Novel layered solid oxide fuel cells with multiple-twinned Ni_0.8Co_0.2 nanoparticles: the key to thermally independent CO₂ utilization and power-chemical cogeneration

Abstract

To energy-efficiently offset our carbon footprint, we herein developed a novel CH₄–CO₂ dry reforming process to co-produce electricity and CO-concentrated syngas, which takes advantage of the selective oxidation of H₂ in high performance proton-conducting solid oxide fuel cells (SOFCs). In these cells, an additional functional layer, consisting of a Ni_0.8Co_0.2–La_0.2Ce_0.8O_1.9 (NiCo–LDC) composite, was successfully incorporated into the anode support, forming a layered SOFC configuration. The multiple-twinned bimetallic nanoparticles were then proven to have superior activity towards in situ dry reforming. In comparison to the conventional design, this layered SOFC demonstrated drastically improved CO₂ resistance as well as internal reforming efficiency (CO₂ conversion reached 91.5% at 700 °C), and up to 100 h galvanostatic stability in a CH₄–CO₂ feedstream at 1 A cm⁻². More importantly, H₂ was effectively and exclusively converted by electrochemical oxidation, yielding no CO₂ but CO concentrated syngas in the anode effluent. The maximum power density exceeded 910 mW cm⁻² at 700 °C with a polarization resistance as low as 0.121 Ω cm². Consequently, the heat released by H₂ electrochemical oxidation fully compensated for that required by the extremely endothermic dry reforming reaction, making the entire process thermally self-sufficient. We also showed that the layered design was beneficial in terms of decreasing coking and increasing CO₂ resistance of the SOFC in the mixed CO₂ and CH₄ feedstock. This novel process promises to play a pivotal role in future CO₂ conversion and utilization.

---

Broader context

Since the trend in the massive exploitation and utilization of fossil fuels has been increasing in step with the economic growth, the resulting unabated emission of greenhouse gases, primarily carbon dioxide, is having a serious negative effect on our ecosystem. It thus becomes urgent to find solutions that can offset our carbon footprint and mitigate climate change. Despite the increasing efforts that have been made, the issue still poses a great challenge, i.e., the emerging technologies tailored towards carbon sequestration via biological, chemical and physical processes are suffering from either low efficiency or the required high energy inputs. Herein, we have demonstrated an innovative way of consuming and utilizing CO₂via proton-conducting solid oxide fuel cells (SOFCs) with a layered structure. In lieu of consuming energy as in conventional CO₂ electrolysis and dry reforming, this novel system converts CO₂ to CO-enriched syngas with superior efficiency and simultaneously cogenerates electrical power with high power density. More interestingly, the entire process can be completely thermally-independent, requiring zero heat input. The produced syngas from the anode effluent can be readily utilized as the feedstock for downstream processes in chemical plants or other industrial areas, e.g., metal refinery.

1. Introduction

CO₂ is a notorious greenhouse gas (GHG) which is being increasingly emitted to our ecosystem as a result of various human activities. In response to the associated societal-environmental impacts ranging from global warming to climate change, how to effectively and efficiently sustain a reasonable atmospheric CO₂ level is becoming an emerging challenge worldwide.¹ Albeit the emission control is indeed a good solution, technology development aiming for CO₂ utilization is actually even more promising. In fact, reducing CO₂ into CO, the most important building block for a wide spectrum of organic syntheses,^1–3 has spurred huge interest in recent years.^4–6 One simple approach is the catalytic dry reforming of methane.^7,8

CO₂ + CH₄ = 2CO + 2H₂ (1)

However, coke formation on the catalysts and the large energy input limit its practical application.^9–11

Solid oxide fuel cells (SOFCs) show the potential as an ideal alternative performing CO₂–CH₄ dry reforming in situ in the anode chamber.^12–14 Unlike the conventional dry reforming process via the chemical route, the heat required for the extremely endothermic reaction can be fully counterbalanced by that cogenerated during the electrochemical oxidation, enabling a thermally autogenous process. Nonetheless, most SOFCs employed for dry reforming studies to date are based on oxygen-conducting electrolytes (O-SOFC), in which O²⁻ can readily oxidize both CO and H₂, the main products of dry reforming, leading to CO₂ emission again in practice.

H₂ + O²⁻ = H₂O + 2e⁻ (2)

CO + O²⁻ = CO₂ + 2e⁻ (3)

Another technical impediment associated is the insufficient reforming efficiency of this one-pot chemical/electrochemical reaction in SOFCs: in most cases, the conventional Ni-based cermet anode suffers from relatively low catalytic activity, causing low conversion of both CO₂ and CH₄.^15,16

Recently, many researchers have paid attention to proton-conducting SOFCs (H-SOFC) which are able to be operated at intermediate temperatures (500–700 °C).^17–22 Theoretically, H-SOFC renders higher efficiency compared with O-SOFC since H₂O generation occurs in the cathode without diluting the fuel in the anode.^23–25 In terms of CH₄ dry reforming, H-SOFC seems more advantageous than O-SOFC: as only protons can migrate through the electrolyte, no CO₂ can be electrochemically generated. Unfortunately, a serious technical problem arises if using H-SOFC for dry reforming: doped BaCeO₃, the most widely applied electrolyte of H-SOFC, has excellent proton conductivity; its chemical stability is, however, challenged in CO₂-containing atmospheres.^26–28 For instance, despite that the composition of doping elements has been well-optimized, phase decomposition of BaZr_0.1Ce_0.7Y_0.1Yb_0.1O_3−δ (BZCYYb) was still found in wet H₂ and CO₂.²⁹ Intuitively, this hinders the implementation of H-SOFC in CH₄ dry reforming in which ideally 50% CO₂ presents as shown in reaction (1).

Herein, we developed a layered H-SOFC incorporated with multiple-twinned Ni_0.8Co_0.2 nanoparticles. This catalyst showed greatly improved reforming activity as well as sintering-resistance, the novel SOFC configuration enables us to successfully achieve three milestones, i.e., CO₂ utilization, electricity generation and syngas production.

2. Experimental

2.1 Material preparation and cell fabrication

BZCYYb powders used in the anode and other portions of the cell were prepared by solid-state reaction (SSR) and sol–gel processes, respectively. The NdBa_0.75Ca_0.25Co₂O_5+δ (NBCaC) cathode was also prepared by the sol–gel process. The chemical compatibility of BZCYYb and NBCaC was investigated beforehand via calcining the mixture at 1000 °C in air for 10 h (see the X-ray diffraction patterns for more details in Fig. S1 in the ESI†). A Ni_0.8Co_0.2–La_0.2Ce_0.8O_1.9 (NiCo–LDC) catalyst was prepared by a glycine–nitrate auto-combustion process (GNP). The anode-supported cells with a configuration of Ni–BZCYYb/BZCYYb/NBCaC–BZCYYb were fabricated by using a sequential dry pressing/pre-calcination, spin coating and final sintering route. The dimension of the fuel cell was ∅ 16 × 1 mm with a cathode area of 0.316 cm². The NiCo–LDC catalyst was paste-coated on the outer surface of the anode support. Additional details are provided in the ESI.†

2.2 Reforming activity evaluation procedure

To compare the CH₄–CO₂ reforming activities of Ni–BZCYYb and NiCo–LDC, catalyst powders of NiO–BZCYYb (weight ratio is 65 : 35) and NiCoO–LDC were calcined at 1420 °C (5 h) and 900 °C (2 h), respectively. The higher calcination temperature of the former is to simulate the particle agglomeration effect during the fuel cell fabrication process. Then, the obtained powders were sieved using 30 to 60 mesh. 0.2 g of each catalyst was mixed with 0.4 g of catalytically inert SiO₂, and then packed into a quartz tube for activity measurements at temperatures from 700 to 550 °C after initial reduction in H₂ at 700 °C for 2 h. CH₄ and CO₂ were concurrently fed to the tubular reactor at a flow rate of 10 mL min⁻¹. The composition of the effluent gas was analyzed using a gas chromatography (GC, Hewlett Packard Series two). CO₂ conversion and CO selectivity were calculated by eqn (4) and (5), respectively.

2.3 Electrochemical performance evaluation procedure

During the SOFC test, Au and Pt pastes were painted on the anode and the cathode of the cells, respectively, as the current collectors, which were initially baked in air at 800 °C for 2 h. Pt meshes were used as the secondary current collector on top of those paste, which were connected with Pt wires. The cell was then sealed on the anode side to an alumina tube using a Ceramabond® glass sealant (Aremco Products, Inc.). After the sealant curing, the cell was reduced in H₂ at 700 °C for 2 h prior to electrochemical measurements. Open-circuit electrochemical impedance spectra (EIS) and the current density–voltage (I–V) curve were obtained at temperatures ranging from 550 to 700 °C in H₂, H₂–Ar, H₂–CO and CH₄–CO₂ (molar ratios of the mixed gases were 1 : 1), the fuel flow rate ranged from 20 mL min⁻¹ to 40 mL min⁻¹. Ambient air (relative humidity = (40 ± 10)%) was used as the oxidant. An impedance/gain phase analyzer (Solartron 1255) and an electrochemical interface (Solartron 1287) were used to record the electrochemical performances. Stability and syngas production tests were carried out via applying a constant current load, and simultaneously monitoring the voltage response as well as the effluent gas composition using a GC. H₂ selectivity during the electro-oxidation process was calculated by eqn (6).

2.4 Material characterization

All the powder materials were subjected to X-ray diffraction analyses (XRD, Rigaku Rotaflex) for phase identification by using Cu Kα radiation generated at 40 kV and 44 mA. The scan was conducted within a 2θ range between 20° and 80° at a rate of 1° min⁻¹. The NiCo–LDC was further analyzed using a transmission electron microscope (TEM, JEOL 2200 FS). The microstructures of the cells were examined using a scanning electron microscope (SEM, JEOL 6301F). X-ray photoelectron spectroscopy (XPS, Kratos AXIS Ultra) was used to study the surface chemistry of NiCo–LDC powder referenced to the adventitious carbon (C 1s with binding energy = 284.5 eV).

3. Results and discussion

3.1 Conventional state-of-the-art H-SOFC: incapable of efficient internal dry reforming

Conventionally, CO₂ utilization requires extremely high energy input via complex chemical or electrochemical processes. While producing syngas or other value-added chemicals, the conventional process unfortunately emits additional GHG and pollutants directly or indirectly as shown in Fig. 1. In this work, we propose to use an H-SOFC to achieve a “one-pot” green process for the conversion of CO₂, cogenerating power and a CO + H₂ mixture that can be used as the feedstock for various chemical syntheses.

Fig. 1 A conceptual comparison of the conventional and the proposed novel methods of CO₂ utilizations for syngas production.

The prepared Ni–BZCYYb/BZCYYb/NBCaC–BZCYYb cell demonstrated superior performances in H₂ feedstream, reaching 1480 mW cm⁻² maximum power density at 700 °C whereas the polarization resistance (R_P) was as low as 0.031 Ω cm². More details can be found in Fig. S2–S4 in the ESI.† By using this state-of-the-art H-SOFC, we then evaluated its feasibility for the internal dry reforming involving CO₂ and CH₄. In a control group, fuel cell performance in Ar diluted H₂ provides a benchmark to examine the possible negative influences of non-H₂ fuels (i.e. CO and CH₄) over H₂ electrochemical oxidation. Fig. 2a and b show the EIS and I–V(P) curves measured in H₂–Ar, H₂–CO and CH₄–CO₂ at 700 °C. The molar ratios of all mixed gases were fixed at 1 : 1. Remarkably, the R_P values of the cell in H₂–Ar and H₂–CO were essentially identical, 0.106 and 0.112 Ω cm², respectively, implying that CO did not affect the electrochemical activation and conversion of H₂ under this experimental condition. This in turn confirms the potential of using H-SOFC in syngas feedstream. Surprisingly, when CH₄–CO₂ was used as the fuel, which was supposed to yield syngas with the H₂–CO ratio close to unity, the polarization resistance of the cell increased substantially to 0.347 Ω cm² whereas the ohmic resistance (R_O) also grew to 0.201 Ω cm². Apparently, these results implied that both the physical and the chemical properties of the triple-phase boundary (TPB) have been altered drastically during internal dry reforming. Consequently, the maximum power density dropped to 562 mW cm⁻² and the cell suffered from degradations during a 12 h galvanostatic operation at 1 A cm⁻².

Fig. 2 Electrochemical performances of conventional H-SOFC in various fuels at 700 °C: (a) impedance spectra, (b) I–V and I–P curves, (c) galvanostatic test at a constant current of 1 A cm⁻²; and (d) XRD patterns of BZCYYb powder after 24 h exposure to either H₂–CO or CH₄–CO₂ at 700 °C.

Thus, it became extremely interesting to understand why the state-of-the-art H-SOFC was incapable of carrying out internal dry reforming. In accordance with the literature, the Ni–BZCYYb composite did exhibit excellent coking resistance in CH₄–CO₂, rendering few carbon deposits visible on the anode from the SEM observations (Fig. S5 in the ESI†). Coke formation is therefore unlikely to cause the lower performance in CH₄–CO₂. In the meantime, the ohmic resistance increase shown in Fig. 2a led us to explore if BZCYYb decomposition has occurred despite that Yang et al.²⁰ reported that these materials could be stable in a 50% CO₂-containing atmosphere. Our conjecture was proven correct conforming to the impurities appearing in the XRD patterns shown in Fig. 2d. In particular, BZCYYb has reacted with CO₂ after the exposure in CH₄–CO₂ at 700 °C for 24 h forming BaCO₃. SEM images in the ESI,† (Fig. S5–S7) also reveal substantial morphological changes of this electrolyte after the designated treatment. Accordingly, we now conclude that the electrochemical performance decrease in CH₄–CO₂ was ascribed solely to the decomposition of BZCYYb rather than the coke formation. Nevertheless, it is still encouraging to find that BZCYYb based SOFC showed excellent stability in syngas (see Fig. 2c), demonstrating the potentials of fueling H-SOFC with dry reforming derived syngas.

3.2 Multiple-twinned NiCo bimetallic nanoparticles: a superior reforming catalyst

Even in the most advanced H-SOFC discussed above, the low internal reforming efficiency brought a high concentration CO₂ stream in the anode causing the deterioration of BZCYYb. This problem makes the direct utilization of CH₄–CO₂ in H-SOFC challenging. Indeed, the development of a more stable electrolyte can be a straightforward solution,³⁰ but this usually undermines the high proton conductivity of the electrolyte materials. Alternatively, the design of an active dry reforming catalyst compatible with SOFC technology becomes critical.

In an attempt to tackle the obstacles, we herein developed a Ni_0.8Co_0.2 bimetallic alloy catalyst supported on La doped CeO₂ (La_0.2Ce_0.8O₂, LDC) via the GNP method. Although Ni exhibits sufficient catalytic activity towards methane dry reforming, it also fosters severe carbon deposition, leading to catalyst deactivation. Interestingly, alloying Ni with Co could simultaneously improve the catalytic activity while suppressing carbon formation.^31,32 Besides, the metal-oxide support interaction might be synergistically boosting the catalytic performance of Ni catalysts.³³ We thus tailored a CeO₂ support using the La₂O₃ dopant which stabilized the cubic structure of CeO₂ at high temperatures, enhanced the oxygen storage capacity and accelerated CO₂ absorption.³⁴

Fig. 3a shows the X-ray diffraction patterns of as-synthesized and reduced NiCoO–LDC, indicating the formations of solid solution in both the metallic and oxide phases. The EDX mappings (Fig. S8 in ESI†) show that the uniform distribution of the bimetallic was obtained in NiCoO–LDC. The dry reforming activity of NiCo–LDC was then compared with Ni–BZCYYb in fixed bed reactors from 550 to 700 °C (Fig. 3b). Noticeably, CO₂ conversion increased as a function of temperature for both catalysts. At 700 °C, the conversion over NiCo–LDC reached 96.9% which was more than 2-fold higher than that obtained using Ni–BZCYYb (48.3%). Meanwhile, the corresponding CO selectivity over NiCo–LDC was 97.6%, significantly surpassing that over Ni–BZCYYb (64.9%). The extremely high conversion of CO₂ on NiCo–LDC strongly supports its excellence for catalyzing dry reforming reaction.

Fig. 3 (a) XRD spectra for as-prepared and reduced NiCoO–LDC; (b) CO₂ conversion and CO selectivity comparison using Ni–BZCYYb and NiCo–LDC catalysts in methane dry reforming reaction at temperatures ranging from 550 °C to 700 °C.

To further understand the NiCo–LDC catalyst, we initially performed TEM analysis to acquire its compositional and morphological characteristics. Fig. S9 in the ESI,† shows the dark field image of a NiCo–LDC sample and the corresponding EDX elemental mappings of Ni, Co, Ce and La. Apparently, Ni and Co formed bimetallic nanoparticles (NPs) of size 15 to 30 nm and well-dispersed over cube-like LDC crystals. Interestingly, all of the observed bimetallic NPs appeared to have twinned structures that have been proven to possess fascinating catalytic properties.^35,36 We then conducted high resolution TEM (HRTEM) analysis to clearly observe the surface structure of the bimetallic NPs. The micrographs in Fig. 4b–d were obtained from the corresponding NPs shown in Fig. 4a. As indicated by the dotted lines in Fig. 4b–d, instead of forming spherical NPs, the NiCo bimetallic NPs had multiple-twinned facets. Stoichiometric amounts of Co and Ni were found on these facets (see Fig. S9 in the ESI†). Possibly, the addition of Co, along with the preparation method, introduced this surface reconstruction stemming from different atomic diffusion rates within the NPs.³⁷ Hence, these well-distributed NPs with multiple-twinned defects can be a key factor pertaining to the high catalytic activity and stability (coking and sintering resistances) of the catalyst.

Fig. 4 TEM analyses of NiCo–LDC catalysts, (a) BF micrograph; (b–d) HRTEM images of NiCo bimetallic nanoparticles at the corresponding sites labeled in (a), the dotted-lines indicate the twin boundaries; and (e) a schematic of a multiple-twinned nanoparticle.

In the control group, a catalyst with the same composition was prepared via the wet impregnation method. Although a Ni_0.8Co_0.2 bimetallic alloy also formed (Fig. S10 in the ESI†), it had different morphologies compared to the GNP counterpart. The spherical NPs suffered from severe agglomerations showing no twinning defects (Fig. S10 and S11 in the ESI†). They thus coherently exhibited much lower activity in dry reforming reactions (Fig. S12 in the ESI†). In addition, more CO₂ molecules can adsorb on the catalyst with twinned NPs relative to that obtained using the wet impregnation method (Fig. S13 in the ESI†).

The electronic properties as well as the surface chemistry of the NiCo–LDC hybrid were then studied via XPS measurements shown in Fig. 5. The survey spectrum shown in Fig. 5a confirms the presence of Ni, Co, La and Ce in the sample. The quantitative information from the XPS data reveals that the atomic ratio of Co/Ni is 17.1/82.9, well matching the chemical stoichiometry. Obviously, the surface was hydrated since a strong hydroxyl peak was observed in the O 1s spectrum (Fig. 5b). The spectrum of Ni 2p_3/2 shown in Fig. 5d suggests that the surface was (partially) covered by Ni(OH)₂. However, these hydroxide shells were unlikely to present under normal SOFC operating conditions. The electronic effect of the bimetallic alloy was determined through studying the Ni⁰ and Co⁰ binding energies shown in Fig. 5c and d. Usually, the 2p_3/2 orbitals of pure Ni and Co metals have binding energies of ∼852 eV and ∼778 eV, respectively.^38–40 When they form an alloy, the Ni⁰ of NiCo–LDC shifts to a lower binding energy, implying an increased electron density and electronegativity; in contrast, Co⁰ shifts to a higher binding energy. This electronic structure optimization by alloying was actually an effective way of increasing the catalyst activity towards methane dry reforming.^41,42

Fig. 5 XPS spectra of (a) all elements, (b) O 1s, (c) Co 2p, and (d) Ni 2p for the NiCo–LDC powder.

3.3 Layered H-SOFC: an ideal single reactor for multi-purposes

Intuitively, one would imagine that in H-SOFC, if majority of CO₂ is converted via methane reforming, the integrity of the BZCYYb electrolyte can be sustained since there is no direct contact between them. Fig. 6 is the schematic illustration of the novel layered SOFC. In this design, an additional reforming layer was incorporated onto the surface of the anode support. The CH₄–CO₂ feedstream is fully reformed on the high performance NiCo–LDC catalyst layer, yielding syngas, before entering the Ni–BZCYYb functional anode. The sequential selective oxidation of H₂ at TPBs not only generates electrical power and CO-enriched syngas, but also provides enough heat completely compensating that required by internal dry reforming. Based on the thermodynamic data shown below, we found that the thermal independency can be acquired when H₂ utilization is higher than 52% (corresponding to the total fuel utilization of 26%). This can significantly counterbalance the energy input for conventional methane reforming.

CH₄(g) + CO₂(g) = 2CO(g) + 2H₂(g), ΔH = 260 kJ mol⁻¹ (700 °C) (7)

2H₂(g) + O₂(g) = 2H₂O(g), ΔH = −495 kJ mol⁻¹ (700 °C) (8)

Fig. 6 A schematic showing the configuration of the novel layered H-SOFC.

Following this design, we fabricated the layered H-SOFC. Fig. 7a and b show the cross-sectional microstructures of the anode, electrolyte and cathode. The thickness of the dense BZCYYb electrolyte was ∼19 μm whereas that for the NBCaC cathode was ∼25 μm. The porous internal reforming layer is demonstrated in Fig. 7c. No delamination was observed in the interface region, ensuring efficient charge transfer during the cell operation. The EDX mapping results in Fig. 7d also indicate the uniform distribution of NiCo bimetallic NPs across the layer after cell fabrication.

Fig. 7 Cross-sectional SEM images of layered H-SOFC. (a) Microstructure of the NiO–BZCYYb/BZCYYb/NBCaC–BZCYYb triple layer; (b) magnified image of the NBCaC–BZCYYb cathode; (c) microstructure of the NiCoO–LDC/NiO–BZCYYb interphase; and (d) EDX elemental mappings for (c).

The internal dry reforming capability of the layered H-SOFC is compared with that of the conventional counterpart in a stream of CH₄–CO₂ in Fig. 8a and b. Both the ohmic and the charge transfer resistances were reduced substantially for the layered H-SOFC, simply owing to the improved BZCYYb stability and higher reforming efficiency. Consequently, its maximum power density was up to 910 mW cm⁻² at 700 °C. Fig. 8c and d contrast the time-dependent exhaust gas compositions for the conventional and layered H-SOFCs at 700 °C under open circuit voltage (OCV). Not surprisingly, the conventional H-SOFC showed much lower conversions of reactants and poorer stability. After only 6 h, the conversion dropped to below 20%, presumably because of the active site blockage in the anode arising from BZCYYb decomposition products and/or carbon deposition. Conversely, the NiCo–LDC incorporated cell sustained a superior reforming activity after 18 h test, and 90% of CO₂ has been reduced.

Fig. 8 Electrochemical performance comparison of conventional and layered SOFCs in CH₄–CO₂ at 700 °C for (a) impedance spectra and (b) I–V and power density profiles. Time dependent exhaust gas compositions for (c) conventional and (d) the NiCo–LDC incorporated SOFC at OCV and 700 °C when feeding with CH₄–CO₂.

We finally examined the electrochemical stability of the layered H-SOFC for cogenerations of power and chemicals. Fig. 9 displays the long-term performance of the cell under a constant current density load of 1 A cm⁻², the voltage stabilized at ∼0.8 V showing negligible degradations, and the power density was as high as 800 mW cm⁻². In the meantime, the gas composition of the anode effluent was examined every 10 h using GC. The H₂ selectivity stabilized at around 95%, implying that the layered cell was fully capable of accomplishing internal dry reforming in the anode chamber (note that BZCYYb shows some oxygen-ionic conductivity at 700 °C). Thanks to the reforming layer, no microstructure change or carbon depositions were observed in the Ni–BZCYYb anode after 100 h test (Fig. S15, ESI†). Additionally, we explored the adjustment of CO concentration in CO₂ derived syngas via electrochemical manipulations in the layered H-SOFC. The practical merit might be compelling: we can use a SOFC reactor to yield syngas with varied H₂/CO ratio upon downstream demands. To make the comparison clearer, we used syngas as the feed (without external CH₄ and CO₂ impurities). Fig. 10a shows H₂ selective oxidation in the cell powered by syngas with different H₂/CO ratios, in which a constant load of 2 A cm⁻² was applied and the anode effluent gas composition was sequentially determined. It can be seen that even when the H₂ inlet concentration was as low as 20%, the cell had a voltage of 0.26 V, giving a power density of 520 mW cm⁻². After electrochemical oxidation, the H₂ concentration was reduced to 10.6% whereas that of CO was enhanced to 88.8%, demonstrating the concept of simultaneously producing CO-enriched syngas using the layered SOFC reactor. Since H₂S is a common impurity in natural gas (source of methane), we also studied the sulfur resistance of the SOFC (see Fig. S16 and S17, ESI†). To have a well-defined pH₂S/pH₂, these tests were performed only in H₂S-containning H₂. The introduction of 50 ppm H₂S (typical concentration⁴³) caused an instant power density drop of the cell. However, during the subsequent galvanostatic measurement, a roughly constant voltage response was recorded, implying it is feasible for practical applications.

Fig. 9 Stability test of in situ dry reforming in layered SOFC fuelled with CH₄–CO₂ at 700 °C. (a) Galvanostatic test under a constant load of 1 A cm⁻² and (b) exhaust gas composition records.

Fig. 10 H₂ selective oxidation evaluations at 700 °C in layered SOFC fed with syngas containing various levels of H₂. (a) H₂ selectivity and the exhaust gas composition as a function of H₂ percentage in syngas and (b) the corresponding voltage responses in different feeds under a constant current load of 2 A cm⁻².

4. Conclusions

The present study demonstrated that the internal dry reforming of methane can be achieved using H-SOFC with a layered structure. The NiCo–LDC internal reforming layer, containing multiple-twinned Ni_0.8Co_0.2 bimetallic nanoparticles, effectively catalyzed the in situ reforming of methane, avoiding the direct contact of high concentration CO₂ with the BZCYYb proton-conducting electrolyte. Hence, the H-SOFC demonstrated excellent electrochemical performances in CH₄–CO₂ feedstream, coproducing CO-enriched syngas with few CO₂ by-products. The present work shed light on the configuration optimizations of SOFC in the context of energy-CO₂ conversions.

Acknowledgements

This research was financially supported by the Climate Change and Emissions Management Corporation, Alberta, Canada. N. Yan acknowledges the support from Sustainable Chemistry Research Priority Area of University of Amsterdam (http://www.suschem.uva.nl).

References

1. A. Appel, J. Bercaw, A. Bocarsly, H. Dobbek, D. DuBois, M. Dupuis, J. Ferry, E. Fujita, R. Hille, P. Kenis, C. Kerfeld, R. Morris, C. Peden, A. Portis, S. Ragsdale, T. Rauchfuss, J. Reek, L. Seefeldt, R. Thauer and G. Waldrop, Chem. Rev., 2013, 113, 6621–6658.
2. E. L. Muetterties and J. Stein, Chem. Rev., 1979, 79, 479–490.
3. C. K. Rofer-DePoorter, Chem. Rev., 1981, 81, 447–474.
4. J. Medina-Ramos, J. L. DiMeglio and J. Rosenthal, J. Am. Chem. Soc., 2014, 136, 8361–8367.
5. M. L. Macnaughtan, H. S. Soo and H. Frei, J. Phys. Chem. C, 2014, 118, 7874–7885.
6. C. Costentin, G. Passard, M. Robert and J.-M. Savéant, J. Am. Chem. Soc., 2014, 136, 11821–11829.
7. M. S. Aw, M. Zorko, I. G. Osojnik Črnivec and A. Pintar, Ind. Eng. Chem. Res., 2015, 54, 3775–3787.
8. P. Tang, Q. Zhu, Z. Wu and D. Ma, Energy Environ. Sci., 2014, 7, 2580–2591.
9. M. Honda, M. Tamura, K. Nakao, K. Suzuki, Y. Nakagawa and K. Tomishige, ACS Catal., 2014, 4, 1893–1896.
10. S. Chan, T. Lam, C. Yang, S. Yan and N. Cheng, Chem. Commun., 2015, 51, 7799–7801.
11. G. Olah, G. Prakash and A. Goeppert, J. Am. Chem. Soc., 2011, 133, 12881–12898.
12. W. Wang, C. Su, Y. Wu, R. Ran and Z. Shao, Chem. Rev., 2013, 113, 8104–8151.
13. S. E. Evans, J. Z. Staniforth, R. J. Darton and R. M. Ormerod, Green Chem., 2014, 16, 4587–4594.
14. K. Lee, C. Gore and E. Wachsman, J. Mater. Chem., 2012, 22, 22405–22408.
15. B. Hua, M. Li, B. Chi and J. Li, J. Mater. Chem. A, 2014, 2, 1150–1158.
16. Z. Zhan and S. A. Barnett, Science, 2005, 308, 844–847.
17. M. Li, B. Hua, S. P. Jiang, J. Pu, B. Chi and J. Li, Int. J. Hydrogen Energy, 2014, 39, 15975–15981.
18. B. Hua, M. Li, J. Pu, B. Chi and J. Li, J. Mater. Chem. A, 2014, 2, 12576–12582.
19. M. Li, B. Hua, J. Pu, B. Chi and J. Li, Sci. Rep., 2015, 5, 7667.
20. L. Yang, S. Wang, K. Blinn, M. Liu, Z. Liu and Z. Cheng, Science, 2009, 326, 126–129.
21. C. Zuo, S. Zha, M. Liu, M. Hatano and M. Uchiyama, Adv. Mater., 2006, 18, 3318–3320.
22. L. Bi, S. Boulfrad and E. Traversa, Chem. Soc. Rev., 2014, 43, 8255–8270.
23. E. Fabbri, L. Bi, D. Pergolesi and E. Traversa, Energy Environ. Sci., 2011, 4, 4984–4993.
24. E. Fabbri, L. Bi, D. Pergolesi and E. Traversa, Adv. Mater., 2012, 24, 195–208.
25. B. Rainwater, M. Liu and M. Liu, Int. J. Hydrogen Energy, 2012, 37, 18342–18348.
26. T. Hibino, A. Hashimoto, M. Suzuki and M. Sano, J. Electrochem. Soc., 2002, 149, A1503–A1508.
27. N. Zakowsky, S. Williamson and J. Irvine, Solid State Ionics, 2005, 176, 3019–3026.
28. R. Kannan, K. Singh, S. Gill, T. Furstenhaupt and V. Thangadurai, Sci. Rep., 2013, 3, 2138.
29. S. Fang, K. Brinkman and F. Chen, J. Membr. Sci., 2014, 467, 85–92.
30. C. Duan, J. Tong, M. Shang, S. Nikodemski, M. Sanders, S. Ricote, A. Almansoori and R. O'Hayre, Science, 2015, 349, 1321–1326.
31. D. San-José-Alonso, J. Juan-Juan, M. Illán-Gómez and M. Román-Martínez, Appl. Catal., A, 2009, 371, 54–59.
32. P. Djinović, I. Črnivec, B. Erjavec and A. Pintar, Appl. Catal., B, 2012, 125, 259–270.
33. S. Takenaka, H. Ogihara, I. Yamanaka and K. Otsuka, Appl. Catal., A, 2001, 217, 101–110.
34. X. Luo, R. Wang, J. Ni, J. Lin, B. Lin, X. Xu and K. Wei, Catal. Lett., 2009, 133, 382–387.
35. M. Kumar and S. Deka, ACS Appl. Mater. Interfaces, 2014, 6, 16071–16081.
36. J. Yang, L. Chng, X. Yang, X. Chen and J. Ying, Chem. Commun., 2014, 50, 1141–1143.
37. G. Li, Q. Wang, Y. Cao, J. Du and J. He, Phys. Lett. A, 2012, 376, 534–537.
38. K. S. Kim and N. Winograd, Surf. Sci., 1974, 43, 625–643.
39. C. Brundle, T. Chuang and D. Rice, Surf. Sci., 1976, 60, 286–300.
40. J. Kim, D. Pugmire, D. Battaglia and M. Langell, Appl. Surf. Sci., 2000, 165, 70–84.
41. J. Zhang, H. Wang and A. Dalai, J. Catal., 2007, 249, 300–310.
42. K. Takanabe, K. Nagaoka, K. Nariai and K. Aika, J. Catal., 2005, 232, 268–275.
43. Z. Cheng, J. Wang, Y. Choi, L. Yang, M. Lin and M. Liu, Energy Environ. Sci., 2011, 4, 4380–4409.

---

Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c5ee03017j

---