earth elements. Magnetic material Fe2O3 can exist in two forms: α-Fe2O3 and γ-Fe2O3. The former has a Curie temperature of 1013 K and the other has an anti-ferromagnetic transition at 952 K; neither of them can possibly lead to a transition at 780 K or 870 K. The Curie temperature of Co2O3 is 1388 K while BGFC Co is paramagnetic above 780 K. So Co2O3 is also unlikely responsible for the transition. Introducing Gd3 þ does not change the magnetic phase transition temperature (TM) of BiFeO3. In addition, Gd2O3 has a Tc of 292 K. Above that the material is paramagnetic. Therefore, Gd2O3 can be excluded from candidate list. If impure phase is trace of Fe3O4, the Curie temperature is 858 K. The value is well consistent with the measured magnetic phase transition temperatures of 870 K for BiFeO3 and BGFO samples, and the magnetic phase transition temperature (TM) of the BGFO3 sample will not
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change (TM ¼870 K) with doping Co3 þ in BGFO samples. The BGFO Co samples exhibited ferromagnetic behavior at 780 K. However, we found the phase transition temperature of BGFO–Co samples moving to lower temperature of 780 K from the experimental data in Fig. 5(b). It has been shown experimentally that the strength and position of the transition at 870 K for BiFeO3 are not affected by these impurities. To sum up, there is no evidence indicating that the high temperature magnetic transition in BFO series arises from impurities. The relatively high magnetic transition temperature (TM ¼870 K) of BiFeO3 remains the same with Gd3 þ doping. Although the physical origin of this ferromagnetic/frrimagnetic transition is not clear at the moment. It seems likely that the interaction between the Fe3 þ ions in the (111) planes is stronger and reaches ferromagnetic order before the Fe3 þ –O2-–Fe3 þ antimagnetic order as the temperature decreases [13]. Substitution of Gd3 þ for Bi3 þ has little effects on the ferromagnetic order of Fe3 þ and the TM of BiFeO3. Another possibility may be that the Fe3 þ –O2–Fe2 þ interaction is stronger than the Fe3 þ –O2 –Fe3 þ interaction [9,16,20,23]. Therefore, BiFeO3 gives a higher ferrimagnetic phase transition temperature of 870 K. When Co3 þ was further introduced and increased from 0.05 to 0.2, the Gd3 þ and Co3 þ co-substituted specimens exhibited a strong ferromagnetic behavior. The magnetic moments of Bi0.95Gd0.05Fe1 xCoxO3 decreased with increasing temperature. We found that the magnetic moments of Bi0.95Gd0.05Fe1 xCoxO3 (x¼0.05, 0.1) also changed abruptly at 655 K, indicating the anti-ferromagnetic phase transition.
The magnetic moments of Bi0.95Gd0.05Fe1 xCoxO3 changed from ferromagnetic to paramagnetic at 780 K. The magnetic phase transition temperature (TM) moving to lower temperature from 870 K to 780 K demonstrates that Co3þ doping reduces the magnetic phase transition temperature (TM) compared with pure BiFeO3 samples. This enhancement in ferrromagnetic ordering and decrease of TM is not due to the impurity of Co2O3 and Fe3O4 which have rather high Curie temperature of TC ¼1388 K, 857 K, respectively. The phase transition cannot occur at 780 K if Co2O3 and Fe3O4 exist in the samples, which are in consistent with the result given in Fig. 1. Three possible explanations may account for the magnetic phase transition temperature of Bi0.95Gd0.05Fe1 xCoxO3 ceramics moving to lower temperature from 870 K to 780 K. i). As the electronic configuration of Co3þ is exactly the same as that of Fe3þ , Fe3þ –O2 –Co3 þ coupling would enhance the ferrimagnetic behavior of the material just as the Fe3 þ –O2–Fe2þ coupling. Naturally, substitution of Co3þ for Fe3 þ will increase the number of Fe3þ –O2 –Co3 þ coupling leading to much stronger magnetization. The bond angle of Fe3þ –O2 –Fe3 þ might become larger and the bond length of Fe3þ –O2 –Co3þ is greater than that of Fe3þ –O2 –Fe3 þ [22,28], resulting in a relatively weak Fe–O–Co superexchange interaction with co-doping of Gd3 þ and Co3þ . It can be a major reason that weak Fe–O–Co superexchange interaction and instable ferrous magnetic structure of Bi0.95Gd0.05Fe1 xCoxO3 lead to the magnetic
Fig. 6. Magnetization vs. magnetic field curves for samples Bi0.95Gd0.05Fe1 x CoxO3 at 700 K, 750 K, 900 K, respectively. (a) BiFeO3. (b) Bi0.95Gd0.05FeO3. (c) Bi0.95Gd0.05Fe0.95Co0.05O3. (d) Bi0.95Gd0.05Fe0.9Co0.1O3. (e) Bi0.95Gd0.05Fe0.85Co0.15O3. (f) Bi0.95Gd0.05Fe0.8Co0.2O3.
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transition temperature shifting toward lower temperature at 780 K. The enhancement of magnetization may also be attributed to the collapse of the spiral spin anti-ferromagnetic magnetic structure to form a new ferrous magnetic structure in BiFeO3 in spite of the existence of magnetic coupling between the local Fe–O–Co with Co3þ substitute Fe3þ . Since the super-exchange interaction is sensitive to bond length and bond angle, the spiral spin structure might be destroyed [30]. ii). With Co3 þ doped, the main diffraction peaks between (104) and (110) of BiFeO3 sample are overlapped (110) to single peak gradually, rhombohedra symmetry of BiFeO3 has transformed to tetragonal perovskite structure from Fig. 1. The structure phase transition was one of the reasons that could lead to the magnetic transition temperature (TM) of BiFeO3 shift from 870 K to 780 K. iii). 4f of Gd3 þ and 3d of Fe3 þ or Co3 þ form a weak exchange interaction (f–d exchange interaction) in Bi0.95Gd0.05Fe1 xCoxO3. But the f–d exchange interaction is relatively weak compared with d–d, a difference of one order of magnitude [38]. Therefore, the f–d exchange interaction does not affect the magnetic transition temperature (TM) of BiFeO3. Considering the three reasons, we may draw a conclusion that the change in magnetic transition temperature (TM) of Bi0.95Gd0.05Fe1 xCoxO3 depends mainly on the Fe3 þ –O2 –Fe3 þ super-exchange strength and the relative stability of the magnetic structure. In order to verify the fact that the ferromagnetic phase transition of BiFeO3 and Bi0.95Gd0.05Fe1 x CoxO3 samples occurred at 870 K and 780 K. We measured the magnetic hysteresis loops for all samples with an applied field up to 3 T at 700 K, 750 K and 900 K, respectively, as shown in Fig. 6. All the samples exhibit symmetric magnetic hysteresis loops at 700 K and 750 K, indicating a ferromagnetic behavior. However, all the samples exhibit linear magnetic hysteresis loops at 900 K, indicating the paramagnetic behavior. This is in consistent with Das's results [36]. Das reported that a large change in the magnetization was observed around 370 1C, which was close to the Neel temperature (TN) of Ba2 þ doped BiFeO3, and another magnetic transition which exhibited the paramagnetism was observed upon 600 1C [36]. It evidences that the ferromagnetic phase transition of Bi0.95Gd0.05Fe1 xCoxO3 samples occur at 780 K in Fig. 5 and no Co2O3 impurity exist in all the samples. Here for the first time the transition temperature has been determined. Naturally, more experimental and theoretical works are needed to clarify the exact origin of this high temperature ferrimagnetic transition. To sum up, we think that Gd3 þ and Co3 þ co-doping in BiFeO3 ceramics is proved to be an effective way to modulate the structure and improve magnetic properties, which is essential for practical applications. 4. Conclusion In conclusion, we have studied the effect of Gd3 þ and Co3 þ doping in the BiFeO3 compound on its crystalline
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structure, electronic transport, ferromagnetism properties and the high temperature magnetic phase transition. The following conclusions have been reached: The main diffraction peak between (104) and (110) of BiFeO3 samples gradually overlapped (110) to single peak with increasing x, implying that rhombohedral symmetry has transformed to orthorhombic perovskite structure. The SEM images indicated that Gd3 þ and Co3 þ doping significantly decreased the grain sizes of BiFeO3 ceramics. Gd3 þ and Co3 þ co-substitution reduced the grain size by approximately 10 times compared to the un-substituted specimen. Gd3 þ and Co3 þ co-doping BiFeO3 enhanced the electrical properties with lower leakage current. The current densities of Bi0.95Gd0.05FeO3 measured at 10 kV are approximately two orders of magnitude lower than that of pure BiFeO3. All samples possess strong ferromagnetism at room temperature except BiFeO3 and Bi0.95Gd0.05FeO3, which show rather weak ferromagnetism. The TN of Bi0.95Gd0.05FeO3 increases from 644 to 655 K, which can be ascribed to the crystal structure change and the exchange interaction between 4f electrons of Gd3 þ and 3d electrons of Fe3 þ . The change in TM of Bi0.95Gd0.05Fe1 xCoxO3 depends mainly on the Fe–O–Fe super-exchange strength and the relative stability of the magnetic structure. All the samples exhibit ferromagnetic behavior below 750 K and paramagnetic behavior above 900 K. It is evidenced that the ferromagnetic phase transition of Bi0.95Gd0.05Fe1 xCoxO3 samples occur at 780 K.
Acknowledgments This work was supported by the National Natural Science Foundation of China (Project no: U120411); Basic and Advanced Technology Research Projects in Henan Province, China (Project no: 122300410203, 122300413208); and Basic Research Program of Education Bureau of Henan Province, China (2011A140014). References [1] W. Eerenstein, N.D. Mathur, J.F. Scott, Nature 442 (2006) 759–765. [2] C.T. Nelson, P. Gao, J.R. Jokisaari, C.M. Folkman, C.B. Eom, D.G. Schlom, Science 334 (2011) 968–971. [3] L.H. Yin, W.H. Song, X.L. Jiao, W.B. Wu, X.B. Zhu, Z.R. Yang, J.M. Dai, R.L. Zhang, Y.P. Sun, J. Phys. D: Appl. Phys. 42 (2009) 205402. [4] Z. Wen, X. Shen, J.X. Wu, D. Wu, A.D. Li., J.L. Wang, Appl Phys Lett. 96 (2010) 202904-1–202904-4. [5] Igor A. Kornev, S. Lisenkov, R. Haumont, B. Dkhil, L. Bellaiche, Phys. Rev. Lett. 99 (2007) 227602-1–227602-4. [6] H. Zhang, Y.J. Liu, L.H. Pan, Acta Phys. Sin. 58 (2009) 7141. [7] P. Uniyal, K.L. Yadav, Mater. Lett. 62 (2008) 2858. (4). [8] K.S. Nalwa, A. Garg, A. Upadhyaya, Mater. Lett. 62 (2008) 878. (4). [9] G.L. Song, G.J. Ma, J. Su, T.X. Wang, H.Y. Yang, F.G. Chang, Ceram. Int. 40 (2014) 3579–3587. [10] X.Q. Zhang, Y. Sui, X.J. Wang, Y. Wang, Z. Wang, J. Alloys Compd. 507 (2010) 157–161. [11] E.M. Choi, S. Patnaik, E. Weal, S.L. Sahonta, H. Wang, Z. Bi, J. Xiong, M.G. Blamire, Q.X. Jia, J.L. Mac Manus Driscoll, Appl. Phys. Lett. 98 (2011) 012509–012509-3. [12] H. Naganum, N. Shimur, J. Miur, H. Shim, S. Yasui, K. Nishida, S. Okamur, J. Appl. Phys. 103 (2008) 07E314. [13] Y.K. Jun, S.H. Hong, Solid State Commun. 144 (2007) 329–333. [14] F.G. Chang, N. Zhang, G.L. Song, J. Phys. D: Appl. Phys. 40 (2007) 7799. [15] F.Z. Qian, J.S. Jiang, D.M. Jiang, C.M. Wang, W.G. Zhang, J. Magn. Magn. Mater. 322 (2010) 3127–3130. [16] K.G. Yang, Y.L. Zhang, S.H. Yang, B. Wang, J. Appl. Phys. 107 (2010) 124109.
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