# Anodic Reaction in Syngas-Fueled Proton-Conducting Solid Oxide Fuel Cells

Authors: Hyun Ho Shin; Junil Choi; Taiho Park; Won Bae Kim
Journal: International Journal of Energy Research
 (2023)

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## Abstract

Anode-supported proton-conducting solid oxide fuel cells (PC-SOFCs) fabricated with two representative proton-conducting oxides, BaCe0.7Zr0.1Y0.1Yb0.1O3-δ (BCZYYb) and BaZr0.8Y0.2O3-δ (BZY), were compared to obtain the insight into the electrochemical performances when fueled with syngas at 700°C and the correlation between the anode thickness (0.4, 0.8, and 1.6 mm) and operational stability. We have demonstrated that, in stability tests, the BCZYYb cells exhibited significantly higher maximum power density (MPD) than the BZY cells when operating on H2, 1.22 and 0.48 W/cm2 for the BCZYYb and BZY cells, but the BCZYYb cells degraded more rapidly than the BZY cells when operating on syngas. In addition, decreasing the anode thickness significantly enhanced the stability of BCZYYb cells operating on syngas, a reduction of 81, 76, and 71% in MPD for 1.6, 0.8, and 0.4 mm anode after 30 min. The electrochemical impedance spectra and X-ray diffraction patterns indicated that the rapid degradation of cerate-based cells with syngas could be mainly attributed to the considerable increase in polarization resistance due to the phase decomposition of the electrolyte powder in the anode. Heterogeneous catalysis was performed to study the catalytic reaction of the H2–CO mixture over anode powders prepared with the two proton-conducting oxides (Ni-BCZYYb and Ni-BZY) in a fixed bed reactor, and the CO conversion and selectivity to CO2 and CH4 were determined. For all anode powders, continuous CO2 production was initially observed with CH4 formation, and no significant difference in catalytic activity trends was observed between both anode powders. An increased residence time substantially decreased the normalized CO2 yield, which was associated with the potential secondary reaction of CO2 with H2, CH4, or perovskite oxide. Based on the results of heterogeneous catalysis with both anode powders, the observed cell degradation on both cells during operation with syngas may be primarily attributed to carbon coking due to CO disproportionation; however, more than 4 times rapid degradation of the cerate-based cell when operating with syngas was clearly demonstrated to be attributed to the decomposition of the electrolyte powder in anode by the resulting CO2 from the catalytic reaction of the H2–CO mixture.


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## Body

## 1. Introduction

Among various fuel cell technologies, solid oxide fuel cells (SOFCs) have attracted increasing interest as alternative energy generators because of their high efficiency. They generate electrical energy through the direct conversion of chemical energy, so they can run with no pollutant emissions and have excellent fuel flexibility [1, 2]. SOFCs are distinguished by the fact that O2− anions are transported from the cathode to the anode through a dense electrolyte such as yttria-stabilized zirconia, which theoretically allows them to operate with any combustible fuel such as hydrocarbons [3]. However, the high activation energy for the transport of O2− anions inevitably requires an undesirable high operating temperature (>800°C), which increases costs and results in challenges with material incompatibility because of the significant mismatch in thermal expansion coefficients and interdiffusion reactions between cell components [4].An alternative to lowering the operating temperature is to utilize solid-state proton conductors with lower activation energy (0.4–0.6 eV) for proton transport, which would enable superior electrochemical performance owing to the high ionic conductivity in the target temperature range of 500–700°C. In contrast to O2− ions, the transport direction of the mobile ions would enhance fuel utilization through water formation at the oxygen electrode and thus minimize fuel dilution [5, 6]. Promising candidates for ceramic proton conductors include perovskite-structure alkaline earth cerates with high conductivity and more chemically stable zirconates in the presence of CO2 and H2O. Many researchers have devoted significant effort to optimizing the stoichiometry for stability and conductivity by making a solid solution between cerates and zirconates and using cation substitution [7, 8]. However, despite their potential, the practical application of proton-conducting SOFCs (PC-SOFCs) has lagged behind that of their traditional O2- anion-conducting counterparts because of fabrication challenges originating from poor sinterability and chemical instability. Recently, Duan et al. provided an essential step toward resolving such fabrication challenges through solid-state reactive sintering; they realized the direct conversion of a raw precursor oxide by using a sintering agent such as nickel oxide to fabricate a complete anode-supported SOFC with minimal ohmic loss in a single sintering step [8].Despite significant advances in terms of material and fabrication, the requirement of hydrogen as the fuel remains one of the biggest challenges for practical implementation of current SOFCs because of the vulnerability of conventional Ni-based anodes to carbon coking and sulfur poisoning [9]. Moreover, switching to a proton transport mechanism would unfortunately induce a potential loss of fuel flexibility because of the absence of an available oxidant on the anode side, which would more likely catalyze carbon coking [3]. Despite the difficulty of directly utilizing hydrocarbons as a fuel for PC-SOFCs, Duan et al. reported that a BaZr0.8Y0.2O3-δ- (BZY-) based single cell performed exceptionally with various hydrocarbon fuel streams and without any modifications to the cell composition or architecture. However, most of the fuel cell tests were conducted at a significantly high inlet steam-to-carbon ratio. As a promising alternative to the vulnerability of conventional Ni-based anodes for hydrocarbon-fueled PC-SOFC, perovskite oxide such as La0.5Sr0.5Fe0.9Mo0.1O3-δ [10] and Sr2Fe1.5Mo0.5O6-δ [11], furthermore, in situ exsolved metallic nanoparticle-structured perovskite such as Sr2Fe1.4Ni0.1Mo0.5O6-δ [11] and Sr2Fe1.4Co0.1Mo0.5O6-δ [12] which has been extensively studied due to their enhanced structural stability and flexible mixed conductivity, while their application to anode-supported cell structure remains a challenge to be solved.One approach to realizing PC-SOFCs fueled by hydrocarbons that has recently gained interest has been the hydrocarbon reforming process, which produces synthetic gas (syngas) including H2 and CO by either steam [13] or CO2 [10, 14]. This process is shown below:
(1)CH4+H2O⟶CO+3H2(2)CH4+CO2⟶2CO+2H2The reforming process offers economic and environmental advantages such as the direct conversion of major greenhouse gases (e.g., CH4 and CO2) to hydrogen fuel and syngas, which are key feedstocks for chemical synthesis. Furthermore, the reforming process is potentially applicable to limiting greenhouse gas emissions and eventually realizing a net-zero emission energy system [15]. To pair the reforming process with SOFC technology, one approach has been to integrate an additional catalytic layer with high activity and coking resistance for the reforming process on a conventional Ni-based anode. The resulting H2 fuel is then electrochemically oxidized to generate electrical power. This approach can potentially minimize the vulnerability of the electrolyte (e.g., doped BaCeO3 perovskite) in the atmosphere, which has high concentrations of CO2 and water [13]. Combining the reforming process with SOFC technology also offers other significant advantages, such as energy compensation between the extremely endothermic nature of the reforming process and exothermic nature of electrochemical oxidation [16].Because of the potential advantages of combining SOFC technology with the reforming process, many studies have been focused on developing catalysts for the on-cell reforming process of direct hydrocarbon PC-SOFCs [13, 17]. However, relatively little is known about the durability of conventional cermet anodes against carbon coking or decomposition of the PC oxide by products of unfavorable side reactions between CO and/or H2 over Ni, such as CO methanation or disproportionation. Furthermore, the extent of the catalytic reaction of the resulting syngas in the microporous anode and how the reaction products affect the electrochemical performance of the anode are still unclear [18, 19].In this study, high-performance anode-supported PC-SOFCs were fabricated with varying anode thicknesses and using two representative PC oxides: BaCe0.7Zr0.1Y0.1Yb0.1O3-δ (BCZYYb) and BZY. The objective was to study how the reaction products of the CO+H2 mixture catalyzed over the microporous cermet anode and the residence time would influence the electrochemical performance when fueled by syngas at 700°C. The catalytic reaction of the CO+H2 mixture over PC anode powder samples was observed to clarify the reaction products and mechanism. In addition to measurements of the CO conversion and the selectivity to CH4 and CO2, the catalytic reaction was with different residence times to elucidate the potential secondary reaction mechanism from the product CO2 with other gas-phase products or the PC oxide.


## 2. Experimental

### 2.1. Materials and Characterization

Precursor powders were prepared by mixing precise amounts of BaCO3 (>99%, Sigma-Aldrich), CeO2 (>99%, Sigma-Aldrich), ZrO2 (>99%, Sigma-Aldrich), Y2O3 (>99%, Sigma-Aldrich), Yb2O3 (>99%, Sigma-Aldrich), NiO (>99%, Kojundo), and starch (Alfa Aesar) as the pore former to realize the desired stoichiometry for different components. For example, NiO, electrolyte, and starch were mixed at a weight ratio of 48 : 32 : 20 to realize the anode composite. The precursor powder mixture was prepared by planetary ball milling (Retsch, PM100) in ethanol at 400 rpm for 1 h and then was dried. The same procedure was used to prepare the anode functional layer (AFL) and electrolyte precursor powders by mixing electrolyte and NiO at the weight ratios of 60 : 40 and 99 : 1, respectively, in the absence of starch. X-ray diffraction (XRD) (D8 Advance (Bruker)) with Rigaku Ultima IV was used to analyze the composite anode powders under the conditions of 40 kV and 20 mA in the 2θ range of 20–80°.


### 2.2. Cell Fabrication

Anode-supported half-cells with a configuration of composite anode support/AFL/electrolyte were fabricated by a single reduced-temperature firing based on solid-state reactive sintering and the drop-coating method. NiO was utilized as a sintering aid for full densification of the electrolyte. Two electrolyte powders were utilized to study their influence in the composite anode on the operation with syngas: BCZYYb and BZY. The anode disc was fabricated by mechanical powder pressing with varying amounts of the prepared anode powders to obtain cells with different anode thicknesses. This was to adjust the residence time of the reactant in the microporous anode. A thin AFL and electrolyte layer were applied on the pressed composite anode pellet by the drop-coating technique. Slurries for the drop coating were prepared by dispersion of the mixed powders in the organic solvent at a fluid-to-solid weight ratio of 10 : 1. The organic solvent comprises a mixture of ethanol and organic binder (V006 (Heraeus)) at a weight ratio of 10 : 0.2. The prepared slurries were mixed by planetary ball milling (Retsch, PM100) for 1 h, which was followed by ultrasonic treatment for 1 h to realize homogeneous dispersion of the powder in the slurry. The well-dispersed AFL slurry was drop-coated onto the composite anode pellet, which was followed by heat treatment at 500°C for 1 h. The same procedure was used to deposit the electrolyte on the coated AFL; then, sintering was carried out at 1400°C for 10 h (1450°C for BZY). The resulting anode thicknesses were characterized using scanning electron microscopy (SEM, Hitachi, S-4300). The paste for the cathode comprised PrBa0.5Sr0.5Co1.5Fe0.5O3-δ (PBSCF, Kceracell) and the organic binder V006 (Heraeus) at a weight ratio of 1 : 1. The paste was screen-printed onto the electrolyte, and the complete cell was then sintered at 900°C for 4 h. The active area of the cathode was 0.28 cm2 with a thickness of ~20 μm. The current collector was fabricated from Ag conductive paste (Alfa Aesar) and wire. The anode thickness of the fabricated cell has been characterized by SEM.


### 2.3. Fuel Cell Testing

The fabricated single cells were mounted on an alumina tube and sealed with an alumina-based ceramic adhesive (552-VFG, Aremco) to fix the cell and prevent gas leakage. The composite anode was reduced under dry H2 at 700°C before electrochemical measurement, where 200 mL/min dry air was supplied to the cathode as the oxidant, and 50 mL/min dry hydrogen was supplied to the anode as the fuel. The AC electrochemical impedance spectra (EIS) and current–voltage characteristics were collected by using a potentiostat (BioLogic SP-300) at an operating temperature of 700°C in H2, H2–N2, or H2–CO (molar ratio of mixed gases was 1 : 1) flowing at 50 mL/min. The AC EIS was measured in the frequency range of 250 kHz to 0.1 Hz with a perturbation AC amplitude of 10 mV and under the open-circuit voltage (OCV) condition.


### 2.4. Catalyst Testing

The catalytic reaction of the CO+H2 mixture was conducted by using a composite anode powder (Ni-BCZYYb or Ni-BZY at a weight ratio of 6 : 4) as a fixed bed in a continuous-flow quartz tube reactor to evaluate the catalytic activity. Prior to catalyst testing, the anode powders were calcined at 1400°C for 1 h. Then, sieved powder (106–150 μm) was utilized to measure the activity at 700°C after reduction in H2 for 1 h. The powder sample was kept at the center of the quartz tube reactor by using quartz wool. Reactants (i.e., 1 : 1 mixture of CO and H2) were fed to the quartz tube reactor at a flow rate of 50 mL/min. The sample size was varied from 0.2 to 1 g to study the influence of the residence time. The concentrations of reactants and products were analyzed for 2 h by using an online gas chromatograph (GC, Agilent 7890) with a 60/80 Carboxen 1000 column and thermal conductivity detector (TCD).


## 2.1. Materials and Characterization

Precursor powders were prepared by mixing precise amounts of BaCO3 (>99%, Sigma-Aldrich), CeO2 (>99%, Sigma-Aldrich), ZrO2 (>99%, Sigma-Aldrich), Y2O3 (>99%, Sigma-Aldrich), Yb2O3 (>99%, Sigma-Aldrich), NiO (>99%, Kojundo), and starch (Alfa Aesar) as the pore former to realize the desired stoichiometry for different components. For example, NiO, electrolyte, and starch were mixed at a weight ratio of 48 : 32 : 20 to realize the anode composite. The precursor powder mixture was prepared by planetary ball milling (Retsch, PM100) in ethanol at 400 rpm for 1 h and then was dried. The same procedure was used to prepare the anode functional layer (AFL) and electrolyte precursor powders by mixing electrolyte and NiO at the weight ratios of 60 : 40 and 99 : 1, respectively, in the absence of starch. X-ray diffraction (XRD) (D8 Advance (Bruker)) with Rigaku Ultima IV was used to analyze the composite anode powders under the conditions of 40 kV and 20 mA in the 2θ range of 20–80°.


## 2.2. Cell Fabrication

Anode-supported half-cells with a configuration of composite anode support/AFL/electrolyte were fabricated by a single reduced-temperature firing based on solid-state reactive sintering and the drop-coating method. NiO was utilized as a sintering aid for full densification of the electrolyte. Two electrolyte powders were utilized to study their influence in the composite anode on the operation with syngas: BCZYYb and BZY. The anode disc was fabricated by mechanical powder pressing with varying amounts of the prepared anode powders to obtain cells with different anode thicknesses. This was to adjust the residence time of the reactant in the microporous anode. A thin AFL and electrolyte layer were applied on the pressed composite anode pellet by the drop-coating technique. Slurries for the drop coating were prepared by dispersion of the mixed powders in the organic solvent at a fluid-to-solid weight ratio of 10 : 1. The organic solvent comprises a mixture of ethanol and organic binder (V006 (Heraeus)) at a weight ratio of 10 : 0.2. The prepared slurries were mixed by planetary ball milling (Retsch, PM100) for 1 h, which was followed by ultrasonic treatment for 1 h to realize homogeneous dispersion of the powder in the slurry. The well-dispersed AFL slurry was drop-coated onto the composite anode pellet, which was followed by heat treatment at 500°C for 1 h. The same procedure was used to deposit the electrolyte on the coated AFL; then, sintering was carried out at 1400°C for 10 h (1450°C for BZY). The resulting anode thicknesses were characterized using scanning electron microscopy (SEM, Hitachi, S-4300). The paste for the cathode comprised PrBa0.5Sr0.5Co1.5Fe0.5O3-δ (PBSCF, Kceracell) and the organic binder V006 (Heraeus) at a weight ratio of 1 : 1. The paste was screen-printed onto the electrolyte, and the complete cell was then sintered at 900°C for 4 h. The active area of the cathode was 0.28 cm2 with a thickness of ~20 μm. The current collector was fabricated from Ag conductive paste (Alfa Aesar) and wire. The anode thickness of the fabricated cell has been characterized by SEM.


## 2.3. Fuel Cell Testing

The fabricated single cells were mounted on an alumina tube and sealed with an alumina-based ceramic adhesive (552-VFG, Aremco) to fix the cell and prevent gas leakage. The composite anode was reduced under dry H2 at 700°C before electrochemical measurement, where 200 mL/min dry air was supplied to the cathode as the oxidant, and 50 mL/min dry hydrogen was supplied to the anode as the fuel. The AC electrochemical impedance spectra (EIS) and current–voltage characteristics were collected by using a potentiostat (BioLogic SP-300) at an operating temperature of 700°C in H2, H2–N2, or H2–CO (molar ratio of mixed gases was 1 : 1) flowing at 50 mL/min. The AC EIS was measured in the frequency range of 250 kHz to 0.1 Hz with a perturbation AC amplitude of 10 mV and under the open-circuit voltage (OCV) condition.


## 2.4. Catalyst Testing

The catalytic reaction of the CO+H2 mixture was conducted by using a composite anode powder (Ni-BCZYYb or Ni-BZY at a weight ratio of 6 : 4) as a fixed bed in a continuous-flow quartz tube reactor to evaluate the catalytic activity. Prior to catalyst testing, the anode powders were calcined at 1400°C for 1 h. Then, sieved powder (106–150 μm) was utilized to measure the activity at 700°C after reduction in H2 for 1 h. The powder sample was kept at the center of the quartz tube reactor by using quartz wool. Reactants (i.e., 1 : 1 mixture of CO and H2) were fed to the quartz tube reactor at a flow rate of 50 mL/min. The sample size was varied from 0.2 to 1 g to study the influence of the residence time. The concentrations of reactants and products were analyzed for 2 h by using an online gas chromatograph (GC, Agilent 7890) with a 60/80 Carboxen 1000 column and thermal conductivity detector (TCD).


## 3. Results and Discussion

Single cells were fabricated with two different PC oxides: BCZYYb (0.4 mm thick anode) and BZY (1.6 mm thick anode). They were then tested under various anodic conditions to examine the potentially adverse effect of non-H2 diluents such as N2 and CO. Figure 1 shows the V–I characteristics for the fabricated anode-supported cells to compare the electrochemical performances when fueled with H2, H2 diluted with N2 to confirm the dilution effect, and syngas including H2+CO (molar ratios of the mixed gases were 1 : 1) at 700°C. When H2 was the fuel, the OCV was 0.94 and 0.935 V for BCZYYb and BZY, respectively, which indicate completely densified electrolytes. This confirms that the 1 wt% NiO sintering aid had no significant effect on the electronic conductivity within the electrolyte under fuel cell conditions [20]. The corresponding maximum power densities (MPDs) of 1.22 and 0.48 W/cm2 for the BCZYYb and BZY cells, respectively, demonstrate the superior performance of the cells fabricated with PC oxide. When the H2 fuel was diluted with N2, the general trends were similar for both the BCZYYb and BZY cells: no significant change in OCV (0.94 and 0.92 V, respectively) and a slight decrease in the MPD (1.19 and 0.46 W/cm2, respectively). With syngas, the BCZYYb cells exhibited an identical OCV to that with the other fuel streams (0.94 V) while the BZY cells exhibited a slight decrease in the OCV (0.9 V). For the MPD, the BZY cells exhibited a small decrease of 9% compared with that of H2 diluted with N2 (0.42 W/cm2). In contrast, a more substantial decrease of 25% was observed for the BCZYYb cells (0.9 W/cm2). For the PC-SOFCs, the identical binary diffusivities when H2 was diluted with N2 and CO, respectively, imply that the fuel mass transport had a limited effect on the electrochemical performance of PC-SOFCs [21].Figure 1
Electrochemical performances of BCZYYb and BZY SOFCs with 0.4 and 1.6 mm thick anodes, respectively, under various anodic conditions (100% H2, 50% H2+50% N2, and 100% syngas): (a, b) I–V–P curves and (c, d) Nyquist plot of the cell performance for BCZYYb and BZY SOFCs at OCV and 700°C.
(a)
BCZYYb(b)
BZY(c)
BCZYYb(d)
BZYThe EIS results were used to evaluate the differences in electrochemical performance between the two cells under various anodic conditions. Because the feed stream on the cathode side was kept the same, the observed changes in electrochemical performance may be attributed solely to the changes in anodic conditions. Thus, the effects of N2 and CO on the anode side on the electrochemical performance were characterized by the EIS for each fuel stream. Figures 1(c) and 1(d) show Nyquist plots of the EIS collected at the OCV under various anodic conditions. The electrolyte ohmic resistance (Ro) was determined from the high-frequency intercept and was 0.11 and 0.24Ω·cm2 for BCZYYb and BZY, respectively, with the H2 feed stream. The polarization resistance (Rp) was the span of the impedance arc and indicated the resistance originating from the electrochemical kinetics on the electrodes. Rp was 0.03 and 0.06 Ω·cm2 for the BCZYYb and BZY cells, respectively, with the H2 feed stream. The smaller Rp for the BCZYYb cells indicates enhanced electrode kinetics. Despite the corresponding electrolyte thickness and cermet anode composition containing 60 wt% Ni, the smaller Ro and Rp for BCZYYb compared to BZY can be attributed to its higher ionic conductivity, which support a substantially higher maximum current density and power density. Dilution of the H2 feed stream with N2 resulted in no significant changes in EIS for both cells at the OCV. The BCZYYb and BZY cells had Rp of 0.04 and 0.058Ω·cm2, respectively, and Ro of 0.11 and 0.24 Ω·cm2, respectively. The slight decrease in power density when H2 was diluted with N2 may be attributed to anodic concentration polarization due to the higher molecular weight of N2 [21]. When syngas was the fuel, noticeably, different trends in the EIS were observed for the two cells. For the BCZYYb cells, Ro and Rp increased significantly to 0.15 and 0.08Ω·cm2, respectively. In contrast, the BZY cells exhibited much smaller increases compared to the other fuels with Ro and Rp of 0.32 and 0.075 Ω·cm2, respectively. These observations imply that the physicochemical properties of the electrolyte and anode for BCZYYb cells were drastically altered when syngas was used.To investigate the effect of syngas on the electrochemical performance of the single cells, stability tests were performed with cells fabricated from the two different PC oxides at various anode thicknesses. The anode thickness of the fabricated cell has been characterized by SEM (FigureS2). Figure 2 shows the cell voltage versus time in galvanostatic mode at a constant current of 0.5 A/cm2 when syngas was the fuel. For the BCZYYb cells, the operation depended on the anode thickness. Increasing the anode thickness induced significantly more rapid degradation, and cell failure occurred within 30 min when the anode thickness was greater than 0.4 mm. A direct comparison between BCZYYb and BZY cells with an anode thickness of 1.6 mm confirmed that the BZY cells degraded more gradually than the BCZYYb cells with rates of 0.24 and 1.2 V/h, respectively.Figure 2
Stability test of (a) BCZYYb and (b) BZY single cells with varying anode thicknesses (1.6, 0.8, and 0.4 mm) during galvanostatic operation at a constant current of 0.5 A/cm2. The cell voltage is shown as a function of time, and syngas (50% H2+50% CO) and flowing air were used as the fuel and oxidant, respectively, at 700°C.
(a)(b)To study the impact of the catalytic reaction products of the H2–CO mixture on the cermet anode, the degradation process was carried out under OCV conditions. In this state, the fabricated cells could be considered a pure chemical reactor without any electrochemical reaction. The I–V–P curves and EIS at OCV were collected in 10 and 30 min intervals for BCZYY and BZY cells, respectively, fueled by syngas to confirm the dominant factor of cell degradation. Figures 3(a) and 3(b) illustrate representative I–V–P curves of a BCZYYb cell with a 0.4 mm thick anode and BZY cell with a 1.6 mm thick anode, respectively, as functions of time. Both cells exhibited similar tendencies for the I–V–P curves, where the MPD decreased with increasing time. However, the BCZYYb cells showed a more significant reduction in MPD from the initial value of 0.92 to 0.265 W/cm2 after 30 min. In contrast, the BZY cells showed a significantly less pronounced decrease from 0.39 to 0.32 W/cm2 after 1 h. As in the stability test, a similar correlation between the anode thickness and cell degradation was observed; a thinner anode resulted in a more gradual reduction in MPD for the BCZYYb cells (Table 1). This indicates that the observed degradation in the stability test in galvanostatic mode was more likely due to the syngas rather than the applied voltage.Figure 3
I–V–P curves of (a) BCZYYb and (b) BZY single cells with 0.4 and 1.6 mm thick anodes, respectively, as a function of time at 700°C. Syngas (50% H2+50% CO) and flowing air were used as the fuel and oxidant, respectively. Time-dependent electrochemical performances of the BCZYYb and BZY single cells: (c) ohmic resistance and (d) polarization resistance at OCV and 700°C.
(a)(b)(c)(d)Table 1
Performances of the BCZYYb and BZY PC-SOFCs at 700°C.
0 min in syngas30 min in syngasElectrolyteElectrodeSintering conditionRo (Ω·cm2)Rp (Ω·cm2)MPD (W/cm2)Ro (Ω·cm2)Rp (Ω·cm2)MPD (W/cm2)BCZYYbNi-BCZYYb1.6 mm1400°C10 h0.1450.1550.7430.270.640.14Ni-BCZYYb0.8 mm0.150.10.760.240.570.18Ni-BCZYYb0.4 mm0.150.080.920.2360.4760.265BZYNi-BZY1.6 mm1450°C10 h0.320.0750.390.340.110.32Figures3(c) and 3(d) show the time-dependent impedance data (Ro and Rp) collected from the fabricated anode-supported BCZYYb and BZY cells at OCV when fueled by syngas. Corresponding Nyquist plot of BCZYYb-based SOFCs with the 0.4 mm and 1.6 mm thick anode as a function of time at OCV at 700°C in syngas fuel is provided in Figure S1. The initial values of Ro for the BCZYYb cells were 0.145, 0.15, and 0.15Ω·cm2 with 1.6, 0.8, and 0.4 mm thick anodes, respectively, while the corresponding Rp values were 0.155, 0.1, and 0.08 Ω·cm2, respectively. Instant increases in Rp and Ro were observed when the balance gas was switched from N2 to CO, and more rapid degradation was observed with thicker anodes. Thus, the differences in the initial EIS values may be attributed to the sensitivity of the BCZYYb anode to the syngas poisoning effect.All BCZYYb cells with various anode thicknesses exhibited a <2-fold increase inRo and >3-fold increase in Rp, even in 30 min. In contrast, the BZY cells showed increases of <10% in 1 h. For example, Ro and Rp were 0.27 and 0.64Ω·cm2, respectively, after 30 min for BCZYYb cells with a 1.6 mm thick anode. In contrast, BZY with a corresponding anode thickness had Ro and Rp of 0.34 and 0.11 Ω·cm2, respectively. In particular, for the BCZYYb cells, a more drastic increase was observed in Rp than in Ro over time, which implies that the degradation of the cell performance with syngas can mainly be attributed to the degradation of the cermet anode. The change in Rp with the anode thickness indicates that reducing the anode thickness will mitigate the degradation, which agrees with the results of the stability test shown in Figure 2.XRD measurements were conducted with crushed anode powders tested in H2 and CO+H2 to investigate the phase decomposition of the anode when fueled by syngas (Figure 4). The fabricated cells were reduced in H2 at 700°C to obtain the metallic Ni phase, which was followed by exposure to syngas for 1 h. Then, the cells were crushed to obtain the powder samples. After the reduction, the XRD patterns of the BCZYYb-NiO and BZY-NiO anode powders revealed only metallic Ni and perovskite oxide phases, which indicated the complete reduction of NiO and no decomposition of the PC oxide. The XRD patterns of the BCZYYb powder samples clearly showed BaCO3 and doped CeO2, which confirmed the decomposition by CO2. The poor chemical stability of ceria-based proton-conducting oxide under CO2 atmosphere has been reported in the literature [22]. In contrast, NiO-BZY only revealed perovskite oxide and metallic Ni, which indicates that the crystal structure of BZY remained unchanged. The PC oxides utilized in this study (i.e., BCZYYb and BZY) have already been confirmed as chemically stable in H2 and H2-CO [23]. However, there is an ongoing discrepancy regarding the stability of BCZYYb in CO2-containing atmosphere [23–25]. In addition, Hua et al. demonstrated the potential of using PC-SOFCs in syngas fuel stream by the absence of negative effects of CO on electrochemical performance [23], while Sun et al. reported that the Ni-BCZYYb cermet anode was poisoned by extremely low percentage level of CO2, supporting our observation.Figure 4
XRD patterns of BCZYYb and BZY anodes powder after 1 h exposure to H2 and H2+CO at 700°C.The catalytic reaction of the CO–H2 mixture was conducted with two different anode powders having 60 wt% NiO, which are denoted as NiO-BCZYYb and NiO-BZY. Figure 5 shows the CO conversion and selectivity to CO2 and CH4 as a function of time. The sample sizes of the two different anode powders were varied to study the influence of the residence time. Note that the calculated conversion and selectivity reported in the figures are for gas-phase products only and do not account for carbonaceous deposits. The gas-phase selectivity to the product was defined as the molar ratio of the product to CO+2CO2+CH4. All anode powder samples (i.e., NiO-BCZYYb and NiO-BZY) showed similar tendencies for the CO conversion and selectivity to CO2 and CH4. The CO conversion and selectivity to CO2 both showed a gradual increase after an initial rapid decrease. The selectivity to CH4 showed a rapid decrease. The analogous trends for NiO-BCZYYb and NiO-BZY imply that the catalytic reaction of the CO–H2 mixture was most likely initiated on the Ni surface, which is supported by the results of Shin et al. [26]. They found that hydrogen dissociation primarily took place on the metallic component of a cermet anode in a pulsed isotope experiment. Increasing the residence time significantly increased the CO conversion and selectivity to CO2 and CH4. However, this may be attributed to the increased catalyst bed size.Figure 5
(a, b) CO conversion and selectivity to (c, d) CO2 and (e, f) CH4 from the reaction mixture of CO and H2 versus time when streamed over BCZYYb and BZY anode powders at 700°C and different residence times: 1.8, 3.6, and 7.2 s.
(a)
BCZYYb-NiO(b)
BZY-NiO(c)
BCZYYb-NiO(d)
BZY-NiO(e)
BCZYYb-NiO(f)
BZY-NiOInterestingly, the selectivity to CO2 revealed a similar trend to that of the CO conversion for both anode powder samples: a slight decrease in the initial 20 min, which follows the observations for CH4. This observation was also reported in the literature [27], which indicates that the methanation of CO depends on the extent of deposited carbonaceous species over the metal surface (Equation (1)). The occurrence of the Boudouard reaction (Equation (2)) in the CO–H2 mixture was evidenced by the formation of CO2, which deposited carbon over the anode powder surface. Yan et al. reported that CO has no influence on the adsorption of H2 based on their work with a CO/H2 competitive adsorption model [19]. Therefore, the deposited carbon may react with dissociated hydrogen over Ni for methanation (Equation (5)). Meanwhile, the continuation of the reaction would lead to the deactivation of Ni catalysts by deposited carbon originating from CO disproportionation, which would block active sites for H2 adsorption. This hypothesis could be indirectly supported by the rapid reduction of the selectivity to CH4 over time and also slow deactivation of BZY cells even without phase decomposition. Despite the deactivation of the anode powder in the initial 20 min, the CO conversion and selectivity to CO2 then gradually increased over time. In addition, a delayed deactivation of CO conversion and selectivity to CO2 and CH4 were initially observed for NiO-BZY. This observed delay in deactivation may have been due to the high H2O uptake capacity of BZY, where the surface hydroxyl compounds can remove adsorbed carbon on Ni [28]. In the present study, no water was utilized; the formation of CH4 confirmed the occurrence of methanation (Equation (2)), which formed water as a product. Furthermore, in contrast to the delayed deactivation observed for NiO-BZY, NiO-BCZYYb exhibited a higher CO conversion and selectivity to CO2 after the initial deactivation. The continuous increase in CO conversion and selectivity to CO2 after the initial deactivation may be attributed to the transition of active sites for CO disproportionation from the metallic Ni to the oxide surface [4]. Furthermore, the higher CO conversion and selectivity to CO2 for NiO-BCZYYb compared to NiO-BZY after the initial deactivation may be attributed to the enhanced catalytic activity of a new oxide phase, such as doped ceria generated from the decomposition of the perovskite oxide (Figure 4) [4].
(3)2CO⟶CO2+C(4)3H2+CO⟶CH4+H2O(5)C+2H2⟶CH4Figure6 shows the influence of the residence time on the CO2 and CH4 yields normalized by the mass of the loaded anode powder at 2 min on stream and 700°C. Increasing the residence time significantly decreased the normalized CO2 yield while having no significant effect on CH4, which clearly indicates a secondary reaction with the product CO2. The significantly smaller reduction in the normalized CH4 yield suggests that the resulting CO2 was unlikely to be consumed for dry reforming, while it is not clear which reaction pathway prevailed for the secondary reaction of CO2 because of the complex reaction network [29]. However, the observation of BaCO3 in the BCZYYb anode powder treated in syngas implies that the resulting CO2 reacted with BCZYYb. Thus, the secondary reaction of CO2 with BCZYYb may have also contributed to the reduction in the normalized CO2 yield with increasing residence time, which would further support the more rapid degradation of BCZYYb cells with increasing anode thickness.Figure 6
Normalized CO2 and CH4 yields at 700°C over (a) BCZYYb and (b) BZY anode powders as a function of residence time.
(a)(b)


## 4. Conclusion

The present study clearly demonstrated that the electrochemical performance of a syngas-fueled PC-SOFC indeed depends on the PC oxide and anode thickness. This may be correlated with the residence time of the resulting product generated via the catalytic reaction of the H2–CO mixture in the microporous Ni-based anode. While CO is known to be an inert gas during the operation of PC-SOFCs, the presence of Ni in the anode clearly led to unfavorable side reactions such as CO disproportionation and methanation, which generated carbon coking over the metallic Ni surface and gas-phase CO2 and CH4 as products. Thus, it might be carefully suggested to utilize water in syngas fuel stream, or perovskite anode might be a solution to mitigate the carbon coking phenomena in the anode. The resulting gas-phase products may have caused further decomposition of the electrolyte materials in the cermet anode according to the residence time of the reactant in the microporous anode, which is correlated with the anode thickness. The catalytic reaction of the CO–H2 mixture is hypothesized to primarily take place on the metallic Ni surface. In addition, the less pronounced degradation observed in the BZY cells, despite the unchanged XRD pattern after treatment in syngas, is likely due to carbon coking via CO disproportionation or methanation. Further work is required to understand the exact nature of the active sites for the catalytic reaction of the CO–H2 mixture and to characterize the deposited carbon on the Ni surface.



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*Source: 9844833-2023-02-03.xml*