Y0.08Sr0.92FexTi1−xO3−δ perovskite for solid oxide fuel cell anodes

Materials Science and Engineering B 177 (2012) 151–156

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Y0.08 Sr0.92 Fex Ti1−x O3−ı perovskite for solid oxide fuel cell anodes Jong Seol Yoon a , Mi Young Yoon a , Chan Kwak b , Hee Jung Park b , Sang Mok Lee b , Kyu Hyoung Lee b , Hae Jin Hwang a,∗ a b

Division of Materials Science and Engineering, Inha University, 253 Yonghyun-dong, Nam-gu, Incheon 402-751, Republic of Korea Samsung Advanced Institute of Technology (SAIT), 14-1 Nongseo-dong, Yongin-si 446-712, Republic of Korea

a r t i c l e

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Article history: Received 11 March 2011 Received in revised form 15 August 2011 Accepted 31 October 2011 Available online 15 November 2011 Keywords: Perovskite Solid oxide fuel cell Oxide anode Doped strontium titanate

a b s t r a c t A single phase perovskite Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.05, 0.1,0 0.20, 0.25, 0.40, and 0.50) was fabricated at 1400 ◦ C in air by a solid state reaction method and its electrical conductivity and electrochemical properties as an anode were investigated as a function of the Fe content. Doping with Y for Sr allowed the SrFex Ti1−x O3−ı perovskite to be stable at 800 ◦ C in a reducing atmosphere. At 900 ◦ C, metallic Fe precipitated and the stability of the perovskite phase under a reducing atmosphere decreased as the Fe content increased. The conductivity of Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.40) was greater than that of the x = 0.20 sample. The conductivity of Y0.08 Sr0.92 Fex Ti1−x O3−ı was found to be 2 × 10−1 Scm−1 at 800 ◦ C in H2 . Sintering the Y0.08 Sr0.92 Fex Ti1−x O3−ı anode at 1200 ◦ C was found to be optimum to obtain not only good interfacial adhesion, but also a fine grain structure. The Y0.08 Sr0.92 Fe0.25 Ti0.75 O3−ı anode exhibited the lowest polarization resistance (0.7 and 1.8 cm2 at 800 and 700 ◦ C). © 2011 Elsevier B.V. All rights reserved.

1. Introduction Ni-YSZ (yttria-stabilized zirconia) cermets have been widely used as the anodes of solid oxide fuel cells (SOFCs), because of their excellent catalytic activity for H2 oxidation, high electrical conductivity and good chemical/mechanical compatibility with the YSZ electrolyte. Their electrochemical characteristics and fabrication techniques for high performance anodes have been extensively investigated. However, Ni-YSZ cermet anodes suffer from performance degradation during long-term cell operation, which is associated with the agglomeration of nickel particles [1,2], the deactivation of nickel by fuel containing sulfur [3,4] or hydrocarbons [5,6], the microstructural changes due to the oxidation of nickel [7–9], and so on. Thus, along with the efforts to overcome the problems mentioned above, alternative anode materials with excellent long-term stability, especially in hydrocarbon fuel atmospheres, are being actively researched. Perovskite-type materials, typically doped strontium titanates (SrTiO3 ), are one of the most promising alternative anode materials for SOFCs because of their excellent structural stability in a reducing atmosphere. Doping with trivalent atoms, such as Y3+ and La3+ , in place of the divalent Sr2+ allows SrTiO3 to become an n-type semiconductor [10–13], while doping with Fe, Cr, Co, etc., results in a p-type semiconductor [14–17]. Although Y or La-doped

∗ Corresponding author. Tel.: +82 32 860 7521; +82 32 862 4482. E-mail address: [email protected] (H.J. Hwang). 0921-5107/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2011.10.016

SrTiO3 exhibits a high electrical conductivity under a reducing atmosphere, it appears that its area specific resistance (ASR) for the anode reaction is high, as compared with that of Ni-YSZ cermets [18]; this indicates that the catalytic activity for H2 oxidation of the Y or La-doped SrTiO3 needs to be improved. In this study, A and B-sites co-doped SrTiO3 -based anode, Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.05, 0.10, 0.20, 0.25, 0.40, and 0.50), was fabricated by a solid state reaction method in air. We selected yttrium and iron for A and B-site doping elements for high electrical conductivity and phase stability in air and H2 atmospheres. In addition, it is expected that the iron substitution for titanium in SrTiO3 can produce oxygen vacancies, which is due to the substitution of Fe3+ for Ti4+ . Oxygen vacancies may provide active sites for hydrogen oxidation in SOFC anode. The phase stability in air and in a reducing atmosphere as well as the electrical conductivity and electrochemical properties of the Y0.08 Sr0.92 Fex Ti1−x O3−ı anodes were investigated as a function of the Fe doping content.

2. Experimental procedure Y0.08 Sr0.92 Fex Ti1−x O3−ı powder was synthesized by a solid state reaction method from commercially available Y2 O3 (99.9%, Aldrich), SrCO3 (99.9%, Aldrich), Fe2 O3 (99.99%, Kojundo Kagaku Kougyo), and TiO2 (99.9%, Kojundo Kagaku Kougyo) powders. Appropriate amounts of the starting powders were mixed in a planetary ball mill for 2 h using ethyl alcohol and ZrO2 balls. The powder mixtures were dried with a rotary evaporator. The dried powder

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mixtures were calcined at 1400 ◦ C for 5 h in air. The obtained powders were then wet and dry-milled in a planetary ball mill. Rectangular bar (4 mm × 5 mm × 35 mm) samples were prepared from the Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.20 and 0.40) powders for the measurement of the electrical conductivity. The powders were pressed uniaxially in a WC-Co alloy mold, followed by cold isostatic pressing (CIP) at 200 MPa. The CIPed bar samples were fired at 1350 ◦ C for 5 h in air. The heating rate was 300 ◦ C h−1 . The sintered bar samples were polished using #220 to #2000-grit SiC papers. The electrical conductivity was measured in air and in H2 atmospheres by the four-probe method, in which a few microamperes were applied to the two outer electrodes of the rectangular bar sample and the resulting potential difference between the two inner electrodes was measured using a high accuracy multimeter (Model 2001, Keithley Instruments Inc.). Pt wires were used as the electrodes. The AC impedance spectra were obtained under open circuit conditions with an excitation potential of 20 mV over the frequency range of 1 MHz to 0.01 Hz by an impedance analyzer (IM6e, Zahner). The three Pt wires from the anode, cathode, and reference electrodes were connected to the working, counter, and reference terminals of the impedance analyzer, respectively. The electrochemical performance of the Y0.08 Sr0.92 Fex Ti1−x O3−ı anode was measured using an electrolyte-supported Pt/YSZ/Y0.08 Sr0.92 Fex Ti1−x O3−ı single cell. First, commercially available YSZ powder (8YSZ, Tosoh) was pressed into a pellet and sintered at 1400 ◦ C for 5 h in air. An YSZ disk with a diameter of 22 mm and a thickness of 0.3 mm was obtained. Y0.08 Sr0.92 Fex Ti1−x O3−ı anode paste was prepared by mixing the Y0.08 Sr0.92 Fex Ti1−x O3−ı powder and an organic vehicle (␣terpinol, n-butyl acetate, and ethyl cellulose). The weight ratio of Y0.08 Sr0.92 Fex Ti1−x O3−ı to the organic vehicle was 60:40. The anode layer was formed using the Y0.08 Sr0.92 Fex Ti1−x O3−ı paste on the YSZ disk via a screen printing technique followed by heat-treatment at 1200 ◦ C for 2 h in air. The area of the electrode was 1 cm2 . A current-collecting Pt layer containing 15 wt% YSZ was further formed on the Y0.08 Sr0.92 Fex Ti1−x O3−ı anode. A Pt-YSZ layer was formed as the cathode on the other side of the YSZ disk via a screen printing method. Pt wire was used to make the reference electrode on the side of the YSZ disk. The alumina tubes on both the top and bottom sides of the fuel cell were sealed with Pyrex glass rings. The fuel cell was then placed inside a furnace and heated at 900 ◦ C to allow for the sealing glass to be softened. H2 containing 3% H2 O and air were supplied to the anode and cathode sides, respectively. The flow rates of both H2 and air were 100 ml min−1 . For phase characterization, X-ray diffraction (XRD, RU-200B, Rigaku Co. Ltd.) was performed using Ni-filtered CuK␣ radiation. The thermal expansion coefficients were measured using a dilatometer (LINSEIS, L75/N1). Cylinder type samples with a diameter of 5 mm and a length of 20 mm were used for the thermal expansion coefficient measurement. The microstructure of the electrode was observed by field emission scanning electron microscopy (FESEM, HITACHI, S-4200).

3. Results and discussion Fig. 1 shows the XRD patterns of the powers obtained by calcining the mixture of SrCO3 , Y2 O3 , TiO2 , and different amounts of Fe2 O3 at 1400 ◦ C for 5 h. All of the peaks are assigned to a perovskite, which means that calcining at 1400 ◦ C is sufficient to synthesize a single phase perovskite with no unreacted or unwanted reaction phases. Since both SrTiO3 and SrFeO3 have a cubic perovskite structure and the ionic radius of the Fe ion (0.065 nm for six coordination) is similar to that of the Ti ion (0.061 nm for six coordination),

Fig. 1. XRD patterns of Y0.08 Sr0.92 Fex Ti1−x O3−ı heat-treated at 1400 ◦ C for 5 h in air.

the observed result is reasonable [19]. In addition, it appears that 8 mol% of Y doping does not negatively affect the formation of a solid solution between the two end members, although the ionic radius of yttrium (0.119 nm for twelve coordination) is much smaller than that of strontium (0.144 nm for twelve coordination) [20]. Fig. 2 shows the XRD patterns of the Y0.08 Sr0.92 Fex Ti1−x O3−ı powders after exposing to a 5% H2 /Ar atmosphere at 800 ◦ C (a) and 900 ◦ C (b) for 5 h. For comparison, the XRD pattern of the SrFe0.20 Ti0.80 O3−ı powder was also shown in Fig. 2(b). As is evident in Fig. 2(a), all of the Y0.08 Sr0.92 Fex Ti1−x O3−ı samples exhibited good phase stability at 800 ◦ C. On the other hand, the Y0.08 Sr0.92 Fex Ti1−x O3−ı perovskite was reduced at 900 ◦ C, most notably in the case of x = 0.6, and formed metallic Fe, Sr3 Ti2 O7 , and Y2 O3 . Once the Fe ions in the Y0.08 Sr0.92 Fex Ti1−x O3−ı were reduced completely and precipitated as metallic Fe, the excess TiO2 reacted with SrO, producing Sr3 Ti2 O7 and, at the same time, the Y, which was substituted for Sr in SrFex Ti1−x O3−ı , diffused out from the perovskite structure, as can be seen in Fig. 2(b). From Fig. 2(b) it appears that the Y doping allows the perovskite to be stable with respect to reduction under an H2 atmosphere. A similar result was reported by Fagg et al. [21] in a lanthanum and iron co-doped SrTiO3 system. They proved from XRD studies that the substitution of La for Sr is effective in improving the phase stability of the Fe doped SrTiO3 perovskite and that La substitution of more than 40% is required to keep the perovskite structure of SrTi0.60 Fe0.40 O3−ı . When a perovskite doped with a transition metal into the B sites, such as SrFex Ti1−x O3−ı , is heat-treated under a reducing atmosphere, oxygen vacancies are created and, at the same time, the transition metal would normally be reduced to a lower oxidation state for the sake of charge compensation. In the case of SrFex Ti1−x O3−ı , it appears that Fe3+ is reduced to Fe0 , i.e., metallic Fe at 800 ◦ C in an H2 atmosphere. Since YFeO3 is known to be much more durable in a reducing atmosphere than SrFeO3 [22], it is considered that the Y doped SrFex Ti1−x O3−ı exhibited better reduction stability than the un-doped SrFex Ti1−x O3−ı . Fig. 3 shows the thermal expansion versus temperature curves of the Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.05, 0.25, and 0.50) perovskites measured in air. Thermal expansion coefficients (TEC) are also shown in Fig. 3. The curve shows a linear relationship for the x = 0.05 sample. On the other hand, the inflection point of the curve was observed in the x = 0.25 and 0.50 samples at around 400 ◦ C, and the slope of the expansion curve slightly increased with respect to the Fe doping content. The slope change, i.e., the increase in the

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Fig. 2. XRD patterns of Y0.08 Sr0.92 Fex Ti1−x O3−ı reduced at 800 ◦ C (a) and 900 ◦ C (b) for 5 h in 5% H2 /Ar.

thermal expansion coefficient, observed in the x = 0.25 and 0.50 samples is thought to be associated with the oxygen vacancy formation and larger ionic radii which is due to the reduction of Fe or Ti in the perovskite structure [23]. The thermal expansion coefficients of the x = 0.05 and x = 0.25 samples were estimated to be 11.7 × 10−6 and 12.5 × 10−6 ◦ C−1 in the temperature range of 100 to 1000 ◦ C. These values are almost identical to Ni-YSZ cermet, Gd-doped ceria (GDC), and ferritic stainless steel, however, they are slightly larger than that of 8YSZ (10.7 × 10−6 ◦ C−1 ). On the other hand, the TEC of the x = 0.50 sample was much higher than that of YSZ and it was found to be 13.3 × 10−6 ◦ C−1 and 17.4 × 10−6 ◦ C−1 for 100 to 400 ◦ C and 400 to 1000 ◦ C, respectively. Fig. 4 presents the electrical conductivity of the Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.20 and 0.40) perovskites measured in air and H2 atmospheres as a function of temperature. Irrespective of the Fe content, the conductivity of Y0.08 Sr0.92 Fex Ti1−x O3−ı increased as the temperature increased and it was higher in air than in H2 in the low and intermediate temperature region, suggesting that the p-type conductivity coming from the substitution of Ti4+ for Fe3+ contribute more effectively to the conductivity in air. In air, the conductivity of Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.40) was higher than that of the x = 0.20 sample. It appears that the higher charge carrier (electron–hole) concentration is responsible for the observed high conductivity in Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.40) [15]. The conductivity of the Y0.08 Sr0.92 Fex Ti1−x O3−ı perovskite increased as the temperature increased in the low temperature range and thereafter became saturated. This conduction behavior could be explained by two factors. The first is the predominant 0.020

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ionic compensation. In other words, the loss of oxygen in the lattice causes a decrease in the concentration of electron holes as the concentration of oxygen vacancies increases at high temperatures [24]. The second is the reduction in the hole mobility at high temperatures; the hole mobility was blocked by oxygen vacancies resulting from the loss of oxygen at high temperatures [25–27]. However, this conductivity break at high temperatures was modest in Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.20), suggesting that it has a relatively lower concentration of oxygen vacancies. The temperature anomalies and slope change in the thermal expansion coefficient shown in Fig. 3 correspond well with the conductivity behavior in air. This indicates that oxygen vacancy formation occurred around 400 ◦ C due to the substitution by Fe and it was accelerated as the Fe content increased. As is similar to that in air, the conductivity in H2 was higher for the x = 0.40 sample than for the x = 0.20 sample. However, the difference between the two was small and the conductivity was found to be almost the same at 800 ◦ C. At a low oxygen partial pressure, Y0.08 Sr0.92 Fex Ti1−x O3−ı exhibits n-type semiconductivity and its electrical conductivity depends on both the Y content (donor concentration) and the reduction Ti or Fe. Since Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.20 and 0.40) have the same donor concentration and part of electrons generated by the donor may be compensated by electron holes by the acceptor, the electrical conduction of Y0.08 Sr0.92 Fex Ti1−x O3−ı is governed by the reduction of Ti4+ to Ti3+ , the electron concentration (electronic conduction), which originates from oxygen losses in H2 , and the oxygen vacancy concentration (ionic conduction). It is considered that oxygen vacancy formation and the resulting electronic conduction in H2 is more preferable in a lightly substituted SrTiO3 , such as Y0.08 Sr0.92 Fe0.2 Ti0.8 O3−ı , than in a highly substituted SrTiO3 . Thus, the activation energy of the x = 0.20 sample in H2 was high (1.02 eV) and its conductivity was almost the same as that of the x = 0.40 sample at 800 ◦ C. On the other hand, it seems that the contribution of the oxygen ion conductivity to the total conductivity is larger in the x = 0.40 sample than in the x = 0.20 sample, especially at low temperatures because the heavily substituted SrTiO3 limits further oxygen vacancy formation in Y0.08 Sr0.92 Fex Ti1−x O3−ı in H2 . Therefore, the activation energy of Y0.08 Sr0.92 Fex Ti1−x O3−ı with x = 0.40 was low (0.58 eV) compared with that of the x = 0.20 sample. The SEM images of the cross sections of the Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.20, 0.25, 0.40, and 0.50)/YSZ interfaces sintered at 1200 ◦ C for 2 h are shown in Fig. 5. In order to make high performance anodes, it is crucial to fabricate Y0.08 Sr0.92 Fex Ti1−x O3−ı particles that not only have a fine microstructure, i.e., a small grain size, but also have strong adhesion to the YSZ electrolyte. As is evident in Fig. 5, the Y0.08 Sr0.92 Fex Ti1−x O3−ı anode exhibited good interfacial contact

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Fig. 4. Electrical conductivity of Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.20 and 0.40) as a function of temperature.

Fig. 5. SEM images showing the cross section of the Y0.08 Sr0.92 Fex Ti1−x O3−ı /YSZ interface sintered at 1200 ◦ C for 2 h; (a) x = 0.20, (b) x = 0.25, (c) x = 0.40, (d) x = 0.50.

with the YSZ electrolyte and there was no sign of delamination or cracking. On the other hand, the microstructure of the Y0.08 Sr0.92 Fex Ti1−x O3−ı anode appears to depend on the Fe content; the x = 0.50 anode exhibited a greater grain size and denser microstructure than the x = 0.20 anode sample. These results mean that a sintering temperature of 1200 ◦ C is too high to obtain the x = 0.50 anode with a fine microstructure. In addition, Sr(Ti,Zr)O3 was found at the interface of Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.50) and YSZ, which may result from the reaction after the interdiffusion of Zr from the YSZ and Sr or Ti from the oxide anode (Fig. 6). The fact that the reaction product is observed in the Y0.08 Sr0.92 Fex Ti1−x O3−ı with a higher Fe content suggests that the phase stability between Y0.08 Sr0.92 Fex Ti1−x O3−ı and YSZ decreases as the Fe content increases. Fig. 7 shows the AC impedance spectra of the Y0.08 Sr0.92 Fe0.40 Ti0.60 O3−ı anode measured at 700, 750, 800 and 850 ◦ C in an H2 atmosphere. The impedance spectra in Fig. 7 were taken at 0 mA, i.e., an open circuit voltage (OCV) condition. The intercept of the impedance semicircle with the real axis at high frequencies (the left intercept of the semicircle) corresponds to the ohmic resistance, including the bulk resistances of the electrolyte and anode. On the other hand, the interfacial resistance

Fig. 6. Thin film XRD patterns for the YSZ electrolyte surface of the Y0.08 Sr0.92 Fe0.5 Ti0.5 O3−ı /YSZ interface sintered at 1200 ◦ C for 5 h in air. The perovskite anode was removed from the cell.

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850 800 750 700

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Z' ( Ω cm ) Fig. 7. AC impedance spectra of the Y0.08 Sr0.92 Fe0.40 Ti0.60 O3−ı /YSZ interface at 700 ◦ C, 750 ◦ C, 800 ◦ C and 850 ◦ C. The lower figure is the enlargement of the upper one.

is determined by subtracting the ohmic resistance from the total resistance, i.e., the intercept of the semicircle with the real axis at low frequencies (right intercept of the semicircle). As is evident in Fig. 7, the impedance spectra of the Y0.08 Sr0.92 Fe0.40 Ti0.60 O3−ı anode consist of two or three semicircles, depending on the temperature. At 850 ◦ C, the spectrum exhibits a small semicircle in the high-frequency (1 × 103 Hz) region, followed two large semicircles in the medium (44 Hz) and lowfrequency (4.7 × 10−2 Hz) region. As the temperature decreases, the semicircles in the high and medium-frequency region grow. In contrast, the semicircle in the low-frequency region become small at 750 ◦ C and disappeared at 700 ◦ C. These results suggest that the semicircle in the high and medium-frequency regions correspond to electrode reactions such as charge transfer or dissociative adsorption of H2 gas. The semicircle observed in the low-frequency region is associated with the depletion of H2 in the triple phase boundary of the anode at high temperatures, i.e., concentration polarization. At 750 and 700 ◦ C, it is thought that the transport of H2 through the anode is sufficiently fast because of relatively low electrode reaction kinetics. Fig. 8 shows the AC impedance spectra of the Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.20, 0.25, 0.40, 0.50) anodes measured at 700 ◦ C in an H2 atmosphere. For comparison, the impedance spectrum of a Ni-YSZ anode prepared with home equipment is also displayed in Fig. 8. The interfacial resistance depended on the Fe content of the Y0.08 Sr0.92 Fex Ti1−x O3−ı anode; it was 12.5  cm2 at x = 0.20, it then decreased to 1.8  cm2 as the Fe content increased to x = 0.25, and then it increased again to 6.9  cm2 (x = 0.40) and 67.1  cm2 (x = 0.50). The Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.25) anode exhibited the lowest polarization resistance (0.7 and 1.8  cm2 at 800 and 700 ◦ C). They are comparable to those of the Ni-YSZ cermet anodes prepared in this study (1.0 and 3.1  cm2 at 800 and 700 ◦ C) or presented by other researchers [28,29]. On the other hand, the large grain size, poorer redox stability, and the reactivity with the YSZ electrolyte are responsible for the high polarization resistance observed in the Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.40 and 0.50) anode.

0 2

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Z' (Ω cm ) Fig. 8. AC impedance spectra of the Y0.08 Sr0.92 Fex Ti1−x O3−ı /YSZ interface at 700 ◦ C as a function of Fe doping content.

In terms of the ohmic resistance, the Y0.08 Sr0.92 Fex Ti1−x O3−ı anodes with x = 0.20, 0.25, and 0.40 exhibited similar values in the range of 2.8 to 4.0  cm2 , which is higher than that of the Ni-YSZ cermet anode. It appears that the high ohmic resistance in the Y0.08 Sr0.92 Fex Ti1−x O3−ı anodes is due to their low electrical conductivity. As can be seen in Fig. 4, the conductivities of dense Y0.08 Sr0.92 Fex Ti1−x O3−ı sintered bodies are 0.09 and 0.15 S cm−1 at 700 ◦ C for x = 0.20 and x = 0.40 samples, respectively. Since the Y0.08 Sr0.92 Fex Ti1−x O3−ı anode has a porous microstructure, the conductivity should decrease, and thus the resistance due to the anode cannot be ignored. On the other hand, the high ohmic resistances in the Y0.08 Sr0.92 Fex Ti1−x O3−ı anodes, however, cannot be explained only by the resistance of the anode. Another possible reason lies on the contact resistance, which is caused by current collection loss due to the electrode resistance. In addition, the high ohmic resistance in the Y0.08 Sr0.92 Fex Ti1−x O3−ı anodes with x = 0.50 (16.9  cm2 ) might be associated with the insulating phase, Sr(Ti,Zr)O3 , which was formed at the Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.50)/YSZ interface. 4. Conclusions An Y0.08 Sr0.92 Fex Ti1−x O3−ı perovskite was synthesized by calcinating a Y2 O3 , SrCO3 , Fe2 O3 , and TiO2 powder mixture at 1400 ◦ C in air. The Y0.08 Sr0.92 Fex Ti1−x O3−ı perovskite exhibited good perovskite phase stability under a reducing atmosphere at 800 ◦ C, which is thought to be due to the substitution of Y for Sr. It was thought that the positive charges induced by the Y substitution may suppress the formation of oxygen vacancies and the resulting reduction of Fe. In addition, Y itself can improve the dissociation pressure of the perovskite. In air, the conductivity of the Y0.08 Sr0.92 Fex Ti1−x O3−ı perovskite was determined by the acceptor concentration, i.e., the Fe content. The greater the Fe content, the higher the conductivity. In contrast, it was less sensitive to the acceptor concentration in H2 , especially at high temperatures, because the contribution of the high carrier concentration was compensated by the formation of oxygen vacancies. For the x = 0.20

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and 0.25 samples, sintering at 1200 ◦ C is optimum to obtain a fine microstructure and good interfacial adhesion. For the x = 0.40 and above samples, low temperature sintering or the presence of a grain growth inhibitor would be required to obtain the optimum microstructure, since the sinterability of Y0.08 Sr0.92 Fex Ti1−x O3−ı increased as the Fe content increased. The polarization resistances of the Y0.08 Sr0.92 Fex Ti1−x O3−ı (x = 0.25) anode were found to be 0.7 and 1.8  cm2 at 800 and 700 ◦ C, respectively. Acknowledgments This research was supported by a grant from the Fundamental R&D Program for Core Technology of Materials funded by the Ministry of Knowledge Economy, Republic of Korea. Part of this work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korea government (MEST) (No.2010-0010744). References [1] D. Simwonis, F. Tietz, D. Stover, Solid State Ionics 132 (2000) 241–251. [2] H. Koide, Y. Someya, T. Yoshida, T. Maruyama, Solid State Ionics 132 (2000) 253–260. [3] A. Lussier, S. Sofie, J. Dvorak, Y.U. Idzerda, Int. J. Hydrogen Energy 33 (2008) 3945–3951. [4] R. Mukundan, E.L. Brosha, F.H. Garzon, Electrochem. Solid-State Lett. 7 (2004) A5–A7. [5] V.A. Restrepo, J.M. Hill, Appl. Catal. A: Gen. 342 (2008) 49–55. [6] J.H. Koh, Y.S. Yoo, J.W. Park, H.C. Lim, Solid State Ionics 149 (2002) 157–166.

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