Sintering effect on electrical properties of Li2TiO3

Sintering effect on electrical properties of Li2TiO3

Solid State Ionics 256 (2014) 29–37 Contents lists available at ScienceDirect Solid State Ionics journal homepage: Sint...

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Solid State Ionics 256 (2014) 29–37

Contents lists available at ScienceDirect

Solid State Ionics journal homepage:

Sintering effect on electrical properties of Li2TiO3 Romakanta Padhy a,⁎, Nagamalleswara Rao A. a, S.K.S. Parashar a, Kajal Parashar a, Paritosh Chaudhuri b a b

Nano Sensor Laboratory, School of Applied Sciences, KIIT University Bhubaneswar 751 024, India TBM Design Section, Institute for Plasma Research, Bhat, Gandhinagar, Gujarat 382 428, India

a r t i c l e

i n f o

Article history: Received 5 August 2013 Received in revised form 25 November 2013 Accepted 24 December 2013 Available online 17 January 2014 Keywords: Sintering AC conductivity activation energy frequency exponents impedance spectroscopy

a b s t r a c t We present in this article, the latest development in electrical properties of lithium titanate (Li2TiO3) with different sintering (800 °C, 900 °C and 1000 °C), which were prepared by high-energy ball milling technique. The transformation of β(monoclinic) → γ(cubic) phase, average crystallite size and lattice strain of Li2TiO3 ceramic at various temperatures are evaluated by X-ray diffraction patterns. The electrical properties are characterized by AC impedance spectroscopy and the experimental results show low activation energy (Ea) and high AC conductivity in the range of 10−3 Scm−1for all samples at 1 MHz frequency. The Ea value, with respect to AC conductivity at 10 kHz frequency, is approximately equal to the Ea value of relaxation time (τ). The surface microstructure analysis of all sintered samples were characterized by SEM, and good crystalline nature was observed at the 1000 °C sintered sample. The temperature dependence of AC conductivity distinctly indicates the negative temperature coefficient of resistance (NTCR) behavior. Conductivity mechanisms were interpreted by the frequency exponents (n1, n2) evaluated with the double power law followed by the Arrhenius relation. © 2014 Elsevier B.V. All rights reserved.

1. Introduction In the development of tritium breeding blankets for fusion reactors, lithium-based ceramics such as lithium orthosilicate (Li4SiO4), lithium titanate (Li2TiO3), lithium zirconate (Li2ZrO3) and lithium oxide (Li2O) have been used. Among them lithium titanate (LT) is one of the most promising tritium breeding materials due to its reasonable lithium atom density, low activation, good compatibility with structural materials, excellent tritium release performance and chemical stability [1–3]. Electrical properties may reflect some characteristic features, hence analysis of electrical charge transport in small-grained Li2TiO3 ceramics, as envisaged for tritium breeding, may contribute to gain information of certain high-energy ball-milling processes. The main attribute of the current study analyzes the electrical conductivity behavior of Li2TiO3 ceramics. In the last decade, low activation energy and higher AC conductivity of lithium titanate has not been reported. The La doped lithium titanate activation energies were reported to be 0.14 to 0.18 eV for bulk conduction and 0.41 to 0.43 eV for grain-boundary conduction prepared with different calcinations (1100 °C and 1200 °C), and sintering at 1200 °C [4]. Typical activation energies of monoclinic β-Li2TiO3 for charge transport, deduced from the DC conductivity σDC, are Ea = 0.6–0.9 eV and the bulk σDC (573 K) = 3 × 10−6 Ω− 1 cm− 1 [5]. The vanadium doped LLTO, Li0.5-xLa0.5Ti1-xVxO3 (x = 0, 0.05, 0.1 and 0.15) have activation energies and AC conductivities at room temperature (0.2570, 0.2954, 0.2939, 0.2741) and (1.21 × 10− 6, 4.438 × 10− 5, 8.546 × 10−5, ⁎ Corresponding author. Tel.: +91 9861913300 (mobile). E-mail address: [email protected] (R. Padhy). 0167-2738/$ – see front matter © 2014 Elsevier B.V. All rights reserved.

9.066 × 10−5) compound was prepared by solid state reaction method at 1300 °C temperature [6]. At high temperatures, the AC conductivity of LLTO are in the range of 1.4 × 10−5 Scm−1. All the above work distinctly indicates the requirement of high AC conductivity and low activation energy of lithium titanate for ceramic breeder blanket applications and solid state electrolytes [1–6]. In the current study, we elucidate the sintering effect on electrical properties, microstructure analysis and frequency exponent (n1 and n2) behavior of LT with temperatures, respectively. Here, we have achieved low activation energy and high AC conductivity with nanoparticles of lithium titanate (LT) by following the high-energy ball mill (HEBM) technique with different sintering temperatures. 2. Experimental procedure A series of Li2TiO3 compounds were synthesized by the solid state reaction method based on the following reaction scheme: Li2 CO3 þ TiO2 →Li2 TiO3 þ CO2 ↑ The mixture of Li2CO3 (Merck 99%) and TiO2 (Merck 99%) powders were first blended by planetary ball milling for 10 h. After every 1 h, of milling, 15 minutes was allotted for cooling the samples to decrease the generated heat during the milling period. The temperature of the powders during milling can be high due to two different reasons. First, due to the kinetic energy of the grinding medium, and second, because of the exothermic processes occurring during the mill [7]. The 10 h milled powder was calcined at 700°C for 2 h. The calcined powder


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Fig. 1. XRD patterns of nano Li2TiO3 samples at various temperatures.

was pressed uniaxially with 3 wt.% PVA (polyvinyl alcohol) solution added as binder. The circular disk samples of 13.5 mm diameter and 2.2 mm thickness by hydraulic press with 411 Mpa of applied pressure. The samples were sintered at 800 °C, 900 °C and 1000 °C for 2 h. Silver contacts were made on the opposite disc faces and heated at 700 °C for 15 minutes with a heating 10 °C per minute for electrical measurement. The phase and surface microstructure analysis of the sintered samples were characterized by X-ray diffraction and scanning electron microscopy (SEM). The electrical properties of the sintered samples were analyzed with a HIOKI 3532-50 Hi-Tester LCR with a frequency range of 100 Hz to 1 MHz with different process temperatures between 300 °C and 500 °C. 3. Results and Discussion 3.1. XRD analysis The Li2TiO3 consists of three polymorphic phases, one is the α-phase (metastable cubic); second is the β-phase (monoclinic), which is generated from the metastable by heating the sample to more than 300 °C and has a narrow homogeneity range and crystallizes in the Li2SnO3-type structure with the space group C2/c; the third one is the high-temperature γ-phase (cubic), which crystallizes in the NaCl type structure and has broad homogeneity range [8]. Laumann reported

that the reflections of the cubic phase are superimposed on the reflections of the monoclinic phase under hydrothermal conditions and assumed that β → γ possible at similar temperatures (or temperature greater than 1150 °C), irrespective of whether a dry powder is heated in a vacuum or in an aqueous lithium hydroxide solution above the critical point [9]. In the present case, the reflections of the 10-h ball-milled powder stipulate the impurity phase (raw powders) shown in Fig. 1, and the monoclinic phase that is aroused whenever the sample temperature reaches 700 °C/2 h calcination, which matches the ICDD profile nos. 33-0831 and 98-011-5055, respectively [10,11]. The lattice parameters of β-phase are a = 4.2481, b = 8.4035 and c = 9.5962 as evaluated by the PowderX software. On the other hand, diffraction patterns of different (800 °C, 900 °C and 1000 °C) sintered Li2TiO3 samples were recorded at room temperature, as shown in Fig. 1. As temperature reached 1000 °C/2 h, diffraction patterns replicated the cubic phase according to the first three major peak intensities shown in Fig. 1, and match with ICDD profile no 98-011-6830. The three major peaks in the monoclinic phase, with respect to the intensity at 18.45° (first), 43.70° (second) and 63.52° (third), these become altered whenever Li2TiO3 is in the cubic phase 43.75° (first), 18.55° (second) and 63.55° (third) and indicates the possible β → γ phase transformation [8–11]. The average crystallite size and lattice strain were calculated from the plots of sin θ versus β cos θ, as shown in Fig. 2. Thus, when β cos θ is plotted against sin θ, a straight line is obtained with the slope as lattice strain and the intercept as Kλ/L, where L is the crystallite size [12]. The increase in crystallite size and decrease in lattice strain by increasing the sintering temperature indirectly indicates stable γ -phase (cubic) of Li2TiO3 ceramic. 3.2. Microstructure analysis The surface microstructures of Li2TiO3 ceramic samples were sintered at 800 °C, 900 °C and 1000 °C for 2 h, as shown in Fig. 3. It was also found that the 1000 °C sintered sample had good crystalline and well-defined granular nature compared with the 800 °C and 900 °C samples. The average granular size was nearly 1 μm and it can be seen that a few pores were found in the interior of the 1000 °C sintered sample. The sizes of the pores were small with distribution in the range of 1 to 2 μm. The granular growth at the 900 °C sample was uniform and there much fewer pores. However, SEM micrographs show the polycrystalline nature of microstructures with good density. 3.3. Impedance analysis Impedance analysis is one of the most interesting phenomena to identify the physical process and characterization of different electrical

Fig. 2. Average crystallite size and lattice strain calculate from the plot of sin θ versus β cos θ, at various temperatures.

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Fig. 3. The surface microstructure of Li2TiO3 sintered samples for 2 h: (a) 800 °C, (b) 900 °C and (c) 1000 °C.

parameters for the appropriate system. It is useful to understand the dielectric behavior of polycrystalline materials. The most commonly used models are equivalent electric circuits consisting of resistors, capacitors, inductors and specialized distributed elements when represented in the Nyquist plot. Polycrystalline materials have both grain and grain boundary effects with different time constants. The high-frequency semicircle corresponds to the bulk property, and the low frequency corresponds to grain boundary property [13]. Complex impedance can be characterized as Z = Z′ + Z″, where Z′ and Z″ is the real and imaginary part of the impedance. The real and imaginary parts of impedance can be written as [14]: ′

Z ¼ ZcosðθÞ

and Z ¼ ZsinðθÞ

The real part of impedance (Z′) decreases with increasing frequency for all the samples at various process temperatures, as shown in Fig. 4. The magnitude value of Z′ decreases with increasing the process temperature for the 800 °C sintered sample indicates the reduction of grains, grain boundaries and electrode interface resistances. Such a

phenomenon indicates the negative temperature coefficient of resistance in Li2TiO3. The frequency dependence of the imaginary part of impedance (Z″) as a function of process temperature for the 800 °C sintered sample shows a peak at a particular frequency called the resonance peak and corresponding maxima is Z″max, as shown in Fig. 5. The magnitude value of the imaginary part of impedance (Z″) decreases at a peak position and shifting towards high-frequency side by changing the process temperatures. Similar types of behavior (Z′ and Z″) were observed for the 900 °C and 1000 °C sintered samples but magnitude values were smaller than that of the 800 °C sintered sample. The asymmetric broadening of the peaks with an increase in temperature suggests the presence of electrical process of the material with a distribution of relaxation times. This will indicate the temperature dependence of electrical relaxation phenomena on the materials [15]. The peak point in Z″ versus frequency graph indicates the relaxation frequency or resonance frequency according to the relation ωmaxτ = 1, where ωmax is the angular frequency at the Z″max. The relaxation times (τ) are calculated using the above relation and the ln(τ) versus 1/T plot for all three sintered samples have linear behavior, as shown in

Fig. 4. The frequency-dependent real part of impedance (Z') as a function of temperature for Li2TiO3 with different sintered samples for 2 h: (a) 325 °C, (b) 375 °C, (c) 425 °C and (d) 450 °C.


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Fig. 5. The frequency-dependent imaginary part of impedance (Z") as a function of temperature for sintered nano Li2TiO3 samples: (a) 325 °C, (b) 375 °C, (c) 425 °C and (d) 450 °C.

Fig. 6. The activation energy was calculated by using the following relation [16]: τ¼τ 0 expðEa =kB T Þ where kB is the Boltzmann constant, τ0 is the pre-exponential factor, T is the measuring temperature and Ea is the activation energy. From the linear fit, the activation energies were estimated to be Ea = 0.244, 0.27 and 0.266 eV for different sintering at 800 °C, 900 °C and 1000 °C, respectively. The complex impedance spectra (Nyquist plots) of different sintered samples measured at different process temperatures between 300 °C and 500 °C are shown in Fig. 7. The complex impedance spectrum (CIS) is composed of a high-frequency semicircle and a low-frequency spike in the temperature. The semicircular arcs gradually decreases with the increasing process temperature and sintering temperatures were shown in Fig. 7(a–d). The samples sintered at 800 °C and

1000 °C for 2 h has bulk (or grain) effect with active–passive transition and, on the other hand, the 900 °C/2 h sintered sample has both bulk and grain boundary effects, this may be due to the symmetric nature of the microstructure. This indicates that the material is purely temperature- and frequency-dependent. The change in bulk and grain boundary resistance leads to a change in capacitance parameters bulk capacitance (Cb), grain boundary capacitance (Cgb) and resonant frequency (fr) by the following relation [16,17]: 2π f r Rb C b ¼ 1


2π f r Rgb C gb ¼ 1

Fig. 8 compares the complex impedance plots with fitted data using commercially available Zview-software for all the samples with equivalent circuit (the parallel combination of resistor R2 and constant phase element (CPE) is connected in series with the resistor R1) [18]. The temperature-dependent electrical parameters Rb, Cb, Rgb and Cgb obtained from the fitting data measured at various temperatures of Li2TiO3

Fig. 6. Variation of relaxation time of Li2TiO3 sintered samples for 2 h: (a) 800 °C, (b) 900 °C and (c) 1000 °C.

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Fig. 7. Complex impedance spectroscopic plots of Li2TiO3 as a function of temperature with different sintered samples for 2 h: (a) 325 °C, (b) 375 °C, (c) 425 °C and (d) 450 °C.

have been shown in Fig. 9. Grain resistance and capacitance can be evaluated with the help of resonant frequency (fr) obtained from frequency versus imaginary (Z) graph (Fig. 5). Similarly, grain boundary capacitance and resistance can be evaluated with the help of resonant frequency (fr) obtained from frequency versus imaginary (M) graphs (not shown in figure), respectively [15,16]. Here, the bulk resistance and grain boundary resistance decreases with the increasing process temperature as well as sintering temperature, as shown in Fig. 9 [19]. It is well known that the Debye type relaxation has a semicircular Argand (Cole–Cole or Nyquist) plots with the center located on the z′ axis. On the other hand, for a non-Debye type relaxation, these Argand complex plane plots are close to semicircular arcs with endpoints on the real axis and the center lying below the abscissa. The complex impedance in such case can be represented as h i  ′ ″ 1‐α Z ðwÞ ¼ Z þ Z ¼ R= 1 þ ðiw=w0 Þ where α represents the magnitude of the departure of electrical response from an ideal condition and this can be evaluated from the location of the center of the semicircle, when α approaches one, that is, [(1 ‐ α) → 0] gives rise to the deviation from Debye’s formalism [20,21]. The proposed system has semicircular Argand plots with the center located below the real axis, precisely indicating the non-Debye relaxation process. The frequency-dependent complex admittance can be written as Y = Y′ + Y″ = G(W) + jB(W), where G(W) is a real part of admittance called conductance and B(W) is an imaginary part of admittance called susceptance. The real and imaginary parts of admittance evaluated by the real and imaginary parts of impedance parameters, are as follows: GðW Þ ¼

‐Z ′

. ðZ Þ

′ 2



and BðW Þ ¼


going to increase (i.e. all dipoles oriented in the same direction) for different sintered samples. In the current study, increasing the sintering temperature reduces the amount of storage energy, which indicates that the orientation of the dipoles are in opposite manner with the grains. So the behavior of susceptance distinctly indicates the polarizability of the sample, which in turn depends on externally applied parameters such as frequency and process temperature.

3.4. Electrical conductivity analysis Conductivity analysis provides significant information related to the transport of charge carriers, i.e., electron/hole or cations/anions that predominate the conduction process and their response as a function of temperature and frequency [22]. The frequency-dependent AC conductivity as a function of temperature for different sintered samples (800 °C, 900 °C and 1000 °C) were shown in Fig. 11. The transition from the DC plateau to AC conductivity dispersion region shifts toward the higher frequency range when temperature increases. The dispersion is due to the high probability for the correlated forward and backward hopping at high frequency, together with the relaxation of the dynamic

. ðZ ′ Þ2 þðZ}Þ2

The temperature dependence of the susceptance (B) as a function of frequency for different sintered samples is shown in Fig. 10. Generally, susceptance indicates the amount of energy that can be stored within a sample. The susceptance increases with increasing frequency and process temperature, indicating that the storage capacity of the sample is

Fig. 8. Complex impedance plot (symbol), fitted data (solid line) and equivalent circuit of Li2TiO3 processing at 325 °C for different sintered samples.


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Fig. 9. Temperature-dependent resistance, capacitance parameters from grain and grain boundary contributions at various sintering temperatures: (a, b) 800 °C, 1000 °C, and (c, d) 900 °C.

cage potential. Therefore, we can say that the AC conductivity is dominant in the high-frequency region [23]. The conductivity behavior obeys the universal power law [24]: σ ðωÞ ¼ σ ð0Þ þ Aω


where σ(0) is the DC conductivity of the sample, A is the temperature-dependent parameter and n represents the power law exponent depending on the temperature or frequency and ω is the angular

frequency of the applied given as


field. The constant parameter (A) can be

2 2

A ¼ πN e =6kβ T ð2α Þ where e is the electronic charge, T is the temperature, α is the polarizability of a pair of sites, and N is the number of sites per unit volume among which hopping takes place [20]. The overall trend of AC

Fig. 10. Temperature dependence of susceptance (B) as a function of frequency for Li2TiO3 with different sintered samples for 2 h: (a) 1 MHz (b) 100 kHz (c) 10 kHz.

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Fig. 11. Frequency-dependent AC conductivity as a function of temperature for Li2TiO3 with different sintered samples for 2 h: (a) 325 °C, (b) 375 °C, (c) 425 °C and (d) 450 °C.

conductivity behavior with frequency in Fig. 11, represents a "universal" power law, which is based on rigorous many-body dielectric interaction. The AC conductivity measurements have begun to appear relating to all manner of materials, and the type of frequency dependence given by the above equation has become the hallmark of hopping conduction [25]. The AC conductivity (σ) can be directly related to the imaginary part of dielectric constant (ε″) as σ = ε0ωε″ where, ε0 and ω is the permittivity of free space and angular frequency, respectively. The tangential loss of dielectric material (ε″) was calculated with the real (Z′) and imaginary parts (Z″) of the impedance parameters. The electrical

conductivity (σac) of the material is thermally activated and obeys the Arrhenius equation [19]: σ ac ¼ σ 0 expð‐Ea =K B T Þ where Ea is the activation energy, KB is the Boltzmann constant, T is the measurement temperature and σ0 is the pre-exponential factor. AC electrical conductivity of the Li2TiO3 was evaluated from complex impedance spectrum data as a function of temperature. The temperature dependence of AC conductivity as a function of frequency for different sintered samples is shown in Fig. 12(a–c). The conductivity increases

Fig. 12. Temperature-dependent AC conductivity as function of frequency for Li2TiO3 with different sintered samples for 2 h: (a) 1 MHz, (b) 100 kHz and (c) 10 kHz.


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Table 1 Activation energy (Ea) of Li2TiO3 with AC conductivity and relaxation time (τ) in terms of eV. Sintering Temperature ( °C)

Ea (eV) using Relaxation Time (S)

800 900 1000

0.24 0.27 0.26

with an increase in process temperature, which indicates the negative temperature coefficient of resistance (NTCR) behavior [26]. The activation energy (Ea) was evaluated at 10 kHz, 100 kHz and 1 MHz frequencies for different sintered pellets, and is shown in Table 1. The Ea values decreased for different sintered samples, with increasing frequency due to increments in the ionic conductivity of the sample. The ionic conductivity is a combination of both macroscopic and microscopic conduction, which is indirectly dependent on the bulk Rb and grain boundary Rgb resistances. The Rb and Rgb can be changed with the processing temperature as well as sintering temperature, which are shown in Fig. 9. At high temperature, only a single semicircle was obtained with highfrequency data, indicating the dominant behavior of grains. The sintering effect on electrical properties of Li2TiO3 due to the reduction of grain, grain boundary resistances and increment in ionic conductivity was examined in Figs. 9 and 12. From Table 1, the lowest activation energy (0.195 eV at 1 MHz frequency and 0.234 eV at 10 kHz) observed for 800 °C sintered sample and the difference of activation energy with other sintering (or with frequency) is low, which is around 0.008 to 0.02 eV. Here, the influence of sintering on activation energy is very low but it has an effect on AC conductivity values. However, low activation and high AC conductivity (10− 3 Scm−1) can be achieved by decreasing the size of the particle (≤ 50 nm) through high-energy ball-milling (HEBM) techniques. The term Aωn can often be explained on the basis of two distinct mechanisms for carrier conduction: quantum mechanical tunneling (QMT) through the barrier separating the localized sites, and correlated barrier hopping (CBH) over the same barrier. In general, the frequency dependence of AC conductivity could not be modeled using the simple Jonscher’s power law but it can be modeled by double power law [15,20]: σ ðωÞ ¼ σ ð0Þ þ A1 ω


þ A2 ω


to describe the different contributions to conductivity. The exponent n1(0 ≤ n1 ≤ 1) characterizes the low-frequency region, corresponding to translational ion hopping, and the exponent n2(0 ≤ n2 ≤ 2) characterizes the high-frequency region [27], indicating the existence of

Ea (eV) using AC conductivity (Ω−1cm−1) σac(1MHz)



0.195 0.206 0.203

0.225 0.241 0.230

0.234 0.260 0.252

well-localized relaxation/re-orientational hopping motion [20,27]. The temperature-dependent frequency exponents (n1, n2) are shown in Fig. 13. The values n1 are found to be a decreasing trend with increasing temperature due to the large polaron hopping conduction (mobility of large polarons is proportional to T-1/2). These large polarons are easily scattered by the ions and phonons [24]. In the case of 900 °C, 1000 °C sintered pellets, have both small and large polaron mechanisms according to the behavior of n2 value with temperature variations. The value of n2 initially increases from 325 °C to 375 °C due to small polaron mechanism, and then decreases up to 500 °C due to a large polaron mechanism. The small polaron hopping mechanism is thermally activated. For small polaron hopping conduction, the value of n2 increases with temperature and, correspondingly, conductivity will also increase. However, the 800 °C sintered sample has no identification of small polaron conduction means n2 value decreases with an increase in temperature due to the mobility of charge carriers.

4. Conclusion The decrement of lattice strain at various temperatures revealed the transformation of β(monoclinic) → stable γ(cubic) phase of nano Li2TiO3 ceramic prepared using the HEBM technique was examined by XRD patterns. It was observed that an increase in sintering temperature (up to a certain extent) will result in good crystalline nature. The fitting data of CIS plots replicate the non-Debye relaxation process with negative temperature coefficient of resistance (NTCR) behavior. In the present case, grain contribution plays a major role in achieving low activation energy (Ea = 0.244, 0.277 and 0.266 eV) for different sintered samples. However, high AC conductivity (10−3 Scm−1) is possible because the rate of successful hopping to unsuccessful hopping is very high, and the present system is suitable for applications such as ceramic breeder blankets and solid state electrolytes. The frequency exponent behavior was analyzed by the double power law, which is a contribution of translation hopping (DC region), short-range and localized hopping, respectively. The above results are more superior than the earlier reports like lanthanum-doped Li2TiO3 properties, which is possible only

Fig. 13. The temperature-dependent frequency exponents (n1, n2) for Li2TiO3 with different sintered samples for 2 h.

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by decreasing the average crystallite size with HEBM technique and without doping any foreign element. Acknowledgments The authors gratefully acknowledge the financial support of the Board of Research in Fusion Science and Technology (BRFST) of India under Project No. NFP-MAT-F10-01. References [1] Xiangwei Wu, Zhaoyin Wen, Xiuyan Wang, Xiaogang Xu, Jiu Lin, Shufeng Song, Fusion Eng. Des. 85 (2010) 1442–1445. [2] Shigeru Saito, Kunihiko Tsuchiya, Hiroshi Kawamura, Takayuki Terai, Satoru Tanaka, J. Nucl. Mater. 253 (1998) 213–218. [3] Zhaoyin Wen, Xiangwei Wu, Xiaogang Xu, Jiu Lin, Zhonghua Gu, Fusion Eng. Des. 85 (2010) 1551–1555. [4] C.H. Chen, K. Amine, Solid State Ionics 144 (2001) 51–57. [5] Th. Fehr, E. Schmidbauer, Solid State Ionics 178 (2007) 35–41. [6] K. Vijaya Babu, V. Veeraiah, P.S.V. Subba Rao, IJSER 3 (2012). [7] C. Suryanarayana, Prog. Mater. Sci. 46 (2001). [8] M. Mohapatra, Y.P. Naik, V. Natarajan, T.K. Seshagiri, Z. Singh, S.V. Godbole, J. Lumin. 130 (2010) 2402–2406. [9] Andreas Laumann, Kirsten Marie Ørnsbjerg Jensen, Christoffer Tyrsted, Martin Bremholm, Karl Thomas Fehr, Michael Holzapfel, Bo Brummerstedt Iversen, Eur. J. Inorg. Chem. (2011) 2221–2226.


[10] Tsuyoshi Hoshino, Kenta Tanaka, Junichi Makita, Takuya Hashimoto, J. Nucl. Mater. 367–370 (2007) 1052–1056. [11] Andrzej Deptula, Wieslawa Lada, Tadeusz Olczak, Bozena Sartowska, Andrze G. Chmielewski, Carlo Alvani, Sergio Casadio, Nukleonika 46 (3) (2001) 95–100. [12] Nagamalleswara Rao Alluri, S.K.S. Parashar, Kajal Parashar, Parthasarthi Mukherjee, B.S. Murty, Metall. Mater. Trans. A 44 (11) (2013) 5241–5250. [13] S. Sen, R.N.P. Choudhary, P. Pramanik, Br. Ceram. Trans. 103 (2004) 250–256. [14] M. Nadeem, M.J. Akhtar, A.Y. Khan, R. Shaheen, M.N. Haque, Chem. Phys. Lett. 366 (2002) 433–439. [15] S.K. Rout, A. Hussain, E. Sinha, C.W. Ahn, I.W. Kim, Solid State Sci. 11 (2009) 1144–1149. [16] D.K. Pradhan, R.N.P. Chowdhury, T.K. Nath, Appl. Nanosci. 2 (2012) 261–273. [17] H. Rahmouni, M. Nouiri, R. Jemai, N. Kallel, F. Rzigua, A. Selmi, K. Khirouni, S. Alaya, J. Magn. Magn. Mater. 316 (2007) 23–28. [18] Shrabanee Sen, R.N.P. Choudhary, Abhijit Tarafdar, P. Pramanik, J. Appl. Phys. 99 (2006) 124114. [19] Adhish Jaiswal, Raja Das, Tuhin Maity, Pankaj Poddar, J. Appl. Phys. 110 (2011) 124301. [20] Ansu K. Roy, Amrita Singh, Karishma Kumari, K. Amar Nath, Ashutosh Prasad, K. Prasad, ISRN Ceramics (2012) 10. [21] M.A.L. Nobre, S. Lanfredi, J. Phys. Chem. Solids 64 (2003) 2457–2464. [22] Archana Shukla, R.N.P. Choudhary, A.K. Thakur, D.K. Pradhan, Physica B 405 (2010) 99–106. [23] S.B.R.S. Adnan1, N.S. Mohamed, Int. J. Electrochem. Sci. 7 (2012) 9844–9858. [24] S. Sumi, P. Prabhakar Rao, M. Deepa, Peter Koshy, J. Appl. Phys. 108 (2010) 063718. [25] S.S.N. Bharadwaja, P. Victor, P. Venkateswarulu, S.B. Krupanidhi, Phys. Rev. B 65 (2002) 174106. [26] Ajay Kumar Behera, Nilaya K. Mohanty, Banarji Behera, Pratibindhya Nayak, Adv. Mat. Lett. 4 (2) (2013) 141–145. [27] A. Ahmed, A. Youssef, Z. Naturforsch. 57a (2002) 263–269.