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Characterization of Sr2.7Ln0.3Fe1.4Co0.6O7 (Ln = La, Nd, Sm, Gd) intergrowth oxides as cathodes for solid oxide fuel cells
Solid State Ionics 180 (2009) 1478–1483
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Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s s i
Characterization of Sr2.7Ln0.3Fe1.4Co0.6O7 (Ln = La, Nd, Sm, Gd) intergrowth oxides as cathodes for solid oxide fuel cells J.-H. Kim, A. Manthiram ⁎ Electrochemical Energy Laboratory & Materials Science and Engineering Program, University of Texas at Austin, Austin, Texas 78712, USA
a r t i c l e
i n f o
Article history: Received 19 June 2009 Received in revised form 18 September 2009 Accepted 21 September 2009 Keywords: Solid oxide fuel cells Cathodes Intergrowth oxides
a b s t r a c t Perovskite-related intergrowth oxides Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) have been investigated as cathode materials for solid oxide fuel cells (SOFC). With decreasing size of the Ln3+ ions, the unit cell volume, oxygen content, thermal expansion coefficient (TEC), and total electrical conductivity decrease from Ln = La to Gd. The decreasing unit cell volume and oxygen content is attributed to the decreasing size of Ln3+ ions from Ln = La to Gd and a consequent preference for lower coordination numbers. While the decrease in the ionicity of the Ln–O bonds from Ln = La to Gd causes a decrease in the TEC, the increasing amount of oxygen vacancies leads to a decrease in electrical conductivity arising from a 1 thermally activated semiconducting behavior. The cathode polarization conductance (R− p ) measured using the ac-impedance spectroscopy and the catalytic activity for the oxygen reduction reaction in SOFC decrease from Ln = La to Gd partly due to the decrease in electrical conductivity. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Transition metal oxides with mixed oxide-ion and electronic conducting (MIEC) properties find unique applications such as cathode materials in SOFC. The cathode in SOFC requires both high electronic and oxide-ion conductivities, thermal expansion compatibility with the electrolyte, good chemical stability with other cell components, and good catalytic activity for the oxygen reduction reaction at high temperatures [1]. Perovskite (La,Sr)MnO3 has traditionally been studied extensively due to its advantage of good thermal expansion coefficient match with other cell components and good chemical stability. However, the low oxide-ion conductivity and consequent low catalytic activity at intermediate-temperatures make (La,Sr)MnO3 unattractive for operation at T < 800 °C [2]. These difficulties have created enormous interest in the development of alternative cathode materials. One class of materials that have drawn some attention is the perovskite-related intergrowth oxides belonging to the Ruddlesden– Popper (R–P) series [3] having the general formula of (AO)(ABO3)n with n = 1, 2, 3. These intergrowth oxides have rock-salt AO layers alternating with a single (n = 1), double (n = 2), or triple (n = 3) perovskite (ABO3)n layers along the c axis [4,5]. For example, Fig. 1 shows the crystal structure of (Sr,Ln)3(Fe,Co)2O7 (n = 2) having (Sr,Ln)O rock-salt layers alternating with two (Sr,Ln)(Fe,Co)O3 perovskite layers along the
⁎ Corresponding author. Tel.: +1 512 471 1791; fax: +1 512 471 7681. E-mail address: [email protected] (A. Manthiram). 0167-2738/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2009.09.007
c axis. Various compositions in this R–P series have been reported with A = Sr and La, and B = Mn, Fe, Co, Ni, Cu, and Ga [4–17]. The MIEC properties of the (Sr,La)n + 1(Fe,Co)nO3n + 1 (n = 1–3) oxides as well as their catalytic activity in SOFC have been studied by our group before and compared with those of perovskite oxides [5–8,17]. The total electrical conductivity and the oxygen permeation flux of the (Sr,La)n + 1(Fe,Co)nO3n + 1 phases increase with n and reach a maximum value for the SrCo0.8Fe0.2O3 − δ perovskite, which can be considered as the n = ∞ member [5,8,9]. However, the R–P phases generally show better structural stability without undergoing undesired phase transitions compared to the perovskite oxides, which tend to lose too much oxygen at high temperatures and undergo a perovskite to brownmillerite transition particularly under reducing conditions. In the n = 2 series, Sr3Fe2O7 − δ suffers from structural and chemical instability at room temperature due to a slow reaction with H2O and CO2 in air. Although our group was able to overcome this instability by a substitution of 10% La for Sr, such a substitution led to a decrease in oxide-ion conductivity [8]. Interestingly, a substitution of 30% Co for Fe to give Sr2.7La0.3Fe1.4Co0.6O7 − δ improved both the electrical and oxideion conductivities and consequently the catalytic activity for ORR in SOFC. Unfortunately, the substitution of Co increases the thermal expansion coefficient (TEC) due to the low spin to the high spin transitions of Co3+ ions, which can lead to mechanical integrity problems in SOFC in contact with the electrolyte. One way to reduce TEC is to replace the La3+ ions in Sr2.7La0.3Fe1.4Co0.6O7 − δ by other Ln3+ (Ln = lanthanide) ions as the iconicity and the thermal expansion of the Ln–O bonds decrease with decreasing size of the Ln3+ ions. Accordingly, we present here an investigation of the high temperature properties of
J.-H. Kim, A. Manthiram / Solid State Ionics 180 (2009) 1478–1483
Fig. 1. Crystal structure of the intergrowth oxide (Sr,Ln)3(Fe,Co)2O7 − δ.
Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) and an exploration of their use as cathodes in SOFC. The influences of the Ln3+ ions on the crystal structure, structural stability, thermal properties, and catalytic activity for the oxygen reduction reaction in SOFC are discussed.
2. Experimental The Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) oxides were synthesized by conventional solid-state reaction. Required amounts of the lanthanide oxides (La2O3, Nd2O3, Sm2O3, or Gd2O3), SrCO3, Fe2O3, and Co3O4 were thoroughly mixed with ball milling in ethanol for 24 h and calcined at 1000 °C for 12 h in air. The calcined powders were then ground, pressed into pellets, and sintered at 1300 °C for 24 h in air. The resulting products were ground and finally heated at 900 °C for 6 h in air, followed by slow cooling to room temperature at a rate of 1 °C/min to maximize the oxygen content value. The La0.8Sr0.2 Ga0.8 Mg0.2O2.8 (LSGM) electrolyte disks were prepared by firing required amounts of La2O3, SrCO3, Ga2O3, and MgO at 1100 °C for 5 h, followed by pelletizing and sintering at 1500 °C for 10 h. Ce0.8Gd0.2O1.9 (GDC), Ce0.9Gd0.1O1.95 (10GDC), and La0.4Ce0.6O1.8 (LDC) powders were synthesized by the glycine–nitrate combustion method [18,19]. For the anode, NiO and 10GDC (Ni:10GDC = 70:30 vol.%) were ball-milled in ethanol for 48 h. The products thus obtained were characterized by X-ray diffraction (XRD) and the XRD data were refined with the Rietveld method using the Fullprof program [20]. Due to the instability problem in air, the Sr2.7Gd0.3Fe1.4Co0.6O7 − δ sample for the XRD experiment was prepared in argon atmosphere and sealed with a kapton film. The average oxidation state of Co and Fe and the room-temperature oxygen content values were determined by iodometric titration [21]. Thermogravimetric analysis (TGA) and thermal expansion data were collected with a Perkin-Elmer Series 7 thermal analysis system. The TGA experiments were carried out with a heating/cooling rate of 3 °C/min in air. The TECs of the sintered samples were measured in the temperature range of 80–900 °C with a heating/cooling rate of 5 °C/min with an intermediate dwelling at 900 °C for 0.5 h. Electrical conductivity of the pellets was measured with a four-probe dc method using a Van der Pauw configuration in the temperature range of 40–900 °C [22]. The specimens for the reactivity tests were obtained by mixing the
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Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) oxides and the GDC powders, followed by heating at 1000 °C for 3 h in air. The polarization resistance (Rp) of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathodes were measured with a half cell consisting of cathode|GDC|Pt. The Pt meshes and wires were attached to each electrode using Pt paste as the current collector. The Pt reference electrode was placed at a distance of 5 mm from the cathode. The acimpedance spectra of the half cells were obtained at different temperatures with the Solartron 1260 FRA equipment. Fuel cell performances of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathodes were evaluated with electrolyte-supported single cells. All the electrode materials were mixed with an organic binder (Heraeus V006) to form slurries and then applied onto the surface of a dense LSGM pellet (400 μm thickness) by screen printing. To prevent the formation of LaNiO3 or La2NiO4 at the anode|electrolyte interface, a LDC buffer layer was first prepared by screen printing onto the anode side of the LSGM electrolyte [23]. After heating the buffer layer at 1000 °C for 1 h, the NiO–10GDC cermet anode was screen printed onto the LDC layer and heated at 1300 °C for 0.5 h. Similarly, to prevent the side reactions at the cathode|electrolyte interface, a GDC buffer layer was prepared onto the cathode side of the LSGM electrolyte and heated at 1200 °C for 3 h. The Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathode layers were subsequently screen printed onto the GDC buffer layer and heated at 1000 °C for 3 h. Pt meshes and wires were attached to each electrode with an area of 0.25 cm2 using Pt paste as the current collector. During the SOFC operation, humidified H2 (~3% H2O at 25 °C) and air were supplied as fuel and oxidant, respectively, at a rate of 100 cm3/min. The cross sections of the single cells were observed with a scanning electron microscope (SEM, JEOL JSM-5610) after the SOFC performance tests. 3. Results and discussion The room-temperature crystal structures of the Sr2.7Ln0.3Fe1.4Co0.6 O7 − δ (Ln = La, Nd, Sm, Gd) samples were refined based on the tetragonal space group I4/mmm. The experimental XRD data, calculated profiles, and the difference between the two shown in Fig. 2 reveal good agreement between the experimental and calculated profiles. Tables 1 and 2 give the refined lattice parameters, atomic positions, occupancies of Sr and Ln, and the quality of the refinements. The lattice parameters a and b and the unit cell volume decrease from Ln= La to Gd due to the decreasing size of Ln3+. In addition, the smaller Ln3+ ions compared to the larger Sr2+ ions prefer the Ln(1) sites in the perovskite layer rather than the Ln(2) sites in the rock-salt layer in Fig. 1. For example, the occupancies of Ln at the Ln(1) site and Sr at the Sr(2) site increase as the size of the Ln3+ ions decreases from La3+ to Gd3+ as seen in Table 2. The average oxidation state values of the transition metal ions and the oxygen contents determined by the iodometric titration are given in Table 1. The room-temperature oxygen content decreases from 7.0 for Ln= La to 6.91 for Ln= Gd. It has been reported that oxygen vacancies are exclusively present on the O(1) site in the Sr(1)/Ln(1)–O(1) plane of Sr3(Fe,Co)2O7 [24,25]. Recent high-temperature neutron diffraction study also revealed that oxygen vacancies in Sr3Fe2O7 − δ are predominantly present at the O(1) site in the temperature range of 20≤T ≤ 900 °C [4]. Thus, the increase in the Ln occupancy at the Ln(1) site and the oxygen vacancies at the O(1) site on going from Ln= La to Gd could be understood to be due to the preference of the Ln3+ ions for lower coordination number with decreasing Ln3+ size. From an earlier study [8], the Sr3Fe2 − yCoyO7 composition was found to suffer from chemical and structural instability at room temperature due to a slow reaction with H2O and CO2 in air. In this study, a substitution of 10% La, Nd, or Sm for Sr in Sr2.7Ln0.3Fe1.4Co0.6O7 − δ was found to be beneficial for suppressing the instability in air. However, the substitution of Gd for Sr could not help to overcome the instability problem in air. For example, the XRD patterns of the Sr2.7Gd0.3Fe1.4 Co0.6O7 − δ powder after storing in air and vacuum indicate that while powder exposed to open air exhibits structural degradation as
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J.-H. Kim, A. Manthiram / Solid State Ionics 180 (2009) 1478–1483 Table 2 Crystallographic data of Sr2.7Ln0.3Fe1.4Co0.6O7 − δ at room temperature.a Ln
La
Nd
Sm
Gd
zSr2/Ln2 zFe/Co zO2 zO3 OccupancySr1 OccupancyLn1 OccupancySr2 OccupancyLn2 Rp Rwp χ2
0.317 0.098 0.198 0.091 0.89 0.11 1.81 0.19 6.01 7.92 2.10
0.318 0.098 0.196 0.090 0.81 0.19 1.89 0.11 5.61 7.49 1.56
0.318 0.099 0.199(1) 0.086(1) 0.76 0.24 1.94 0.06 7.47 9.51 2.69
0.317(1) 0.098(1) 0.196 0.089 0.75 0.25 1.95 0.05 9.15 12.3 1.30
a Refinements were performed based on the space group I4/mmm with the atomic positions Sr(1)/Ln(1): 0, 0, 0.5; Sr(2)/Ln(2): 0, 0, z; Fe/Co: 0, 0, z; O(1): 0, 0, 0; O(2): 0, 0, z; O(3): 0, 0.5, z.
Fig. 2. XRD patterns, calculated profiles, peak positions, and the difference between the observed and calculated profiles of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ samples: (a) Ln = La, (b) Ln = Nd, (c) Ln = Sm, and (d) Ln = Gd.
evidenced by the impurity reflections in Fig. 3(a), that stored in vacuum maintains the original intergrowth structure without any degradation as seen in Fig. 3(b). Fig. 4 shows the variations of the oxygen content and oxidation state of Fe and Co ions with temperature in air for the various Sr2.7Ln0.3Fe1.4 Co0.6O7 − δ (Ln = La, Nd, Sm, Gd) samples. These curves were derived using the initial oxygen content values (Table 1) determined by the iodometric titration and the TGA data. The Sr2.7Ln0.3Fe1.4Co0.6O7 − δ samples lose 0.4–0.45 oxygen atoms per formula unit on heating to
920 °C. At a given temperature, the oxygen content and the average oxidation state of Fe and Co ions decrease from Ln= La to Gd due to the differences in the initial, room-temperature oxygen content values (Table 1). Fig. 5 shows thermal expansion curves of the Sr2.7Ln0.3Fe1.4Co0.6O7 −δ (Ln = La, Nd, Sm, Gd) specimens measured in air. All the specimens show a linear thermal expansion behavior as indicated by a dashed line at T < 300 °C, where the oxygen loss is negligible (Fig. 4). However, they show an increase in the slope at T > 300 °C due to the loss of oxygen from the lattice (Fig. 4) and the consequent reduction of smaller Fe4+/ Co4+ to larger Fe3+/Co3+ ions. The thermal expansion curves were found to be reversible in the subsequent heating and cooling cycles and the average TEC values are given in Table 1 for two different temperature ranges, 80–900 °C and 500–900 °C. The TEC value decreases from Ln= La to Gd in both the temperature ranges. The decrease in TEC is due to the decrease in the ionicity of the Ln–O bonds as the ionic bonds generally exhibit larger thermal expansion than covalent bonds. Similar trends with the Ln–O bond ionicity have been observed in the (La,A)MnO3 [26] and LnBaCo2O5 + δ [27] systems. Fig. 6 shows the variations of the electrical conductivity with temperature for the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, Gd) samples. All the samples show an increase in conductivity with temperature for T < 500 °C, indicating a thermally activated semiconducting behavior. However, the conductivity decreases with increasing temperature for T > 500 °C due to the increasing concentration of oxygen vacancies (Fig. 4) and a consequent perturbation of the O–(Fe, Co)–O interaction and carrier localization [28]. For the same reason, at a given temperature, the electrical conductivity decreases from Ln= La to Gd due to the decreasing oxygen content and carrier concentration. The Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, Gd) samples show much lower electrical conductivities compared to the perovskite oxides such as La0.6Sr0.4CoO3 − δ (~1200 S/cm) [29]. While the perovskite structure facilitates three-dimensional electrical pathways through the O–Co–O network, the double (Sr,Ln)O rock-salt layers in between the double perovskite layers along the c axis in the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ samples breaks down the O–(Fe,Co)–O interaction along the c axis (Fig. 1) and limits the electrical conduction pathway. For this reason, the electrical
Table 1 Structural parameters, chemical analysis data, and TECs of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ samples. Sample
La0.3Sr2.7Fe1.4Co0.6O7 − δ Nd0.3Sr2.7Fe1.4Co0.6O7 − δ Sm0.3Sr2.7Fe1.4Co0.6O7 − δ Gd0.3Sr2.7Fe1.4Co0.6O7 − δ
Space group
I4/mmm I4/mmm I4/mmm I4/mmm
a (Å)
3.860 3.857 3.857 3.849
c (Å)
20.171 20.120 20.103 20.075
V (Å3)
300.577 299.270 299.084 297.417
Oxidation state of Fe,Co
Oxygen content (7 − δ)
TEC × 106 (°C− 1) 80–900 °C
500–900 °C
3.85 3.84 3.78 3.76
7.00 6.99 6.93 6.91
18.2 18.0 18.0 17.3
21.2 20.2 20.0 19.1
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Fig. 3. XRD patterns of the Sr2.7Gd0.3Fe1.4Co0.6O6.91 powder stored in (a) air and (b) vacuum. The vertical bars show the peak positions of phase-pure Sr2.7Gd0.3Fe1.4Co0.6O6.91.
Fig. 6. (a) Temperature dependence of the electrical conductivity in the temperature range of 40–900 °C and (b) Arrhenius plots of the electrical conductivity in the temperature range of 40–280 °C for the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) samples in air.
plots show the Arrhenius relation as seen in Fig. 6(b), indicating small polaron hopping conduction, which is generally expressed as follows [30], Fig. 4. Variations of the oxygen content and the average oxidation state of (Fe,Co) in Sr2.7Ln0.3Fe1.4Co0.6O7 − δ with temperatures in air: (a) Ln = La, (b) Ln = Nd, (c) Ln = Sm, and (d) Ln = Gd.
conductivity of the (Sr,La)n + 1(Fe,Co)nO3n + 1 samples increases with n and reaches the highest value for the perovskite (n = ∞ member) [8,9]. At T < 300 °C where oxygen loss is negligible (Fig. 4), the log (σT) vs. 1/T
Fig. 5. Thermal expansion (dL/Lo) curves of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ specimens in the temperature range of 80–900 °C in air: (a) Ln = La, (b) Ln = Nd, (c) Ln = Sm, and (d) Ln = Gd.
σ=
A E exp − a kT T
ð1Þ
where k is the Boltzmann constant, T is temperature, Ea is activation energy for the hopping of the small polarons, and A is the preexponential factor. The Ea for the small polaron hopping conduction
Fig. 7. XRD patterns of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ and GDC mixtures after firing at 1000 °C for 3 h in air: (a) Ln = La, (b) Ln = Nd, (c) Ln = Sm, and (d) Ln = Gd.
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1 Fig. 8. (a) Temperature dependence of the total polarization conductance (R− p ) and (b) the normalized ac-impedance spectra measured at 550 °C of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathodes in air.
increases from 0.14 eV for Ln = La to 0.17 eV for Ln = Gd in the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ system. We have reported previously that a GDC interlayer is necessary in between the (Sr,La)3(Fe,Co)2O7 cathode and LSGM electrolyte to prevent their side reaction at high temperature during the single cell fabrication [17]. Therefore, the chemical stability of Sr2.7Ln0.3Fe1.4Co0.6 O7 − δ (Ln = La, Nd, Sm, Gd) in contact with GDC was assessed by heating the mixed powders at 1000 °C for 3 h. Fig. 7 compares the XRD patterns recorded after heat treating the mixed powders. While the Ln= La sample shows a trace amount of unknown reaction product as indicated
by the reflection at 2θ ≈ 31.7°, the amount of this reaction product decreases or becomes negligible for the Ln= Nd, Sm, and Gd samples. Fig. 8(a) compares the temperature dependence of the total 1 polarization conductance (R− p ) of the Sr 2.7Ln0.3Fe 1.4Co 0.6 O7 − δ (Ln = La, Nd, Sm, Gd) cathodes on the GDC electrolyte in air. The activation energy (Ea) of the cathode polarization conductance increases with decreasing size of the Ln3+ ions from La (Ea = 1.51 eV) to Gd (Ea = 1.68 eV). In addition, at a given temperature, the cathode polarization conductance decreases from Ln = La to Gd. Fig. 8(b) compares the normalized impedance spectra of the Sr2.7Ln0.3Fe1.4Co0.6 O7 − δ (Ln = La, Nd, Sm, Gd) cathodes at 550 °C in air. The increasing resistance and Ea for polarization conductance on going from Ln= La to Gd is partly attributed to the decreasing electrical conductivity and the increasing Ea for the small polaron hopping conduction (Fig. 6). The electrochemical performance of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathodes in SOFC for the oxygen reduction reaction was evaluated using LSGM as an electrolyte. The GDC interlayer is employed between the cathode and electrolyte to prevent side reactions between the cathode and the LSGM electrolyte. Fig. 9 shows the SEM images of the cross sections of the cathode|GDC interlayer| LSGM electrolyte portion after the SOFC single cell performance tests. While the bottom of the micrograph indicates a dense, well-sintered LSGM electrolyte, the upper portion shows the GDC interlayer and the porous Sr2.7Ln0.3Fe1.4Co0.6O7 − δ cathode. The GDC interlayer with a thickness of ~ 5 μm provides a good separation between the cathode and the electrolyte. Fig. 10 compares the current–voltage (I–V) curves and the corresponding power density curves measured at 800 °C for the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathodes. The single cells show maximum power densities (Pmax) of 469, 447, 430, and 299 mW/cm2, respectively, for the Ln = La, Nd, Sm, and Gd cathodes. Thus, the cathode performance decreases from Ln = La to Gd in accordance with the cathode polarization resistance data in Fig. 8. The low electrical conductivity (≤80 S/cm) as seen in Fig. 6 can limit the charge-transfer rate in the overall oxygen reduction process with the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ cathode at 800 °C. Accordingly, the decreasing electrical conductivity from Ln= La to Gd leads to a decrease in the cathode performance. However, the Ln= Gd cathode exhibits much
Fig. 9. SEM micrographs showing the cross sections of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ cathode|GDC interlayer|LSGM electrolyte portion after the SOFC single cell performance tests: (a) Ln = La, (b) Ln= Nd, (c) Ln= Sm, and (d) Ln= Gd.
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and the catalytic activity for the oxygen reduction reaction in SOFC decrease from Ln= La to Gd partly due to the decreasing electrical conductivity. Acknowledgement Financial support by the Welch Foundation grant F-1254 is gratefully acknowledged. References
Fig. 10. SOFC performance data of single cells consisting of Sr2.7Ln0.3Fe1.4Co0.6O7 − δ| GDC|LSGM|LDC|Ni-GDC at 800 °C.
inferior performance in SOFC compared to the Ln= La, Nd, and Sm cathodes. This can be understood by the structural instability of the Ln= Gd cathode in air (Fig. 3). Although the vacuum-stored powders were employed to prepare the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ cathodes, the Ln= Gd cathode experiences structural instability problems in contact with air during the single cell fabrication process. 4. Conclusions The influence of the size of Ln3+ ions on the crystal structure, oxygen content, thermal expansion, electrical conductivity, and catalytic activity for the oxygen reduction reaction in SOFC of Sr2.7Ln0.3Fe1.4 Co0.6O7 − δ with Ln= La, Nd, Sm, and Gd have been characterized. Although the substitution of 10% La, Nd, or Sm for Sr in Sr2.7Ln0.3Fe1.4 Co0.6O7 − δ is beneficial for suppressing the structural instability of Sr3Fe2O7 − δ, the Gd substitution for Sr could not help to suppress the structural degradation in air. While maintaining the tetragonal structure (space group: I4/mmm), the unit cell volume and oxygen content (7−δ) decrease from Ln= La to Gd due to the decreasing size of Ln3+ and a consequent preference for lower coordination numbers. Also, both the TEC and electrical conductivity decrease from Ln= La to Gd due to, respectively, the decrease in the ionicity of the Ln–O bonds and the increasing amount of oxygen vacancies. The cathode polarization 1 conductance (R− p ) measured using the ac-impedance spectroscopy
[1] N.Q. Minh, T. Takahashi, Science and Technology of Ceramic Fuel Cells, Elsevier, Oxford, 1995, p. 119. [2] S.C. Singhal, K. Kendall, High Temperature Solid Oxide Fuel Cells: Fundamentals, Design and Applications, Elsevier, Oxford, 2003, p. 127. [3] S.N. Ruddlesden, P. Popper, Acta Crystallogr. 11 (1958) 54. [4] F. Prado, L. Mogni, G.J. Cuello, A. Caneiro, Solid State Ionics 178 (2007) 77. [5] A. Manthiram, F. Prado, T. Armstrong, Solid State Ionics 152–153 (2002) 647. [6] K.T. Lee, A. Manthiram, Chem. Mater. 18 (2006) 1621. [7] T. Armstrong, F. Prado, A. Manthiram, Solid State Ionics 140 (2001) 89. [8] F. Prado, T. Armstrong, A. Caneiro, A. Manthiram, J. Electrochem. Soc. 148 (2001) J7. [9] G. Amow, S.J. Skinner, J. Solid State Electrochem. 10 (2006) 538. [10] F. Prado, K. Gurunathan, A. Manthiram, J. Mater. Chem. 12 (2002) 2390. [11] A. Manthiram, J.B. Goodenough, Solid State Chem. 92 (1991) 231. [12] Z. Zhang, M. Greenblatt, J.B. Goodenough, J. Solid State Chem. 108 (1994) 402. [13] I.D. Fawcett, G.M. Veith, M. Greenblatt, J. Solid State Chem. 155 (2000) 96. [14] V.V. Kharton, E.V. Tsipis, E.N. Naumovich, A. Thursfield, M.V. Patrakeev, V.A. Kolotygin, J.C. Waerenborgh, I.S. Metcalfe, J. Solid State Chem. 181 (2008) 1425. [15] J.A. Kilner, C.K.M. Shaw, Solid State Ionics 154–155 (2002) 523. [16] K. Yamaura, Q. Huang, R.J. Cava, J. Solid State Chem. 146 (1999) 277. [17] K.T. Lee, D.M. Bierschenk, A. Manthiram, J. Electrochem. Soc. 153 (2006) A1255. [18] L.A. Chick, L.R. Pedersen, G.D. Maupin, J.L. Bates, L.E. Thomas, G.J. Exarhos, Mater. Lett. 10 (1990) 6. [19] H. Taguchi, D. Mastuda, M. Nagao, K. Tanihata, Y. Miyamoto, J. Am. Ceram. Soc. 75 (1992) 201. [20] J. Rodriguez-Carjaval, Phys. B 192 (1993) 55. [21] A. Manthiram, J.S. Swinnea, Z.T. Sui, H. Steinfink, J.B. Goodenough, J. Am. Chem. Soc. 109 (1987) 6667. [22] L.J. Van der Pauw, Philips Res. Repts. 13 (1958) 1. [23] J. Wan, J.H. Zhu, J.B. Goodenough, Solid State Ionics 177 (2006) 1211. [24] S.E. Dann, M.T. Weller, D.B. Currie, J. Solid State Chem. 97 (1992) 179. [25] Y. Bread, C. Michel, F. Hervieu, Studer, A. Maignan, B. Raveau, Chem. Mater. 14 (2002) 3128. [26] M. Mori, Y. Hiei, N.M. Sammes, G.A. Tompsett, J. Electrochem. Soc. 147 (2000) 1295. [27] J.-H. Kim, A. Manthiram, J. Electrochem. Soc. 155 (2008) B385. [28] H. Takahashi, F. Munakata, M. Yamanaka, Phys. Rev. B 57 (1998) 15211. [29] K.T. Lee, A. Manthiram, J. Electrochem. Soc. 153 (2006) A794. [30] K. Huang, H.Y. Lee, J.B. Goodenough, J. Electrochem. Soc. 145 (1998) 3220.
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Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s s i
Characterization of Sr2.7Ln0.3Fe1.4Co0.6O7 (Ln = La, Nd, Sm, Gd) intergrowth oxides as cathodes for solid oxide fuel cells J.-H. Kim, A. Manthiram ⁎ Electrochemical Energy Laboratory & Materials Science and Engineering Program, University of Texas at Austin, Austin, Texas 78712, USA
a r t i c l e
i n f o
Article history: Received 19 June 2009 Received in revised form 18 September 2009 Accepted 21 September 2009 Keywords: Solid oxide fuel cells Cathodes Intergrowth oxides
a b s t r a c t Perovskite-related intergrowth oxides Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) have been investigated as cathode materials for solid oxide fuel cells (SOFC). With decreasing size of the Ln3+ ions, the unit cell volume, oxygen content, thermal expansion coefficient (TEC), and total electrical conductivity decrease from Ln = La to Gd. The decreasing unit cell volume and oxygen content is attributed to the decreasing size of Ln3+ ions from Ln = La to Gd and a consequent preference for lower coordination numbers. While the decrease in the ionicity of the Ln–O bonds from Ln = La to Gd causes a decrease in the TEC, the increasing amount of oxygen vacancies leads to a decrease in electrical conductivity arising from a 1 thermally activated semiconducting behavior. The cathode polarization conductance (R− p ) measured using the ac-impedance spectroscopy and the catalytic activity for the oxygen reduction reaction in SOFC decrease from Ln = La to Gd partly due to the decrease in electrical conductivity. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Transition metal oxides with mixed oxide-ion and electronic conducting (MIEC) properties find unique applications such as cathode materials in SOFC. The cathode in SOFC requires both high electronic and oxide-ion conductivities, thermal expansion compatibility with the electrolyte, good chemical stability with other cell components, and good catalytic activity for the oxygen reduction reaction at high temperatures [1]. Perovskite (La,Sr)MnO3 has traditionally been studied extensively due to its advantage of good thermal expansion coefficient match with other cell components and good chemical stability. However, the low oxide-ion conductivity and consequent low catalytic activity at intermediate-temperatures make (La,Sr)MnO3 unattractive for operation at T < 800 °C [2]. These difficulties have created enormous interest in the development of alternative cathode materials. One class of materials that have drawn some attention is the perovskite-related intergrowth oxides belonging to the Ruddlesden– Popper (R–P) series [3] having the general formula of (AO)(ABO3)n with n = 1, 2, 3. These intergrowth oxides have rock-salt AO layers alternating with a single (n = 1), double (n = 2), or triple (n = 3) perovskite (ABO3)n layers along the c axis [4,5]. For example, Fig. 1 shows the crystal structure of (Sr,Ln)3(Fe,Co)2O7 (n = 2) having (Sr,Ln)O rock-salt layers alternating with two (Sr,Ln)(Fe,Co)O3 perovskite layers along the
⁎ Corresponding author. Tel.: +1 512 471 1791; fax: +1 512 471 7681. E-mail address: [email protected] (A. Manthiram). 0167-2738/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2009.09.007
c axis. Various compositions in this R–P series have been reported with A = Sr and La, and B = Mn, Fe, Co, Ni, Cu, and Ga [4–17]. The MIEC properties of the (Sr,La)n + 1(Fe,Co)nO3n + 1 (n = 1–3) oxides as well as their catalytic activity in SOFC have been studied by our group before and compared with those of perovskite oxides [5–8,17]. The total electrical conductivity and the oxygen permeation flux of the (Sr,La)n + 1(Fe,Co)nO3n + 1 phases increase with n and reach a maximum value for the SrCo0.8Fe0.2O3 − δ perovskite, which can be considered as the n = ∞ member [5,8,9]. However, the R–P phases generally show better structural stability without undergoing undesired phase transitions compared to the perovskite oxides, which tend to lose too much oxygen at high temperatures and undergo a perovskite to brownmillerite transition particularly under reducing conditions. In the n = 2 series, Sr3Fe2O7 − δ suffers from structural and chemical instability at room temperature due to a slow reaction with H2O and CO2 in air. Although our group was able to overcome this instability by a substitution of 10% La for Sr, such a substitution led to a decrease in oxide-ion conductivity [8]. Interestingly, a substitution of 30% Co for Fe to give Sr2.7La0.3Fe1.4Co0.6O7 − δ improved both the electrical and oxideion conductivities and consequently the catalytic activity for ORR in SOFC. Unfortunately, the substitution of Co increases the thermal expansion coefficient (TEC) due to the low spin to the high spin transitions of Co3+ ions, which can lead to mechanical integrity problems in SOFC in contact with the electrolyte. One way to reduce TEC is to replace the La3+ ions in Sr2.7La0.3Fe1.4Co0.6O7 − δ by other Ln3+ (Ln = lanthanide) ions as the iconicity and the thermal expansion of the Ln–O bonds decrease with decreasing size of the Ln3+ ions. Accordingly, we present here an investigation of the high temperature properties of
J.-H. Kim, A. Manthiram / Solid State Ionics 180 (2009) 1478–1483
Fig. 1. Crystal structure of the intergrowth oxide (Sr,Ln)3(Fe,Co)2O7 − δ.
Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) and an exploration of their use as cathodes in SOFC. The influences of the Ln3+ ions on the crystal structure, structural stability, thermal properties, and catalytic activity for the oxygen reduction reaction in SOFC are discussed.
2. Experimental The Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) oxides were synthesized by conventional solid-state reaction. Required amounts of the lanthanide oxides (La2O3, Nd2O3, Sm2O3, or Gd2O3), SrCO3, Fe2O3, and Co3O4 were thoroughly mixed with ball milling in ethanol for 24 h and calcined at 1000 °C for 12 h in air. The calcined powders were then ground, pressed into pellets, and sintered at 1300 °C for 24 h in air. The resulting products were ground and finally heated at 900 °C for 6 h in air, followed by slow cooling to room temperature at a rate of 1 °C/min to maximize the oxygen content value. The La0.8Sr0.2 Ga0.8 Mg0.2O2.8 (LSGM) electrolyte disks were prepared by firing required amounts of La2O3, SrCO3, Ga2O3, and MgO at 1100 °C for 5 h, followed by pelletizing and sintering at 1500 °C for 10 h. Ce0.8Gd0.2O1.9 (GDC), Ce0.9Gd0.1O1.95 (10GDC), and La0.4Ce0.6O1.8 (LDC) powders were synthesized by the glycine–nitrate combustion method [18,19]. For the anode, NiO and 10GDC (Ni:10GDC = 70:30 vol.%) were ball-milled in ethanol for 48 h. The products thus obtained were characterized by X-ray diffraction (XRD) and the XRD data were refined with the Rietveld method using the Fullprof program [20]. Due to the instability problem in air, the Sr2.7Gd0.3Fe1.4Co0.6O7 − δ sample for the XRD experiment was prepared in argon atmosphere and sealed with a kapton film. The average oxidation state of Co and Fe and the room-temperature oxygen content values were determined by iodometric titration [21]. Thermogravimetric analysis (TGA) and thermal expansion data were collected with a Perkin-Elmer Series 7 thermal analysis system. The TGA experiments were carried out with a heating/cooling rate of 3 °C/min in air. The TECs of the sintered samples were measured in the temperature range of 80–900 °C with a heating/cooling rate of 5 °C/min with an intermediate dwelling at 900 °C for 0.5 h. Electrical conductivity of the pellets was measured with a four-probe dc method using a Van der Pauw configuration in the temperature range of 40–900 °C [22]. The specimens for the reactivity tests were obtained by mixing the
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Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) oxides and the GDC powders, followed by heating at 1000 °C for 3 h in air. The polarization resistance (Rp) of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathodes were measured with a half cell consisting of cathode|GDC|Pt. The Pt meshes and wires were attached to each electrode using Pt paste as the current collector. The Pt reference electrode was placed at a distance of 5 mm from the cathode. The acimpedance spectra of the half cells were obtained at different temperatures with the Solartron 1260 FRA equipment. Fuel cell performances of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathodes were evaluated with electrolyte-supported single cells. All the electrode materials were mixed with an organic binder (Heraeus V006) to form slurries and then applied onto the surface of a dense LSGM pellet (400 μm thickness) by screen printing. To prevent the formation of LaNiO3 or La2NiO4 at the anode|electrolyte interface, a LDC buffer layer was first prepared by screen printing onto the anode side of the LSGM electrolyte [23]. After heating the buffer layer at 1000 °C for 1 h, the NiO–10GDC cermet anode was screen printed onto the LDC layer and heated at 1300 °C for 0.5 h. Similarly, to prevent the side reactions at the cathode|electrolyte interface, a GDC buffer layer was prepared onto the cathode side of the LSGM electrolyte and heated at 1200 °C for 3 h. The Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathode layers were subsequently screen printed onto the GDC buffer layer and heated at 1000 °C for 3 h. Pt meshes and wires were attached to each electrode with an area of 0.25 cm2 using Pt paste as the current collector. During the SOFC operation, humidified H2 (~3% H2O at 25 °C) and air were supplied as fuel and oxidant, respectively, at a rate of 100 cm3/min. The cross sections of the single cells were observed with a scanning electron microscope (SEM, JEOL JSM-5610) after the SOFC performance tests. 3. Results and discussion The room-temperature crystal structures of the Sr2.7Ln0.3Fe1.4Co0.6 O7 − δ (Ln = La, Nd, Sm, Gd) samples were refined based on the tetragonal space group I4/mmm. The experimental XRD data, calculated profiles, and the difference between the two shown in Fig. 2 reveal good agreement between the experimental and calculated profiles. Tables 1 and 2 give the refined lattice parameters, atomic positions, occupancies of Sr and Ln, and the quality of the refinements. The lattice parameters a and b and the unit cell volume decrease from Ln= La to Gd due to the decreasing size of Ln3+. In addition, the smaller Ln3+ ions compared to the larger Sr2+ ions prefer the Ln(1) sites in the perovskite layer rather than the Ln(2) sites in the rock-salt layer in Fig. 1. For example, the occupancies of Ln at the Ln(1) site and Sr at the Sr(2) site increase as the size of the Ln3+ ions decreases from La3+ to Gd3+ as seen in Table 2. The average oxidation state values of the transition metal ions and the oxygen contents determined by the iodometric titration are given in Table 1. The room-temperature oxygen content decreases from 7.0 for Ln= La to 6.91 for Ln= Gd. It has been reported that oxygen vacancies are exclusively present on the O(1) site in the Sr(1)/Ln(1)–O(1) plane of Sr3(Fe,Co)2O7 [24,25]. Recent high-temperature neutron diffraction study also revealed that oxygen vacancies in Sr3Fe2O7 − δ are predominantly present at the O(1) site in the temperature range of 20≤T ≤ 900 °C [4]. Thus, the increase in the Ln occupancy at the Ln(1) site and the oxygen vacancies at the O(1) site on going from Ln= La to Gd could be understood to be due to the preference of the Ln3+ ions for lower coordination number with decreasing Ln3+ size. From an earlier study [8], the Sr3Fe2 − yCoyO7 composition was found to suffer from chemical and structural instability at room temperature due to a slow reaction with H2O and CO2 in air. In this study, a substitution of 10% La, Nd, or Sm for Sr in Sr2.7Ln0.3Fe1.4Co0.6O7 − δ was found to be beneficial for suppressing the instability in air. However, the substitution of Gd for Sr could not help to overcome the instability problem in air. For example, the XRD patterns of the Sr2.7Gd0.3Fe1.4 Co0.6O7 − δ powder after storing in air and vacuum indicate that while powder exposed to open air exhibits structural degradation as
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J.-H. Kim, A. Manthiram / Solid State Ionics 180 (2009) 1478–1483 Table 2 Crystallographic data of Sr2.7Ln0.3Fe1.4Co0.6O7 − δ at room temperature.a Ln
La
Nd
Sm
Gd
zSr2/Ln2 zFe/Co zO2 zO3 OccupancySr1 OccupancyLn1 OccupancySr2 OccupancyLn2 Rp Rwp χ2
0.317 0.098 0.198 0.091 0.89 0.11 1.81 0.19 6.01 7.92 2.10
0.318 0.098 0.196 0.090 0.81 0.19 1.89 0.11 5.61 7.49 1.56
0.318 0.099 0.199(1) 0.086(1) 0.76 0.24 1.94 0.06 7.47 9.51 2.69
0.317(1) 0.098(1) 0.196 0.089 0.75 0.25 1.95 0.05 9.15 12.3 1.30
a Refinements were performed based on the space group I4/mmm with the atomic positions Sr(1)/Ln(1): 0, 0, 0.5; Sr(2)/Ln(2): 0, 0, z; Fe/Co: 0, 0, z; O(1): 0, 0, 0; O(2): 0, 0, z; O(3): 0, 0.5, z.
Fig. 2. XRD patterns, calculated profiles, peak positions, and the difference between the observed and calculated profiles of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ samples: (a) Ln = La, (b) Ln = Nd, (c) Ln = Sm, and (d) Ln = Gd.
evidenced by the impurity reflections in Fig. 3(a), that stored in vacuum maintains the original intergrowth structure without any degradation as seen in Fig. 3(b). Fig. 4 shows the variations of the oxygen content and oxidation state of Fe and Co ions with temperature in air for the various Sr2.7Ln0.3Fe1.4 Co0.6O7 − δ (Ln = La, Nd, Sm, Gd) samples. These curves were derived using the initial oxygen content values (Table 1) determined by the iodometric titration and the TGA data. The Sr2.7Ln0.3Fe1.4Co0.6O7 − δ samples lose 0.4–0.45 oxygen atoms per formula unit on heating to
920 °C. At a given temperature, the oxygen content and the average oxidation state of Fe and Co ions decrease from Ln= La to Gd due to the differences in the initial, room-temperature oxygen content values (Table 1). Fig. 5 shows thermal expansion curves of the Sr2.7Ln0.3Fe1.4Co0.6O7 −δ (Ln = La, Nd, Sm, Gd) specimens measured in air. All the specimens show a linear thermal expansion behavior as indicated by a dashed line at T < 300 °C, where the oxygen loss is negligible (Fig. 4). However, they show an increase in the slope at T > 300 °C due to the loss of oxygen from the lattice (Fig. 4) and the consequent reduction of smaller Fe4+/ Co4+ to larger Fe3+/Co3+ ions. The thermal expansion curves were found to be reversible in the subsequent heating and cooling cycles and the average TEC values are given in Table 1 for two different temperature ranges, 80–900 °C and 500–900 °C. The TEC value decreases from Ln= La to Gd in both the temperature ranges. The decrease in TEC is due to the decrease in the ionicity of the Ln–O bonds as the ionic bonds generally exhibit larger thermal expansion than covalent bonds. Similar trends with the Ln–O bond ionicity have been observed in the (La,A)MnO3 [26] and LnBaCo2O5 + δ [27] systems. Fig. 6 shows the variations of the electrical conductivity with temperature for the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, Gd) samples. All the samples show an increase in conductivity with temperature for T < 500 °C, indicating a thermally activated semiconducting behavior. However, the conductivity decreases with increasing temperature for T > 500 °C due to the increasing concentration of oxygen vacancies (Fig. 4) and a consequent perturbation of the O–(Fe, Co)–O interaction and carrier localization [28]. For the same reason, at a given temperature, the electrical conductivity decreases from Ln= La to Gd due to the decreasing oxygen content and carrier concentration. The Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, Gd) samples show much lower electrical conductivities compared to the perovskite oxides such as La0.6Sr0.4CoO3 − δ (~1200 S/cm) [29]. While the perovskite structure facilitates three-dimensional electrical pathways through the O–Co–O network, the double (Sr,Ln)O rock-salt layers in between the double perovskite layers along the c axis in the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ samples breaks down the O–(Fe,Co)–O interaction along the c axis (Fig. 1) and limits the electrical conduction pathway. For this reason, the electrical
Table 1 Structural parameters, chemical analysis data, and TECs of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ samples. Sample
La0.3Sr2.7Fe1.4Co0.6O7 − δ Nd0.3Sr2.7Fe1.4Co0.6O7 − δ Sm0.3Sr2.7Fe1.4Co0.6O7 − δ Gd0.3Sr2.7Fe1.4Co0.6O7 − δ
Space group
I4/mmm I4/mmm I4/mmm I4/mmm
a (Å)
3.860 3.857 3.857 3.849
c (Å)
20.171 20.120 20.103 20.075
V (Å3)
300.577 299.270 299.084 297.417
Oxidation state of Fe,Co
Oxygen content (7 − δ)
TEC × 106 (°C− 1) 80–900 °C
500–900 °C
3.85 3.84 3.78 3.76
7.00 6.99 6.93 6.91
18.2 18.0 18.0 17.3
21.2 20.2 20.0 19.1
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Fig. 3. XRD patterns of the Sr2.7Gd0.3Fe1.4Co0.6O6.91 powder stored in (a) air and (b) vacuum. The vertical bars show the peak positions of phase-pure Sr2.7Gd0.3Fe1.4Co0.6O6.91.
Fig. 6. (a) Temperature dependence of the electrical conductivity in the temperature range of 40–900 °C and (b) Arrhenius plots of the electrical conductivity in the temperature range of 40–280 °C for the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) samples in air.
plots show the Arrhenius relation as seen in Fig. 6(b), indicating small polaron hopping conduction, which is generally expressed as follows [30], Fig. 4. Variations of the oxygen content and the average oxidation state of (Fe,Co) in Sr2.7Ln0.3Fe1.4Co0.6O7 − δ with temperatures in air: (a) Ln = La, (b) Ln = Nd, (c) Ln = Sm, and (d) Ln = Gd.
conductivity of the (Sr,La)n + 1(Fe,Co)nO3n + 1 samples increases with n and reaches the highest value for the perovskite (n = ∞ member) [8,9]. At T < 300 °C where oxygen loss is negligible (Fig. 4), the log (σT) vs. 1/T
Fig. 5. Thermal expansion (dL/Lo) curves of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ specimens in the temperature range of 80–900 °C in air: (a) Ln = La, (b) Ln = Nd, (c) Ln = Sm, and (d) Ln = Gd.
σ=
A E exp − a kT T
ð1Þ
where k is the Boltzmann constant, T is temperature, Ea is activation energy for the hopping of the small polarons, and A is the preexponential factor. The Ea for the small polaron hopping conduction
Fig. 7. XRD patterns of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ and GDC mixtures after firing at 1000 °C for 3 h in air: (a) Ln = La, (b) Ln = Nd, (c) Ln = Sm, and (d) Ln = Gd.
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1 Fig. 8. (a) Temperature dependence of the total polarization conductance (R− p ) and (b) the normalized ac-impedance spectra measured at 550 °C of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathodes in air.
increases from 0.14 eV for Ln = La to 0.17 eV for Ln = Gd in the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ system. We have reported previously that a GDC interlayer is necessary in between the (Sr,La)3(Fe,Co)2O7 cathode and LSGM electrolyte to prevent their side reaction at high temperature during the single cell fabrication [17]. Therefore, the chemical stability of Sr2.7Ln0.3Fe1.4Co0.6 O7 − δ (Ln = La, Nd, Sm, Gd) in contact with GDC was assessed by heating the mixed powders at 1000 °C for 3 h. Fig. 7 compares the XRD patterns recorded after heat treating the mixed powders. While the Ln= La sample shows a trace amount of unknown reaction product as indicated
by the reflection at 2θ ≈ 31.7°, the amount of this reaction product decreases or becomes negligible for the Ln= Nd, Sm, and Gd samples. Fig. 8(a) compares the temperature dependence of the total 1 polarization conductance (R− p ) of the Sr 2.7Ln0.3Fe 1.4Co 0.6 O7 − δ (Ln = La, Nd, Sm, Gd) cathodes on the GDC electrolyte in air. The activation energy (Ea) of the cathode polarization conductance increases with decreasing size of the Ln3+ ions from La (Ea = 1.51 eV) to Gd (Ea = 1.68 eV). In addition, at a given temperature, the cathode polarization conductance decreases from Ln = La to Gd. Fig. 8(b) compares the normalized impedance spectra of the Sr2.7Ln0.3Fe1.4Co0.6 O7 − δ (Ln = La, Nd, Sm, Gd) cathodes at 550 °C in air. The increasing resistance and Ea for polarization conductance on going from Ln= La to Gd is partly attributed to the decreasing electrical conductivity and the increasing Ea for the small polaron hopping conduction (Fig. 6). The electrochemical performance of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathodes in SOFC for the oxygen reduction reaction was evaluated using LSGM as an electrolyte. The GDC interlayer is employed between the cathode and electrolyte to prevent side reactions between the cathode and the LSGM electrolyte. Fig. 9 shows the SEM images of the cross sections of the cathode|GDC interlayer| LSGM electrolyte portion after the SOFC single cell performance tests. While the bottom of the micrograph indicates a dense, well-sintered LSGM electrolyte, the upper portion shows the GDC interlayer and the porous Sr2.7Ln0.3Fe1.4Co0.6O7 − δ cathode. The GDC interlayer with a thickness of ~ 5 μm provides a good separation between the cathode and the electrolyte. Fig. 10 compares the current–voltage (I–V) curves and the corresponding power density curves measured at 800 °C for the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ (Ln = La, Nd, Sm, and Gd) cathodes. The single cells show maximum power densities (Pmax) of 469, 447, 430, and 299 mW/cm2, respectively, for the Ln = La, Nd, Sm, and Gd cathodes. Thus, the cathode performance decreases from Ln = La to Gd in accordance with the cathode polarization resistance data in Fig. 8. The low electrical conductivity (≤80 S/cm) as seen in Fig. 6 can limit the charge-transfer rate in the overall oxygen reduction process with the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ cathode at 800 °C. Accordingly, the decreasing electrical conductivity from Ln= La to Gd leads to a decrease in the cathode performance. However, the Ln= Gd cathode exhibits much
Fig. 9. SEM micrographs showing the cross sections of the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ cathode|GDC interlayer|LSGM electrolyte portion after the SOFC single cell performance tests: (a) Ln = La, (b) Ln= Nd, (c) Ln= Sm, and (d) Ln= Gd.
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and the catalytic activity for the oxygen reduction reaction in SOFC decrease from Ln= La to Gd partly due to the decreasing electrical conductivity. Acknowledgement Financial support by the Welch Foundation grant F-1254 is gratefully acknowledged. References
Fig. 10. SOFC performance data of single cells consisting of Sr2.7Ln0.3Fe1.4Co0.6O7 − δ| GDC|LSGM|LDC|Ni-GDC at 800 °C.
inferior performance in SOFC compared to the Ln= La, Nd, and Sm cathodes. This can be understood by the structural instability of the Ln= Gd cathode in air (Fig. 3). Although the vacuum-stored powders were employed to prepare the Sr2.7Ln0.3Fe1.4Co0.6O7 − δ cathodes, the Ln= Gd cathode experiences structural instability problems in contact with air during the single cell fabrication process. 4. Conclusions The influence of the size of Ln3+ ions on the crystal structure, oxygen content, thermal expansion, electrical conductivity, and catalytic activity for the oxygen reduction reaction in SOFC of Sr2.7Ln0.3Fe1.4 Co0.6O7 − δ with Ln= La, Nd, Sm, and Gd have been characterized. Although the substitution of 10% La, Nd, or Sm for Sr in Sr2.7Ln0.3Fe1.4 Co0.6O7 − δ is beneficial for suppressing the structural instability of Sr3Fe2O7 − δ, the Gd substitution for Sr could not help to suppress the structural degradation in air. While maintaining the tetragonal structure (space group: I4/mmm), the unit cell volume and oxygen content (7−δ) decrease from Ln= La to Gd due to the decreasing size of Ln3+ and a consequent preference for lower coordination numbers. Also, both the TEC and electrical conductivity decrease from Ln= La to Gd due to, respectively, the decrease in the ionicity of the Ln–O bonds and the increasing amount of oxygen vacancies. The cathode polarization 1 conductance (R− p ) measured using the ac-impedance spectroscopy
[1] N.Q. Minh, T. Takahashi, Science and Technology of Ceramic Fuel Cells, Elsevier, Oxford, 1995, p. 119. [2] S.C. Singhal, K. Kendall, High Temperature Solid Oxide Fuel Cells: Fundamentals, Design and Applications, Elsevier, Oxford, 2003, p. 127. [3] S.N. Ruddlesden, P. Popper, Acta Crystallogr. 11 (1958) 54. [4] F. Prado, L. Mogni, G.J. Cuello, A. Caneiro, Solid State Ionics 178 (2007) 77. [5] A. Manthiram, F. Prado, T. Armstrong, Solid State Ionics 152–153 (2002) 647. [6] K.T. Lee, A. Manthiram, Chem. Mater. 18 (2006) 1621. [7] T. Armstrong, F. Prado, A. Manthiram, Solid State Ionics 140 (2001) 89. [8] F. Prado, T. Armstrong, A. Caneiro, A. Manthiram, J. Electrochem. Soc. 148 (2001) J7. [9] G. Amow, S.J. Skinner, J. Solid State Electrochem. 10 (2006) 538. [10] F. Prado, K. Gurunathan, A. Manthiram, J. Mater. Chem. 12 (2002) 2390. [11] A. Manthiram, J.B. Goodenough, Solid State Chem. 92 (1991) 231. [12] Z. Zhang, M. Greenblatt, J.B. Goodenough, J. Solid State Chem. 108 (1994) 402. [13] I.D. Fawcett, G.M. Veith, M. Greenblatt, J. Solid State Chem. 155 (2000) 96. [14] V.V. Kharton, E.V. Tsipis, E.N. Naumovich, A. Thursfield, M.V. Patrakeev, V.A. Kolotygin, J.C. Waerenborgh, I.S. Metcalfe, J. Solid State Chem. 181 (2008) 1425. [15] J.A. Kilner, C.K.M. Shaw, Solid State Ionics 154–155 (2002) 523. [16] K. Yamaura, Q. Huang, R.J. Cava, J. Solid State Chem. 146 (1999) 277. [17] K.T. Lee, D.M. Bierschenk, A. Manthiram, J. Electrochem. Soc. 153 (2006) A1255. [18] L.A. Chick, L.R. Pedersen, G.D. Maupin, J.L. Bates, L.E. Thomas, G.J. Exarhos, Mater. Lett. 10 (1990) 6. [19] H. Taguchi, D. Mastuda, M. Nagao, K. Tanihata, Y. Miyamoto, J. Am. Ceram. Soc. 75 (1992) 201. [20] J. Rodriguez-Carjaval, Phys. B 192 (1993) 55. [21] A. Manthiram, J.S. Swinnea, Z.T. Sui, H. Steinfink, J.B. Goodenough, J. Am. Chem. Soc. 109 (1987) 6667. [22] L.J. Van der Pauw, Philips Res. Repts. 13 (1958) 1. [23] J. Wan, J.H. Zhu, J.B. Goodenough, Solid State Ionics 177 (2006) 1211. [24] S.E. Dann, M.T. Weller, D.B. Currie, J. Solid State Chem. 97 (1992) 179. [25] Y. Bread, C. Michel, F. Hervieu, Studer, A. Maignan, B. Raveau, Chem. Mater. 14 (2002) 3128. [26] M. Mori, Y. Hiei, N.M. Sammes, G.A. Tompsett, J. Electrochem. Soc. 147 (2000) 1295. [27] J.-H. Kim, A. Manthiram, J. Electrochem. Soc. 155 (2008) B385. [28] H. Takahashi, F. Munakata, M. Yamanaka, Phys. Rev. B 57 (1998) 15211. [29] K.T. Lee, A. Manthiram, J. Electrochem. Soc. 153 (2006) A794. [30] K. Huang, H.Y. Lee, J.B. Goodenough, J. Electrochem. Soc. 145 (1998) 3220.