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The effect of cation nonstoichiometry on the electrical conductivity of acceptor-doped CaZrO3
Solid State Ionics 177 (2006) 3099 – 3103 www.elsevier.com/locate/ssi
The effect of cation nonstoichiometry on the electrical conductivity of acceptor-doped CaZrO3 Soon Cheol Hwang a,b , Gyeong Man Choi b,⁎ a
New Materials and Components Research Center, Research Institute of Industrial Science and Technology, San 32, Hyoja-dong, Pohang 790-330, Korea b Fuel Cell Research Center and Department of Materials Science and Engineering, Pohang University of Science and Technology, San 31, Hyoja-dong, Pohang 790-784, Korea Received 12 May 2006; received in revised form 6 July 2006; accepted 2 August 2006
Abstract The electrical properties of acceptor-doped Ca1−xZr0.99M0.01O3−δ (M = Mg2+, In3+) systems were investigated as a function of cation nonstoichiometry (0 ≤ x ≤ 0.05). The characterization was carried out using the impedance spectroscopy between 550 °C and 1100 °C in dry air. The contributions of the grain and grain boundary conductivity to the total conductivity were obtained from the impedance data. When the Ca deficiency (x) increased, the total conductivity rapidly decreased with the corresponding increase in activation energy. Although the grain conductivity increased slightly with increasing x, the total conductivity is mostly determined by the highly resistive grain boundary. With varying x, the activation energy of total conductivity showed the percolation behavior. The percolation threshold values vary according to the doped species. It may be due to the difference in concentration of oxygen vacancies of the specimens. © 2006 Elsevier B.V. All rights reserved. Keywords: CaZrO3; Nonstoichiometry; Electrical conductivity; Acceptor; Impedance
1. Introduction Perovskite oxides of the formula ABO3 are of interest for their application in electrochemical devices, such as solid oxide fuel cells, oxygen separation membranes and oxygen sensors [1–3]. Among them, calcium zirconate (CaZrO3) has received much attention for high temperature applications, because it has high thermal and chemical stability, and good thermal shock resistance [4]. It has been widely used as a refractory material, a microwave dielectric, a solid electrolyte and a proton conductor [5,6]. Several investigators reported that it is a mixed p-type and oxygen-ion conductor in high PO2 but a stable oxygen-ion conductor in low PO2 [7–9]. When the zirconium in the oxides is partially replaced by aliovalent cation, it exhibits protonic conduction under hydrogen containing atmosphere at elevated temperature [10– 12]. Although the nonstoichiometry of this compound has been
⁎ Corresponding author. Tel.: +82 54 279 2146; fax: +82 54 279 2399. E-mail address: [email protected] (G.M. Choi). 0167-2738/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2006.08.002
reported to be very limited, there are still controversies regarding to the electrical conductivity as a function of Ca/Zr ratio. Recently, Rog et al. and Dudek and Bucko reported the relation between the conductivity and nonstoichiometry in calcium zirconate system [13,14]. However, the conductivity that they reported was the total conductivity. In order to understand the electrical conduction mechanism in polycrystalline materials, it is needed to evaluate the contributions of grain and grain boundary to the total conductivity. In a previous study, we reported that the total conductivities of Ca1−xZrO3−δ (0 ≤ x ≤ 0.1) systems decreased with increasing x up to 0.02 and suggested that the total resistance was mainly determined by the grain boundary resistance [15]. However, the deconvolution of grain and grain boundary contribution to impedance was not possible for x above 0.02 since one large semicircle was only shown. The origin of grain boundary resistance was not clear since the factors affecting the grain boundary resistance were very complex. In the present study, we have prepared acceptor-doped Ca1−xZr0.99M0.01O3−δ (M = Mg2+, In3+) specimens, expecting to minimize the influence of grain boundary on the electrical property. The electrical characterization was
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obtain impedance spectra with a frequency range between 5 Hz and 13 MHz. The temperature dependence of electrical conductivity was measured while cooling with 50 °C step and the specimens were maintained for 1 h at each temperature. The impedance spectra were analyzed by commercial impedance curve-fitting software.
Fig. 1. The temperature dependence of the total electrical conductivity (σt) of Ca1−xZr0.99Mg0.01O3−δ specimen was shown in air for 0 ≤ x ≤ 0.05. The activation energies of conductivity (Ea) were also shown.
carried out using the impedance spectroscopy. It is a useful technique for studying the conductivity of ionic or mixed conductors in polycrystalline materials, because the resistance due to grain and grain boundaries can be separated [16]. From the impedance measurements, the effect of cation nonstoichiometry on the electrical property of calcium zirconate systems was discussed. 2. Experimental procedure Ca1−xZr0.99M0.01O3−δ (M = Mg2+, In3+) ceramics (0≤x ≤ 0.05) were prepared by a solid state reaction. Appropriate amount of CaCO3 (99.9%, Showa, Japan), ZrO2 (99.9%, Tosoh, Japan), (MgCO3)4Mg(OH)2·5H2O (Extra Pure, Junsei Chemical, Japan) and In2O3 (99.97%, High Purity Chemicals, Japan) powders were ball-milled with zirconia balls in ethanol for 24 h and calcined at 1350 °C for 4 h. The calcined powders were analyzed by X-ray diffractometry (Model D-Max 1400, Rigaku, Japan) to confirm the formation of the perovskite phase. These calcined powders were ball-milled again to destroy agglomerates, and then formed into disc-shape by uniaxial pressing, followed by cold isotatic pressing (CIP) at a pressure of 200 MPa. The pellets were sintered at 1700 °C for 4 h in air. The microstructures of fractured surface were examined by FE-SEM (Model S-4300SE, Hitachi, Japan). Specimens for electrical measurements were prepared from sintered pellets by polishing the faces flat. Platinum paste (Engelhard, #6926, USA) was painted on both sides of the pellet and fired at 1000 °C for 1 h. The impedance was measured between 550 °C and 1100 °C in dry air. The LF Impedance analyzer (Model 4192A, Hewlett-Packard, USA) was used to
Fig. 2. The temperature dependence of the σg and σgb of Ca1−xZr0.99Mg0.01O3−δ specimen was shown in air for 0 ≤ x ≤ 0.035: (a) grain conductivity and (b) grain boundary conductivity.
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3. Results and discussion The temperature dependence of the total electrical conductivity (σt) of Ca1−xZr0.99Mg0.01O3−δ specimen in air for 0 ≤ x ≤ 0.05 was shown in Fig. 1. The σt was obtained from the total resistance (Rt) value on impedance spectrum. The activation energy (Ea) values, calculated by fitting the data, were also shown for each composition. When the Ca deficiency (x) increased up to x = 0.04, the σt rapidly decreased, about 1000 times at 1000 °C (for x = 0 and 0.04, σt = 1.87 × 10− 3 S/cm and 1.46 × 10− 6 S/cm, respectively), with the corresponding increase in Ea (1.04 eV to 2.46 eV). With further Ca deficiency, as shown for x = 0.05, the σt increased slightly. In order to examine the contributions of grain and grain boundary to the total conductivity, both the grain (σg) and the grain boundary conductivity (σgb) were calculated from the grain (Rg) and the grain boundary resistance (Rgb) value on impedance spectrum, respectively. In our previous paper, we reported that, when Ca deficiency was small, both the Rg and Rgb slightly increased with nonstoichiometry (x). However, when Ca deficiency was large and excess zirconia was found as a second phase, the Rgb and the contribution of the Rgb to the Rt significantly increased with x. Fig. 2 shows the temperature dependence of the σg and σgb of Ca1−xZr0.99Mg0.01O3−δ specimen in air for 0 ≤ x ≤ 0.035. As shown in Fig. 2, both σg and σgb of Ca1−xZr0.99Mg0.01O3−δ specimen decreased with increasing x. For x = 0, the value of σg, σgb and σt was 4.36 × 10− 3 S/cm, 4.95 × 10− 3 S/cm and 2.32 × 10− 3 S/cm at 600 °C, respectively. The contribution of the grain to the total conductivity was nearly equal to that of the grain boundary, since the σg and the σgb values revealed the same order of magnitude. However, it is clearly shown that the
Fig. 3. The activation energy values of grain, grain boundary and total conductivity Ca1−xZr0.99Mg0.01O3 − δ specimen in air were plotted as a function of x. Ea values increase rapidly above x = 0.025.
Fig. 4. The total electrical conductivity values (σt) of Ca1−xZr0.99In0.01O3−δ specimen in air were plotted as a function of x for 0 ≤ x ≤ 0.05. The activation energies of conductivity (Ea) were also shown.
decrease in σgb is faster than that in σg with increasing x. Note that the scale of Fig. 2a and b is the same. Also, the increase in Egb (1.32 eV to 2.23 eV) is much greater than that in Eg (0.93 eV to 1.37 eV). For x = 0.035, the value of σg, σgb and σt was 2.31 × 10− 5 S/cm, 7.26 × 10− 7 S/cm and 7.04 × 10− 7 S/cm at 800 °C, respectively. With increasing x, the σt value became very close to the σgb even at high temperature. These results indicate that, for x = 0.035, the main contribution to the total conductivity is the grain boundary conductivity, thus the σt is mostly determined by the σgb. It has been generally recognized that a highly resistive grain boundary might be induced by excess amorphous phases located at grain boundaries, or secondary phases and precipitates having lower conductivity in perovskite structure [16–18]. A similar observation was made in this study and is summarized below. From the XRD analysis and SEM observation, it was revealed that the Ca1−xZr0.99Mg0.01O3−δ can tolerate a very limited amount of nonstoichiometry. When the Ca deficiency increased up to x = 0.03, Ca1−xZr0.99Mg0.01O3−δ specimen showed single phase of the orthorhombic perovskite similar to CaZrO3. However, the specimen with x ≥ 0.035 showed second phase ZrO2. In addition, the relative dielectric constant of grain boundary (εgb), obtained from impedance spectrum, decreased markedly with increasing x (εgb = 1692 for x = 0, εgb = 263 for x = 0.035), whereas the grain size of the sintered specimen slightly decreased. Thus, it is suggested that the grain boundary compositions of Ca1−xZr0.99Mg0.01O3−δ specimen may be variable with increasing x and responsible for the significant decrease in the σgb. On the other hand, the decrease in the σg can be explained by the relevant defect model in Ca1−xZr0.99Mg0.01O3−δ systems. At
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conductivity than the stoichiometric composition. However, this prediction does not agree with the observed conductivity results shown in Figs. 1 and 2. Therefore, this model is not appropriate. Another possible assumption is substitution of excess zirconium ion to A-site in ABO3 perovskite structure. This defect model involves the annihilation of oxygen vacancies according to the following reaction (2), CaO
ZrO2 þ VddO Y ZrddCa þ 2OO
ð2Þ
Shima and Haile proposed that Ba deficiencies could be accommodated by the redistribution of doped ion (Gd3+) to Ba2+ site in BaxCe1−yGdyO3 systems [19]. The similar behavior was also reported by Ma et al. [20]. They suggested that, in the Ba-deficient BaxCeO3, the excess Ce4+ transfers to Ba2+ site until the oxygen vacancy concentration is adjusted to zero. Thus, it can be considered that, in Ca1−xZr0.99Mg0.01O3−δ systems, the transfer of a small amount of excess Zr4+ to Ca2+ site may reduce the concentration of oxygen vacancies which was generated by acceptor doping. This model is in accordance with the decreasing σg with x in Fig. 2a. Fig. 3 shows the activation energy of Ca1−xZr0.99Mg0.01O3−δ specimen in air as a function of x. The entire composition range may be divided into two regions following the shape of curve and the XRD and the SEM observations. The percolation phenomena were clearly observed for σg, σgb and σt. The boundary line may be drawn between x = 0.03 and x = 0.035 where percolation threshold concentration is defined for all three conductivities. Below x = 0.025, the value of Ea was almost the same. However, from x = 0.03, the Ea significantly increased with increasing x. For x above 0.04, the Ea change was small. It is apparent that the Ea of σt was mainly determined by Ea of σgb.
Fig. 5. The temperature dependence of the σg and σgb of Ca1−xZr0.99Mg0.01O3−δ specimen was shown in air for 0 ≤x ≤ 0.035, respectively, for (a) grain conductivity and (b) grain boundary conductivity.
first, one can consider the possibility that the calcium deficiency creates cation vacancies at Ca2+ site, ð1−xÞCaO þ ZrO2 →ð1−xÞCaCa þ xVCa ″ þ xVddO þ ZrZr þ ð3−xÞOO
ð1Þ
The introduction of cation vacancies should lead to the formation of oxygen vacancies. With increasing x, it should exhibit higher
Fig. 6. The activation energies of σt of Ca1−xZrO3 and Ca1−xZr0.99M0.01O3−δ specimen (M = Mg, In) in air were compared as a function of x. Rapid increase in Ea value was shown at higher x values for Mg- and In-doped samples than that for undoped sample.
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Figs. 4 and 5 show the temperature dependence of the σt, σg and σgb of Ca1−xZr0.99In0.01O3−δ specimen in air for 0 ≤ x ≤ 0.05, respectively. Similar with Mg-doped specimen, the σt rapidly decreased with increasing x up to x = 0.03, about 1000 times at 1000 °C (for x = 0 and x = 0.03, σ t = 5.78 × 10 − 3 and 2.52 × 10− 6 S/cm, respectively), with the corresponding increase in Ea (1.03 eV to 2.48 eV). Both the σg and the σgb values also decreased with increasing x. In In-doped specimen, unfortunately, the deconvolution of the σg and σgb for x above 0.03 was not possible due to the severe overlapping between the grain and the grain boundary semicircles on impedance spectrum. Fig. 6 shows the comparison of the activation energy of σt of Ca1−xZr0.99M0.01O3−δ specimen (M = Mg2+, In3+) in air as a function of x. For all compositions, the percolation behavior in Ea thus in σt was clearly shown. As previously mentioned, it is mainly due to the increasing Ea in σgb. However, the percolation threshold value was somewhat different, and it was x ~ 0.015, 0.026, 0.032, respectively, for undoped, In-doped and Mgdoped specimens. The difference may be due to the difference in concentration of oxygen vacancies of the samples. When the acceptors (Mg2+, In3+) were substituted for Zr4+ ion in CaZrO3, the charge imbalance creates oxygen vacancies as shown below for Mg or In, ZrO2
ð3Þ
2ZrO2
ð4Þ
MgO Y Mg WZr þ OO þ VddO In2 O3 Y 2InVZr þ 3OO þ VddO
Since doped or undoped CaZrO3 is a mixed (oxygen ion and hole) conductor, the main charge carriers, electron holes, are subsequently generated by the following reaction. 1=2O2 þ VddO → OO þ 2hd
ð5Þ
That is, the high concentration of oxygen vacancies results in the high hole-conductivity. From this point of view, the divalent ion (Mg2+) is more efficient to generate oxygen vacancies and thus electron holes than trivalent ion (In3+). Consequently, since the increase in Ea is related to the oxygen vacancy and thus electron-hole annihilation, it is suggested that the acceptor doping is very effective method to enlarge the composition region of conductive Ca1−xZr0.99M0.01O3−δ phase, up to x = 0.01 for undoped to x N ~ 0.02 for Mg- or In-doped CaZrO3. 4. Conclusions The electrical property of acceptor-doped Ca1−xZr0.99M0.01O3−δ (M = Mg2+, In3+) systems have been investigated as a function of cation nonstoichiometry (0 V x V 0.05) and temperature (550 °C– 1100 °C). When the Ca deficiency (x) increased, the σt rapidly decreased with the corresponding increase in Ea. With increasing x, the σt is mostly determined by the σgb. From the XRD analysis and SEM observation, it is revealed that the Ca1−xZr0.99Mg0.01O3−δ
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can tolerate a very limited amount of nonstoichiometry. It is suggested that the grain boundary compositions is variable with x. The decrease in the σg with x was explained by the relevant defect model. The transfer of a small amount of excess Zr4+ to Ca2+ site may be responsible for decreasing the concentration of oxygen vacancies and electron holes, and thus decreasing σg. From the comparison of the activation energy values of σt for undoped, Mgdoped and In-doped samples as a function of x, the percolation behavior in the Ea and thus σt was observed. The different threshold values may be due to the difference in concentration of oxygen vacancies and thus electron holes of the samples. Consequently, it is suggested that acceptor doping is very effective method to enlarge the composition region of conductive Ca1−xZr0.99M0.01O3−δ phase, up to x = 0.01 for undoped to x N ~ 0.02 for Mg- or In-doped CaZrO3. Acknowledgement This work was supported by POSCO, Korea. References [1] H. Iwahara, T. Esaka, H. Uchida, N. Maeda, Solid State Ionics 3/4 (1981) 359. [2] H.J.M Bouwmeester, A.J. Burggraaf, in: P.J. Gellings, H.J.M. Bouwmeester (Eds.), The CRC Handbook of Solid State Electrochemistry, CRC Press, Inc., Boca Raton, 1997, p. 481. [3] D. Janke, Metallurgical Transactions B (1982) 227. [4] V. Longo, F. Marchini, F. Ricciardiello, A. de Pretis, La Ceramica 34 (1981) 23. [5] T. Yajima, K. Koide, H. Takai, N. Fukatsu, H. Iwahara, Solid State Ionics 79 (1995) 33. [6] N. Fukatsu, N. Kurita, K. Koide, T. Ohashi, Solid State Ionics 113–115 (1998) 219. [7] V. Longo, F. Ricciardiello, Silicate Industries 1 (1981) 3. [8] C. Wang, X. Xu, H. Yu, Y. Wen, K. Zhao, Solid State Ionics 28–30 (1988) 542. [9] S.S. Pandit, A. Weyl, D. Janke, Solid State Ionics 69 (1994) 93. [10] T. Yajima, H. Kazeoka, T. Yogo, H. Iwahara, Solid State Ionics 47 (1991) 271. [11] H. Iwahara, T. Yajima, T. Hibino, K. Ozaki, H. Suzuki, Solid State Ionics 61 (1993) 65. [12] N. Kurita, N. Fukatsu, K. Ito, T. Ohashi, Journal of the Electrochemical Society 142 (1995) 1552. [13] G. Rog, M. Dudek, A. Kozlowska-Rog, M. Bucko, Electrochemica Acta 47 (2002) 4523. [14] M. Dudek, M. Bucko, Solid State Ionics 157 (2003) 183. [15] S.C. Hwang, G.M. Choi, Journal of the European Ceramic Society 25 (2005) 2609. [16] S.M. Haile, D.L. West, J. Campbell, Journal of Materials Research 13 (1998) 1576. [17] J.W. Stevenson, T.R. Armstrong, L.R. Pederson, J. Li, C.A. Lewinsohn, S. Baskaran, Solid State Ionics 113–115 (1998) 571. [18] B. Gharbage, F.M. Marques, J.R. Frade, Journal of the European Ceramic Society 16 (1996) 1149. [19] D. Shima, S.M. Haile, Solid State Ionics 97 (1997) 443. [20] G. Ma, H. Matsumoto, H. Iwahara, Solid State Ionics 122 (1999) 237.
The effect of cation nonstoichiometry on the electrical conductivity of acceptor-doped CaZrO3 Soon Cheol Hwang a,b , Gyeong Man Choi b,⁎ a
New Materials and Components Research Center, Research Institute of Industrial Science and Technology, San 32, Hyoja-dong, Pohang 790-330, Korea b Fuel Cell Research Center and Department of Materials Science and Engineering, Pohang University of Science and Technology, San 31, Hyoja-dong, Pohang 790-784, Korea Received 12 May 2006; received in revised form 6 July 2006; accepted 2 August 2006
Abstract The electrical properties of acceptor-doped Ca1−xZr0.99M0.01O3−δ (M = Mg2+, In3+) systems were investigated as a function of cation nonstoichiometry (0 ≤ x ≤ 0.05). The characterization was carried out using the impedance spectroscopy between 550 °C and 1100 °C in dry air. The contributions of the grain and grain boundary conductivity to the total conductivity were obtained from the impedance data. When the Ca deficiency (x) increased, the total conductivity rapidly decreased with the corresponding increase in activation energy. Although the grain conductivity increased slightly with increasing x, the total conductivity is mostly determined by the highly resistive grain boundary. With varying x, the activation energy of total conductivity showed the percolation behavior. The percolation threshold values vary according to the doped species. It may be due to the difference in concentration of oxygen vacancies of the specimens. © 2006 Elsevier B.V. All rights reserved. Keywords: CaZrO3; Nonstoichiometry; Electrical conductivity; Acceptor; Impedance
1. Introduction Perovskite oxides of the formula ABO3 are of interest for their application in electrochemical devices, such as solid oxide fuel cells, oxygen separation membranes and oxygen sensors [1–3]. Among them, calcium zirconate (CaZrO3) has received much attention for high temperature applications, because it has high thermal and chemical stability, and good thermal shock resistance [4]. It has been widely used as a refractory material, a microwave dielectric, a solid electrolyte and a proton conductor [5,6]. Several investigators reported that it is a mixed p-type and oxygen-ion conductor in high PO2 but a stable oxygen-ion conductor in low PO2 [7–9]. When the zirconium in the oxides is partially replaced by aliovalent cation, it exhibits protonic conduction under hydrogen containing atmosphere at elevated temperature [10– 12]. Although the nonstoichiometry of this compound has been
⁎ Corresponding author. Tel.: +82 54 279 2146; fax: +82 54 279 2399. E-mail address: [email protected] (G.M. Choi). 0167-2738/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2006.08.002
reported to be very limited, there are still controversies regarding to the electrical conductivity as a function of Ca/Zr ratio. Recently, Rog et al. and Dudek and Bucko reported the relation between the conductivity and nonstoichiometry in calcium zirconate system [13,14]. However, the conductivity that they reported was the total conductivity. In order to understand the electrical conduction mechanism in polycrystalline materials, it is needed to evaluate the contributions of grain and grain boundary to the total conductivity. In a previous study, we reported that the total conductivities of Ca1−xZrO3−δ (0 ≤ x ≤ 0.1) systems decreased with increasing x up to 0.02 and suggested that the total resistance was mainly determined by the grain boundary resistance [15]. However, the deconvolution of grain and grain boundary contribution to impedance was not possible for x above 0.02 since one large semicircle was only shown. The origin of grain boundary resistance was not clear since the factors affecting the grain boundary resistance were very complex. In the present study, we have prepared acceptor-doped Ca1−xZr0.99M0.01O3−δ (M = Mg2+, In3+) specimens, expecting to minimize the influence of grain boundary on the electrical property. The electrical characterization was
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obtain impedance spectra with a frequency range between 5 Hz and 13 MHz. The temperature dependence of electrical conductivity was measured while cooling with 50 °C step and the specimens were maintained for 1 h at each temperature. The impedance spectra were analyzed by commercial impedance curve-fitting software.
Fig. 1. The temperature dependence of the total electrical conductivity (σt) of Ca1−xZr0.99Mg0.01O3−δ specimen was shown in air for 0 ≤ x ≤ 0.05. The activation energies of conductivity (Ea) were also shown.
carried out using the impedance spectroscopy. It is a useful technique for studying the conductivity of ionic or mixed conductors in polycrystalline materials, because the resistance due to grain and grain boundaries can be separated [16]. From the impedance measurements, the effect of cation nonstoichiometry on the electrical property of calcium zirconate systems was discussed. 2. Experimental procedure Ca1−xZr0.99M0.01O3−δ (M = Mg2+, In3+) ceramics (0≤x ≤ 0.05) were prepared by a solid state reaction. Appropriate amount of CaCO3 (99.9%, Showa, Japan), ZrO2 (99.9%, Tosoh, Japan), (MgCO3)4Mg(OH)2·5H2O (Extra Pure, Junsei Chemical, Japan) and In2O3 (99.97%, High Purity Chemicals, Japan) powders were ball-milled with zirconia balls in ethanol for 24 h and calcined at 1350 °C for 4 h. The calcined powders were analyzed by X-ray diffractometry (Model D-Max 1400, Rigaku, Japan) to confirm the formation of the perovskite phase. These calcined powders were ball-milled again to destroy agglomerates, and then formed into disc-shape by uniaxial pressing, followed by cold isotatic pressing (CIP) at a pressure of 200 MPa. The pellets were sintered at 1700 °C for 4 h in air. The microstructures of fractured surface were examined by FE-SEM (Model S-4300SE, Hitachi, Japan). Specimens for electrical measurements were prepared from sintered pellets by polishing the faces flat. Platinum paste (Engelhard, #6926, USA) was painted on both sides of the pellet and fired at 1000 °C for 1 h. The impedance was measured between 550 °C and 1100 °C in dry air. The LF Impedance analyzer (Model 4192A, Hewlett-Packard, USA) was used to
Fig. 2. The temperature dependence of the σg and σgb of Ca1−xZr0.99Mg0.01O3−δ specimen was shown in air for 0 ≤ x ≤ 0.035: (a) grain conductivity and (b) grain boundary conductivity.
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3. Results and discussion The temperature dependence of the total electrical conductivity (σt) of Ca1−xZr0.99Mg0.01O3−δ specimen in air for 0 ≤ x ≤ 0.05 was shown in Fig. 1. The σt was obtained from the total resistance (Rt) value on impedance spectrum. The activation energy (Ea) values, calculated by fitting the data, were also shown for each composition. When the Ca deficiency (x) increased up to x = 0.04, the σt rapidly decreased, about 1000 times at 1000 °C (for x = 0 and 0.04, σt = 1.87 × 10− 3 S/cm and 1.46 × 10− 6 S/cm, respectively), with the corresponding increase in Ea (1.04 eV to 2.46 eV). With further Ca deficiency, as shown for x = 0.05, the σt increased slightly. In order to examine the contributions of grain and grain boundary to the total conductivity, both the grain (σg) and the grain boundary conductivity (σgb) were calculated from the grain (Rg) and the grain boundary resistance (Rgb) value on impedance spectrum, respectively. In our previous paper, we reported that, when Ca deficiency was small, both the Rg and Rgb slightly increased with nonstoichiometry (x). However, when Ca deficiency was large and excess zirconia was found as a second phase, the Rgb and the contribution of the Rgb to the Rt significantly increased with x. Fig. 2 shows the temperature dependence of the σg and σgb of Ca1−xZr0.99Mg0.01O3−δ specimen in air for 0 ≤ x ≤ 0.035. As shown in Fig. 2, both σg and σgb of Ca1−xZr0.99Mg0.01O3−δ specimen decreased with increasing x. For x = 0, the value of σg, σgb and σt was 4.36 × 10− 3 S/cm, 4.95 × 10− 3 S/cm and 2.32 × 10− 3 S/cm at 600 °C, respectively. The contribution of the grain to the total conductivity was nearly equal to that of the grain boundary, since the σg and the σgb values revealed the same order of magnitude. However, it is clearly shown that the
Fig. 3. The activation energy values of grain, grain boundary and total conductivity Ca1−xZr0.99Mg0.01O3 − δ specimen in air were plotted as a function of x. Ea values increase rapidly above x = 0.025.
Fig. 4. The total electrical conductivity values (σt) of Ca1−xZr0.99In0.01O3−δ specimen in air were plotted as a function of x for 0 ≤ x ≤ 0.05. The activation energies of conductivity (Ea) were also shown.
decrease in σgb is faster than that in σg with increasing x. Note that the scale of Fig. 2a and b is the same. Also, the increase in Egb (1.32 eV to 2.23 eV) is much greater than that in Eg (0.93 eV to 1.37 eV). For x = 0.035, the value of σg, σgb and σt was 2.31 × 10− 5 S/cm, 7.26 × 10− 7 S/cm and 7.04 × 10− 7 S/cm at 800 °C, respectively. With increasing x, the σt value became very close to the σgb even at high temperature. These results indicate that, for x = 0.035, the main contribution to the total conductivity is the grain boundary conductivity, thus the σt is mostly determined by the σgb. It has been generally recognized that a highly resistive grain boundary might be induced by excess amorphous phases located at grain boundaries, or secondary phases and precipitates having lower conductivity in perovskite structure [16–18]. A similar observation was made in this study and is summarized below. From the XRD analysis and SEM observation, it was revealed that the Ca1−xZr0.99Mg0.01O3−δ can tolerate a very limited amount of nonstoichiometry. When the Ca deficiency increased up to x = 0.03, Ca1−xZr0.99Mg0.01O3−δ specimen showed single phase of the orthorhombic perovskite similar to CaZrO3. However, the specimen with x ≥ 0.035 showed second phase ZrO2. In addition, the relative dielectric constant of grain boundary (εgb), obtained from impedance spectrum, decreased markedly with increasing x (εgb = 1692 for x = 0, εgb = 263 for x = 0.035), whereas the grain size of the sintered specimen slightly decreased. Thus, it is suggested that the grain boundary compositions of Ca1−xZr0.99Mg0.01O3−δ specimen may be variable with increasing x and responsible for the significant decrease in the σgb. On the other hand, the decrease in the σg can be explained by the relevant defect model in Ca1−xZr0.99Mg0.01O3−δ systems. At
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conductivity than the stoichiometric composition. However, this prediction does not agree with the observed conductivity results shown in Figs. 1 and 2. Therefore, this model is not appropriate. Another possible assumption is substitution of excess zirconium ion to A-site in ABO3 perovskite structure. This defect model involves the annihilation of oxygen vacancies according to the following reaction (2), CaO
ZrO2 þ VddO Y ZrddCa þ 2OO
ð2Þ
Shima and Haile proposed that Ba deficiencies could be accommodated by the redistribution of doped ion (Gd3+) to Ba2+ site in BaxCe1−yGdyO3 systems [19]. The similar behavior was also reported by Ma et al. [20]. They suggested that, in the Ba-deficient BaxCeO3, the excess Ce4+ transfers to Ba2+ site until the oxygen vacancy concentration is adjusted to zero. Thus, it can be considered that, in Ca1−xZr0.99Mg0.01O3−δ systems, the transfer of a small amount of excess Zr4+ to Ca2+ site may reduce the concentration of oxygen vacancies which was generated by acceptor doping. This model is in accordance with the decreasing σg with x in Fig. 2a. Fig. 3 shows the activation energy of Ca1−xZr0.99Mg0.01O3−δ specimen in air as a function of x. The entire composition range may be divided into two regions following the shape of curve and the XRD and the SEM observations. The percolation phenomena were clearly observed for σg, σgb and σt. The boundary line may be drawn between x = 0.03 and x = 0.035 where percolation threshold concentration is defined for all three conductivities. Below x = 0.025, the value of Ea was almost the same. However, from x = 0.03, the Ea significantly increased with increasing x. For x above 0.04, the Ea change was small. It is apparent that the Ea of σt was mainly determined by Ea of σgb.
Fig. 5. The temperature dependence of the σg and σgb of Ca1−xZr0.99Mg0.01O3−δ specimen was shown in air for 0 ≤x ≤ 0.035, respectively, for (a) grain conductivity and (b) grain boundary conductivity.
first, one can consider the possibility that the calcium deficiency creates cation vacancies at Ca2+ site, ð1−xÞCaO þ ZrO2 →ð1−xÞCaCa þ xVCa ″ þ xVddO þ ZrZr þ ð3−xÞOO
ð1Þ
The introduction of cation vacancies should lead to the formation of oxygen vacancies. With increasing x, it should exhibit higher
Fig. 6. The activation energies of σt of Ca1−xZrO3 and Ca1−xZr0.99M0.01O3−δ specimen (M = Mg, In) in air were compared as a function of x. Rapid increase in Ea value was shown at higher x values for Mg- and In-doped samples than that for undoped sample.
S.C. Hwang, G.M. Choi / Solid State Ionics 177 (2006) 3099–3103
Figs. 4 and 5 show the temperature dependence of the σt, σg and σgb of Ca1−xZr0.99In0.01O3−δ specimen in air for 0 ≤ x ≤ 0.05, respectively. Similar with Mg-doped specimen, the σt rapidly decreased with increasing x up to x = 0.03, about 1000 times at 1000 °C (for x = 0 and x = 0.03, σ t = 5.78 × 10 − 3 and 2.52 × 10− 6 S/cm, respectively), with the corresponding increase in Ea (1.03 eV to 2.48 eV). Both the σg and the σgb values also decreased with increasing x. In In-doped specimen, unfortunately, the deconvolution of the σg and σgb for x above 0.03 was not possible due to the severe overlapping between the grain and the grain boundary semicircles on impedance spectrum. Fig. 6 shows the comparison of the activation energy of σt of Ca1−xZr0.99M0.01O3−δ specimen (M = Mg2+, In3+) in air as a function of x. For all compositions, the percolation behavior in Ea thus in σt was clearly shown. As previously mentioned, it is mainly due to the increasing Ea in σgb. However, the percolation threshold value was somewhat different, and it was x ~ 0.015, 0.026, 0.032, respectively, for undoped, In-doped and Mgdoped specimens. The difference may be due to the difference in concentration of oxygen vacancies of the samples. When the acceptors (Mg2+, In3+) were substituted for Zr4+ ion in CaZrO3, the charge imbalance creates oxygen vacancies as shown below for Mg or In, ZrO2
ð3Þ
2ZrO2
ð4Þ
MgO Y Mg WZr þ OO þ VddO In2 O3 Y 2InVZr þ 3OO þ VddO
Since doped or undoped CaZrO3 is a mixed (oxygen ion and hole) conductor, the main charge carriers, electron holes, are subsequently generated by the following reaction. 1=2O2 þ VddO → OO þ 2hd
ð5Þ
That is, the high concentration of oxygen vacancies results in the high hole-conductivity. From this point of view, the divalent ion (Mg2+) is more efficient to generate oxygen vacancies and thus electron holes than trivalent ion (In3+). Consequently, since the increase in Ea is related to the oxygen vacancy and thus electron-hole annihilation, it is suggested that the acceptor doping is very effective method to enlarge the composition region of conductive Ca1−xZr0.99M0.01O3−δ phase, up to x = 0.01 for undoped to x N ~ 0.02 for Mg- or In-doped CaZrO3. 4. Conclusions The electrical property of acceptor-doped Ca1−xZr0.99M0.01O3−δ (M = Mg2+, In3+) systems have been investigated as a function of cation nonstoichiometry (0 V x V 0.05) and temperature (550 °C– 1100 °C). When the Ca deficiency (x) increased, the σt rapidly decreased with the corresponding increase in Ea. With increasing x, the σt is mostly determined by the σgb. From the XRD analysis and SEM observation, it is revealed that the Ca1−xZr0.99Mg0.01O3−δ
3103
can tolerate a very limited amount of nonstoichiometry. It is suggested that the grain boundary compositions is variable with x. The decrease in the σg with x was explained by the relevant defect model. The transfer of a small amount of excess Zr4+ to Ca2+ site may be responsible for decreasing the concentration of oxygen vacancies and electron holes, and thus decreasing σg. From the comparison of the activation energy values of σt for undoped, Mgdoped and In-doped samples as a function of x, the percolation behavior in the Ea and thus σt was observed. The different threshold values may be due to the difference in concentration of oxygen vacancies and thus electron holes of the samples. Consequently, it is suggested that acceptor doping is very effective method to enlarge the composition region of conductive Ca1−xZr0.99M0.01O3−δ phase, up to x = 0.01 for undoped to x N ~ 0.02 for Mg- or In-doped CaZrO3. Acknowledgement This work was supported by POSCO, Korea. References [1] H. Iwahara, T. Esaka, H. Uchida, N. Maeda, Solid State Ionics 3/4 (1981) 359. [2] H.J.M Bouwmeester, A.J. Burggraaf, in: P.J. Gellings, H.J.M. Bouwmeester (Eds.), The CRC Handbook of Solid State Electrochemistry, CRC Press, Inc., Boca Raton, 1997, p. 481. [3] D. Janke, Metallurgical Transactions B (1982) 227. [4] V. Longo, F. Marchini, F. Ricciardiello, A. de Pretis, La Ceramica 34 (1981) 23. [5] T. Yajima, K. Koide, H. Takai, N. Fukatsu, H. Iwahara, Solid State Ionics 79 (1995) 33. [6] N. Fukatsu, N. Kurita, K. Koide, T. Ohashi, Solid State Ionics 113–115 (1998) 219. [7] V. Longo, F. Ricciardiello, Silicate Industries 1 (1981) 3. [8] C. Wang, X. Xu, H. Yu, Y. Wen, K. Zhao, Solid State Ionics 28–30 (1988) 542. [9] S.S. Pandit, A. Weyl, D. Janke, Solid State Ionics 69 (1994) 93. [10] T. Yajima, H. Kazeoka, T. Yogo, H. Iwahara, Solid State Ionics 47 (1991) 271. [11] H. Iwahara, T. Yajima, T. Hibino, K. Ozaki, H. Suzuki, Solid State Ionics 61 (1993) 65. [12] N. Kurita, N. Fukatsu, K. Ito, T. Ohashi, Journal of the Electrochemical Society 142 (1995) 1552. [13] G. Rog, M. Dudek, A. Kozlowska-Rog, M. Bucko, Electrochemica Acta 47 (2002) 4523. [14] M. Dudek, M. Bucko, Solid State Ionics 157 (2003) 183. [15] S.C. Hwang, G.M. Choi, Journal of the European Ceramic Society 25 (2005) 2609. [16] S.M. Haile, D.L. West, J. Campbell, Journal of Materials Research 13 (1998) 1576. [17] J.W. Stevenson, T.R. Armstrong, L.R. Pederson, J. Li, C.A. Lewinsohn, S. Baskaran, Solid State Ionics 113–115 (1998) 571. [18] B. Gharbage, F.M. Marques, J.R. Frade, Journal of the European Ceramic Society 16 (1996) 1149. [19] D. Shima, S.M. Haile, Solid State Ionics 97 (1997) 443. [20] G. Ma, H. Matsumoto, H. Iwahara, Solid State Ionics 122 (1999) 237.