The performance of Ni/SrO-doped CeO₂ (SrDC) anode-supported cells, combined with the Sm₂O₃-doped CeO₂ (SDC) electrolyte layer, has been investigated for the direct supply of a dry CH₄. The SrO content in the Ce_1−xSr_xO_2−x solid solution affects the porous structure of the starting anode support comprising of the sintered NiO and Ce_1−xSr_xO_2−x grains and denseness of the SDC layer by co-sintering. The cell using the 5 mol% SrO-doped CeO₂ shows the highest cell performance without time-dependent degradation even at 750°C. This cell (the 5SrDC cell) also provides better cell performance with decreasing operating temperature than the conventional Ni/YSZ anode-supported cell (the YSZ cell); the maximum power density of the 5SrDC cell is higher than that of the YSZ cell by 73% at 650°C. This enhancement effect is discussed in light of the mixed ionic-electronic conduction of the CeO₂-based solid electrolytes.

Development of oxide-ion (O²⁻) conducting electrolytes has been contributing to realizing oxygen sensors and solid oxide fuel cells (SOFCs), originating with the fundamental studies by Kiukkola and Wagner on galvanic cells involving solid electrolytes.^1,2) Yttria-stabilized zirconia (YSZ) has been most frequently used as the electrolyte for the oxygen sensors and SOFCs, and a historical summary on an application of YSZ to them are, for example, concisely reviewed in a book on SOFCs.³⁾ Etsell and Flengas reported a review on the electrical properties of solid oxide electrolytes as early as 1970.⁴⁾ They pointed out that the most successful oxide electrolytes are the group 4 oxides ZrO₂, HfO₂, CeO₂, or ThO₂ having either alkaline earth oxide, Sc₂O₃, Y₂O₃, or a rare earth oxide. Solid solutions of these systems provide the O²⁻ conduction within specific ranges of temperature and oxygen pressure. Although both the oxygen sensors and SOFCs consist of the electrolyte, anode, and cathode, the requirements for the SOFC component materials are different from the oxygen sensors. The electrolyte of SOFC requires high O²⁻ conductivity and low electrical resistivity, generally used as a membrane-type ceramic. The anode requires a sufficiently porous structure for the fuel gas transportation and electrochemical activity to oxidize the supplied fuel. The cathode requires a sufficiently porous structure for oxidant gas transportation and electrochemical activity to reduce oxygen molecules supplied from the oxidant gas. The electrochemical performance of SOFCs can be improved by fulfilling these requirements in choosing the most suitable component materials and their fabrication.

Among the types of SOFCs classified by the electrolyte materials used, ZrO₂-based SOFCs are the most studied and engineered system so far.⁵⁾ As an anode material, a cermet consisting of metallic Ni and ZrO₂ based-electrolyte has been most frequently used for oxidation of H₂ and hydrocarbon (HC) fuels due to catalytic activity of Ni to the oxidation. The ZrO₂-based SOFC system, however, seems to allow stable operation only for H₂ fuels due to the susceptibility to deactivation of the Ni and ZrO₂-based cermet anode when HC fuels are used.^6,7) Wei et al. recently reported a review on the progress in Ni-based anode materials for direct HC SOFCs.⁸⁾ They suggested that promising Ni-based anode materials include the noble metals, CeO₂, Ba-containing oxides, and titanium oxides. CeO₂-based electrolytes have been extensively studied, and their fundamentals and applications are summarized in review papers for solid electrolytes and catalysts including applications to automobile exhaust gas purification, oxygen storage capacity (OSC), HC reforming, photocatalysis, etc.^4,9–19)

The direct utilization of HC fuels, which does not require endothermic reforming processes, has an advantage of overall efficiency in the electric power generation by SOFCs. CH₄ is the most favorable among the HC fuels due to the lowest C:H ratio resulting in furnishing the largest combustion energy per emitted CO₂. It is the main constituent of natural gas and abundant on the earth. Moreover, transportation and storage technologies for CH₄ have been well established. In addition, CH₄ hydrate and biogenic CH₄ can be economically feasible in the future.

Murray, Tsai, and Barnett reported the first direct-CH₄ SOFC, which was a cathode-supported type cell consisting of the porous cathode La_0.8Sr_0.2MnO₃ (LSM) pellet, coated in order with the porous 15 mol%Y₂O₃–CeO₂ (YDC) film, YSZ electrolyte film, thin YDC film, and porous Ni/YSZ anode layer.²⁰⁾ This cell exhibited a good cell performance with a low polarization resistance for the supply of wet CH₄ at 550–650°C due to inserting the YDC layer between the Ni/YSZ anode and YSZ electrolyte. However, the role of CeO₂-based oxides, which can be used as a constituent of the Ni-based cermet anode or the electrolyte, is still unclear on the cell-performance stability for the direct supply of dry CH₄ and its oxidation mechanism, as suggested by Wei et al.⁸⁾ Sm₂O₃-doped CeO₂ (SDC) has been most successfully used as a component of the Ni-based cermet anode for direct supply of CH₄, probably due to its excellent O²⁻ conduction with concurrent electronic conduction property.^21–24)

CeO₂ solid solutions doped with an alkaline earth oxide having divalent cations (Mg²⁺, Ca²⁺, Sr²⁺, Ba²⁺) also show mixed ionic-electronic conduction; their electrical conductivities are, however, somewhat lower than the ones doped with a rare earth oxide having trivalent cations (Y³⁺, La³⁺, Sm³⁺, Gd³⁺).¹⁵⁾ Although electrical properties and structural characterization have been extensively studied for the CeO₂ solid solution doped with an alkaline earth oxide,^{4,9–16,24–34)} their application to the anode material of direct-CH₄ SOFCs is very limited. Zhao and Du reported that anode-supported single cells consisting of a starting NiO–Ce_1−xCa_xO_2−x (x = 0–0.3) anode support, YSZ-SDC bi-layer electrolyte, and La_0.6Sr_0.4Co_0.2Fe_0.8O_3−y (LSCF) cathode showed an excellent cell performance at 800°C for the supply of humidified CH₄.³⁵⁾ Since the CH₄ supply with water vapor at high temperatures accompanies reforming reaction producing H₂ and CO, the reaction mechanism in the electrochemical oxidation of CH₄ is difficult to understand. CaO-doped CeO₂ and SrO-doped CeO₂ (SrDC) showed excellent ionic conductivities and high ionic transference numbers for O²⁻, and the ionic conduction was not significantly reduced in the presence of a second phase, e.g. SrCeO₃ in the SrDC, even exceeding the solid solubility limit.²⁶⁾ This property is favorable for the application of the anode material of the direct-CH₄ SOFCs. The OSC property in the SrDC is further expected to enhance electrochemical oxidation of CH₄, as found in the improvement of the OSC in CeO₂–ZrO₂ mixed oxides by the SrO addition.³⁶⁾ In this work, we fabricated Ni/Ce_1−xSr_xO_2−x (x = 0–0.1) anode supported single cells combined with the SDC electrolyte and Sm_0.5Sr_0.5CoO₃ (SSC) cathode. Effects of the SrO content on the microstructure in fabricating single cells and their cell performance have been investigated for the direct supply of CH₄. The cell performance was compared with that of the Ni/YSZ anode-supported single cell.³⁷⁾

The SDC powder was prepared by the coprecipitation method. Ce(NO₃)₃·6H₂O (98%, Nacalai Tesque) and Sm(NO₃)₃·6H₂O (95%, Nacalai Tesque) were mixed in a molar ratio of 4:1, and dissolved in deionized water. CO(NH₂)₂ (99%, Nacalai Tesque) was added little by little to keep the solution temperature at 90°C, and the precipitate was formed during 3 h at a stirring speed of 300 rpm. The precipitate was filtered and washed three times with distilled water. After drying at 100°C, the precipitate was ground in an agate mortar with a pestle, and calcined at 800°C for 3 h. The calcined powder was ground in the same manner and subsequently milled at 300 rpm for 24 h using 10 mm diameter zirconia balls added with an appropriate amount of ethanol. After the ball-milling, the slurry was dried at room temperature and sieved to obtain the uniformly sized SDC powder, whose nominal composition is denoted as Ce_0.8Sm_0.2O_1.9.

The SrDC powders having four different compositions were prepared by the solid-state reaction method. Appropriate molar ratios of CeO₂ powder (99%, Nacalai Tesque) and SrCO₃ powder (95%, Nacalai Tesque) were weighed, and the same procedure as described to prepare the SDC powder was applied to obtain dried powder mixtures. The powder mixtures were calcined at 900°C for 12 h and ground in an agate mortar with a pestle after the calcination. The calcined powders were heated at 1450°C for 30 h and milled under the same ball-milling condition as described above. The nominal molar concentrations of SrO to CeO₂ were 2.5 mol%, 5.0 mol%, 7.5 mol%, and 10.0 mol%; we denote the prepared SrDC powders as the 2.5SrDC, 5SrDC, 7.5SrDC, and 10SrDC powders, respectively.

Each nSrDC (n = 2.5, 5, 7.5, 10) powder was mixed with NiO (99%, Wako Pure Chemical) in a mass ratio of 4:6. The mixture was added with cornstarch at a ratio of 15 mass% as a pore former and stearic acid (95%, Sigma-Aldrich) at a ratio of 5 mass% as a binder, and they were milled under the same ball-milling condition as described in 2.1. After drying at room temperature, 0.6 g of the milled powder was put in a cylindrical die with an inner diameter of 16.5 mm and uniaxially pressed at 10 kN to make a pellet. The pellet was subjected to a two-step pre-sintering, heating at 800°C for 1 h to eliminate the pore former followed by heating at 1100°C for 3 h to obtain a pre-sintered NiO–nSrDC pellet. The heating rate was fixed at 100 K h⁻¹.

The SDC powder was mixed with an ethyl cellulose solvent-soluble polymer (ETHOCEL STD-100, Dow Chemical) in a mass ratio of 1:1 to form a slurry. The one face of the pre-sintered anode pellet was uniformly coated with the slurry using a masking tape having a circular hole with a diameter of 13 mm. The coated pellet was dried at 80°C for 30 min and subsequently heated at 100 K h⁻¹ and kept heating at 1400°C for 10 h. A cathode slurry consisting of the SDC powder, Sm_0.5Sr_0.5CoO₃ (SSC, AGC Seimi Chemical) powder, and STD-100 was prepared. The SDC and SSC powders were mixed in a mass ratio of 3:7 in agate mortar with a pestle, and they were mixed with STD-100 in a mass ratio of 2:1. The prepared slurry was applied to the surface of the SDC layer using a masking tape with a diameter of 6 mm, and a circular Pt mesh (#100 mesh) with the same diameter was attached to the slurry surface. The same size Pt mesh was attached to the bottom face of the anode pellet using a Pt paste (TR-7070, Tanaka Kikinzoku Kogyo), and the prepared cell sample was heated at 1200°C for 3 h.

Another type of anode-supported cell consisting of the Ni and 8 mol% yttria-stabilized zirconia (YSZ), YSZ electrolyte, and La_0.8Sr_0.2MnO₃ (LSM)-based cathode (Ni/YSZ|YSZ|LSM) was prepared. The same preparation procedure was applied as we described previously.³⁷⁾

The anode face and the cathode face of the prepared cell were respectively sealed with a PYREX^® glass ring gasket to an end of the alumina tube. The cell was heated at 850°C for 1 h to soften the glass ring gasket, and the NiO–nSrDC anode support was reduced by H₂ at 750°C for 1 h to obtain the Ni/nSrDC anode. The prepared anode-supported cell (Ni/nSrDC|SDC|SSC) is denoted as the nSrDC cell here. The same reducing treatment was applied to the YSZ electrolyte-based anode-supported cell, denoted as the YSZ cell. The cell’s electrochemical performance was evaluated using current-voltage (I-V) characteristics. A Pt wire with the 0.3 mm diameter was respectively spot-welded to the Pt mesh, each attached to the anode and cathode. Since a cell performance measured by the two-terminal lead wire configuration is affected by the ohmic resistance of the Pt wire, we used the same wire length and sample placement in the measurement. The cell performance was evaluated by the current-voltage (I-V) characteristics. H₂ and O₂ were supplied as fuel and oxidant, respectively, at a flow rate of 100 cm³ min⁻¹. After the cell performance measurement by supplying H₂, the fuel gas was switched to CH₄ (20 v/v%CH₄ diluted with He) with a total flow rate of 100 cm³ min⁻¹, and the same measurement was conducted. The cell performance measurement was conducted two times to confirm the reproducibility. Impedance spectra of the prepared cells were recorded with a frequency response analyzer (FRA5097, NF Corp.) under the open-circuit voltage (OCV) condition.

Microstructural observation was made with a scanning electron microscope (TM3000, Hitachi High-Tech). The sample surface was sputter-coated with Au to provide electrical conductivity. Phase identification of the samples was made by X-ray diffraction (XRD, RINT-2200, Rigaku) with the monochromated Cu-Kα emission powered at 40 kV and 30 mA.

Figure 1 shows cross-sectional microstructures of the co-sintered anode and electrolyte layers in the NiO–5SrDC|SDC and NiO–YSZ|YSZ samples. The two samples both showed about the same total thickness. The other NiO–nSrDC|SDC samples also had about the same thickness. The average thickness of the SDC and YSZ electrolyte layers was about 25 µm. The interface between the NiO–5SrDC anode and SDC electrolyte was not flat but well connected, and the interface between the NiO–YSZ anode and YSZ electrolyte was relatively flat showing a well-connected structure. Large and elongated pores were observed in the NiO–YSZ layer. The uniaxial pressing appears to form this specific porous sintered structure due to the fluidity of the fine YSZ (TZ8Y, Tosoh) powder with the average grain size of 0.3 µm and NiO powder with the pore-former and binder. In contrast to the sintered NiO–YSZ layer, the NiO–5SrDC layer exhibited a uniform porous structure, probably due to low fluidity in the pressed mixture containing the 5SrDC particles, which has a larger average size with a broader size distribution than the YSZ particles.

Fig. 1

Cross-sections of the co-sintered anode and electrolyte layers in the NiO–5SrDC|SDC and NiO–YSZ|YSZ samples, whose images were observed with back-scattered electrons (BSE).

Figure 2 shows a comparison of the microstructures among the NiO–nSrDC anode layers. These images revealed that the average size of the nSrDC grains increased with the SrO content, indicating that an enhancement of the grain growth occurred during the heating at 1450°C with the substitutional incorporation of the Sr²⁺ ions for Ce⁴⁺ ions in the host CeO₂ lattice by diffusion. The average size of NiO grains in the NiO–nSrDC anode layers was not affected by the SrO content after the heating. Studies of sintered Ce_1−xSr_xO_2−x samples have revealed their electrical, electrochemical, and thermal expansion properties and the phase equilibria and crystal structural characterization in the CeO₂–SrO system.^26–33) However, effect of the SrO content on the grain growth during the solid-state reaction seems unclear. The Ce_1−xSr_xO_2−x (x = 0.05–0.20) sintered pellets, whose precursor was a mixture of the metal nitrates, exhibited almost the same grain size after sintering at 1200°C.³²⁾ This result does not agree with our result on the growth of the SrDC grains with an increase in the SrO content during the heating at 1450°C. By contrast, Zheng et al. reported the enhancement effect of Sr addition (x = 0.01–0.10) on the densification and reducing sintering temperature of Ce_0.8Sm_0.2O_1.9 in the solid-state reaction among the CeO₂, Sm₂O₃, and SrCO₃ powders at 1400–1600°C.³⁴⁾ They did not refer to the grain growth in the sintered SrO-doped Ce_0.8Sm_0.2O_1.9; however, the effect of SrO addition on the sintering of Ce_0.8Sm_0.2O_1.9 by the solid-state reaction can account for the growth of the nSrDC grains found in our study. The X-ray diffraction patterns for the nSrDC powders, shown in Fig. 3, exhibited a reflection peak broadening and its shift moving to lower angles on increasing the SrO content. The 422 reflection peak position of the 7.5SrDC powder (Fig. 3(d)) is close to that of the 10SrDC (Fig. 3(e)), which can be interpreted from the solid solubility limit of x = 0.09 in the CeO₂–SrO system.^27,29,30) The reflection peak broadening and the peak shift with the SrO content indicates that the diffusion and displacement of the Sr²⁺ and Ce⁴⁺ ions occur in the host CeO₂ lattice by the solid-state reaction. This reaction provides the lattice distortion accompanying a low crystallization in the cubic fluorite structure and probably an enhancement in the growth of the SrDC grains during the heating due to a significant difference in the ionic radius between the Sr²⁺ ions (0.126 nm) and Ce⁴⁺ ions (0.097 nm).³⁸⁾

Fig. 2

Cross-sections of the NiO–nSrDC anode layers observed with BSE.

Fig. 3

X-ray diffraction patterns for the nSrDC powders heated at 1450°C for 30 h in air.

Figure 4 shows a comparison of the cross-sectional microstructures among the NiO–nSrDC|SDC samples, which revealed that the SrO content higher than n > 2.5 made the SDC layer dense enough as an electrolyte. Densification of the SDC layer during co-sintering is controlled by shrinkage of the anode-support layer consisting of the nSrDC and NiO particles. The smaller particle size in the CeO₂ and 2.5SrDC appears to result in a low shrinkage of the NiO–nSrDC (n = 0 and 2.5), providing insufficient sintering of the coated SDC overlayer.

Fig. 4

Cross-sections of the co-sintered anode and electrolyte layers in the NiO–nSrDC|SDC samples observed with BSE.

The 5SrDC cell showed the highest cell performance among the nSrDC cells investigated. Table 1 summarizes values of the OCV and maximum power density (P_max). Figure 5 shows a comparison of the I-V characteristics at 750°C and 600°C between the 5SrDC and YSZ cells. This cell-performance measurement was made within approximately 1 h after supplying CH₄ to the anode, confirming that the cell voltage as a function of the current density provided a stable value. The 5SrDC cell exhibited a lower P_max than the YSZ cell by 5% at 750°C, whereas the 5SrDC cell exhibited a higher P_max than the YSZ cell by 73% at 600°C. The OCV of the YSZ cell at 750°C for the CH₄ supply provided an acceptable value, taking into account that the transference number of oxide ions in the YSZ electrolyte is unity. The OCV of this cell at 600°C was lower than that at 750°C, which does not agree with the theoretically expected temperature dependence of the electromotive force. Since the electrochemical measurement we made started at 750°C and finished at 600°C, the thermal decomposition of CH₄, which can quickly occur at temperatures higher than approximately 700°C,³⁹⁾ leads to carbon deposition on the surface of the Ni particles. Lowering in the electrochemical activity at the electrode-electrolyte interface due to the carbon deposition can decrease the OCV with the temperature decrease under the supply of CH₄. By contrast, the OCV of the 5SrDC cell at 600°C was higher than that at 750°C, indicating an increase in the oxide-ion transference number of the SDC electrolyte as the mixed ionic conductor with lowering the temperature.²⁷⁾ The I-V characteristics of the 5SrDC were much higher than that of the YSZ cell at 600°C, revealing that the Ni/5SrDC anode combined with the SDC electrolyte has higher activity electrochemical oxidation of CH₄ than the Ni/YSZ anode with the YSZ electrolyte. Comparing the decreasing ratio in P_max from 750°C to 600°C between the nSrDC and YSZ cells can be a criterion for evaluating the temperature-dependent cell performance decrement. This comparison revealed that the nSrDC (n = 5, 7.5, 10) cells showed a lower decreasing ratio of 0.40–0.45 than the YSZ cell showing the ratio of 0.67. Since the oxidation of CH₄ is exothermic, the reaction favors products on decreasing temperature provided that a sufficient amount of oxygen transported through the electrolyte is supplied to CH₄. Although the CeO₂-based electrolytes have mixed ionic-electronic conduction, their high electrical conductivity and catalytic oxidation activity are suitable for efficient electrochemical oxidation of CH₄ at low temperatures, as pointed out by Montini et al.¹⁷⁾

Table 1 OCV and P_max of the nSrDC and YSZ cells for a dry CH₄ fuel.

Fig. 5

I-V characteristics of the 5SrDC and YSZ cells supplied with CH₄ fuel at (a) 750°C and (b) 600°C.

Figure 6 shows a comparison of the impedance arcs between the 5SrDC and YSZ cells. The 5SrDC cell exhibited a smaller arc than the YSZ cell at 750°C and 600°C, respectively. The intersection of the arc with the real axis at high frequencies represents the total ohmic resistance of the measurement circuit, including a cell sample and a Pt lead wire. The results in Figs. 6(a) and (b) show that the Ni/5SrDC|SDC layers have lower ohmic resistances than the Ni/YSZ|YSZ layers, and lowering in the ohmic resistance reflects a smaller voltage drop from the OCV in the 5SrDC cell than in the YSZ cell. The polarization resistance (R_p), showing a sum of the charge-transfer resistance at the electrode-electrolyte interface and the mass-transport resistance, is given by the arc intercept on the real axis. The 5SrDC cell provided R_p of 0.13 Ω cm² at 750°C and 1.0 Ω cm² at 600°C, and the YSZ cell provided R_p of 0.58 Ω cm² at 750°C and 2.8 Ω cm² at 600°C. The difference in R_p between the 5SrDC and YSZ cells at the same temperature was not directly related to the cell performance. The smaller R_p in the 5SrDC cell, thus, suggests that the charge-transfer resistance decrease involves electronic conduction at the anode-electrolyte interface specific to the CeO₂-based electrolytes due to their mixed conduction properties. We observed that consumption of the lattice oxygen of the SDC electrolyte in the partial oxidation of CH₄, producing a nearly stoichiometric ratio of CO and H₂, can significantly release electrons from the SDC at the anode-electrolyte interface at OCV.²⁴⁾ This finding accounts for lowering R_p in the 5SrDC cell.

Fig. 6

Impedance spectra for the 5SrDC and YSZ cells supplied with CH₄ fuel at (a) 750°C and (b) 600°C.

Figure 7 shows the cell-voltage change with time on supplying CH₄ at a constant current at 750°C; we fixed the current density to provide the initial cell voltage to 0.20 V. The fixed current density was 0.53 A cm⁻² for the 2.5SrDC cell, 0.87 A cm⁻² for the 5SrDC cell, 0.83 A cm⁻² for the 7.5SrDC cell, and 0.84 A cm⁻² for the 10SrDC cell. The nSrDC (n = 5, 7.5, 10) cells showed a gradual increase in the cell voltage with time, whereas the 2.5SrDC cell showed a significant decrease mainly due to the insufficient dense structure of the SDC layer, as shown in Fig. 4(b). The cell performance of the nSrDC cells (n = 2.5, 5, 7.5) exhibited increases in P_max following the cell voltage increase. The increases in P_max shown in percentage after 12 h were 6% for the 2.5SrDC cell, 7% for the 5SrDC cell, and 11% for the 10SrDC cell. These increases in the cell voltage and P_max corresponded to a decrease in R_p observed by the impedance measurement, indicating a reduction in the charge-transfer resistance at the anode-electrolyte interface caused by the current passing effect. The 10SrDC cell showed time-dependent increases in the cell voltage and P_max slightly higher than the 5SrDC and 7.5SrDC cells. The reason for this cell performance difference cannot be precisely presented within the accuracy of experimental errors; the porous microstructure difference among the sintered Ni/nSrDC cermets, shown in Figs. 2 and 4, may account for the difference. Further prolonged cell-performance measurement can reveal this difference. We have recently reported that the YSZ cell showed a stable cell voltage at a constant current up to about 20 h at 750°C for CH₄, accompanying a decrease in R_p.³⁷⁾ In the more prolonged cell-performance measurement, we found that the YSZ cell showed a cell-voltage decrease from 0.53 V at 0 h to 0.30 V after 250 h, corresponding to an increase in R_p about twice.³⁷⁾ An enhancement effect on the oxidation of CH₄ by the lattice oxygen of the SrDC, accompanying the cell-performance stability, can be interpreted by the mixed ionic-electronic property of the SrDC. The release of the lattice oxygen from the SrDC for the oxidation of CH₄ will be compensated by the electron supply forming a tentative short circuit in the SDC electrolyte near the interface connected to SrDC grains of the anode. The sufficient supply of O²⁻ from the SDC to the SrDC at high current densities can result in a prolonged stable cell performance without degradation due to the carbon deposition.

Fig. 7

Cell voltage change with time for the nSrDC cells supplied with CH₄ fuel at 750°C.

Our finding has revealed that the nSrDC cells consisting of the porous anode support, dense SDC electrolyte, and SSC-based cathode are promising as a direct CH₄ feeding SOFC at low temperatures. Further study on the improvement mechanism of the electrochemical oxidation of CH₄ in the Ni/SrDC anode supported cell is in progress.

We compared the cell performance for the direct supply of CH₄ between the two types of anode-supported single SOFCs; nSrDC (n = 0–10) and YSZ cells. The SrO content in the CeO₂–SrO system affected the formation of anode support porous structure and denseness of the SDC electrolyte layer by co-sintering. The nSrDC (n > 2.5) cells showed a prolonged stable cell performance for 12 h at 750°C. We found that the 5SrDC cell shows the highest cell performance and much higher power density despite a lower OCV than the YSZ cell at 600°C. This result indicates that the Ni/5SrDC cermet anode combined with the SDC electrolyte provides a good cell performance for the direct utilization of a CH₄ fuel. The significant enhancement in the electrochemical oxidation of CH₄ by the nSrDC cells can be interpreted in light of their mixed ionic-electronic property.

This work was supported by JSPS KAKENHI Grant Number 18H01746. Nicharee Wongsawatgul would like to acknowledge the financial support in the form of a grant from the Pacific Rim Green Innovation Hub Project (Nagaoka University of Technology (NUT)) supported by MEXT. The advanced methane research utilization project of NUT, supported by MEXT, provided some analytical instruments.

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