Recent advances in carbon-resistant anodes for solid oxide fuel cells

Abstract

Solid oxide fuel cells (SOFCs) can efficiently satisfy the power supply demand at any time and place in an environmentally friendly manner. However, considering the commercial application of SOFCs, the deposition of carbon on conventional SOFC anodes (nickel-based anodes) during their operation with hydrocarbon fuels, which leads to a degradation in the cell performance, needs to be well addressed. In this review, we discuss the carbon deposition process in SOFCs, carbon detection methods, and strategies to solve anode carburization, with a primary focus on alternative anode materials. Specifically, the coking mechanism, carbon-resistant strategies, and research development of bimetallic-cermet materials are reviewed in detail. In addition, the principle of in situ nanoparticle exsolution from perovskite materials and factors affecting the growth of dissolved nanoparticles are introduced, and the application of in situ exsolution in anti-carbon anodes is also discussed. Furthermore, we present the carbon deposition phenomenon in SOFC anodes and alternative anode design ideas in terms of simulation, calculation, and reaction kinetic models. Finally, we also give future research directions, especially proposing the potential application of single-atom based anode catalysts in hydrocarbon-fueled SOFCs.

---

Wei Zhang

Wei Zhang is currently a PhD candidate under the supervision of Prof. Chunwen Sun at China University of Mining and Technology (Beijing). He received his BS Degree in Applied Chemistry from Shandong University of Technology. His research interests include solid oxide fuel cells operating at intermediate and low temperatures, mainly focusing on designing anodes of bimetallic-cermet materials, in situ exsolution of perovskite materials and single-atom catalysts.

Jialu Wei

Jialu Wei is currently pursuing her MS Degree under the supervision of Prof. Chunwen Sun at China University of Mining and Technology (Beijing). She received her BS Degree in Applied Chemistry from China University of Mining and Technology (Beijing). Her research interests are mainly focused on developing anode materials for solid oxide fuel cells.

Fusheng Yin

Fusheng Yin is currently pursuing his MS Degree under the supervision of Prof. Chunwen Sun and Dr Zhijun Zhang at China University of Mining and Technology (Beijing). He received his BS Degree in Chemical Engineering and Technology from Jiangsu University. His research interests are first principles computational chemistry and all-solid-state batteries.

Chunwen Sun

Prof. Chunwen Sun received his PhD in Condensed Matter Physics from the Institute of Physics (IOP), Chinese Academy of Sciences (CAS) in 2006. He is a Full Professor and Group Leader of Energy Storage Materials and Devices at the School of Chemical and Environmental Engineering at China University of Mining & Technology (Beijing). His research interests include lithium-ion batteries, all-solid-state batteries, solid oxide fuel cells and electrocatalysis. He has published 150 peer-reviewed papers with citations of more than 14 000 times (Google Scholar), edited 7 book chapters and filed 22 Chinese patents.

---

1. Introduction

The sustainable development of modern society is hindered by the issues of energy shortage and environmental pollution. Therefore, solving or alleviating these problems has become a notable research direction. In this case, the development of renewable energy, the improvement of storage technologies and advanced energy conversion are essential to solve the above-mentioned challenges.^1–4

Fuel cells (FCs) are electrochemical energy conversion devices that directly convert chemical energy into electrical energy,⁵ which are highly efficient compared to other conventional energy conversion devices.^6,7 Further, FCs can operate in the opposite mode, that is, they can also convert electrical energy generated by modern renewable technologies such as solar and wind energy back into chemical energy.⁸ Therefore, FCs are a promising technology in the clean utilization of energy. FCs are desirable all-solid-state structures, i.e., there are no moving parts, and thus there is no risk of leakage in FCs. Consequently, these systems are expected to possess high reliability and long life. Moreover, unlike regular batteries, FCs allow arbitrary scaling between power (determined by fuel cell size) and capacity (determined by fuel storage size). Thus, based on their significant advantages, various types of FCs with different electrolyte materials have been developed.^5,6 Among them, proton-exchange membrane fuel cells (PEMFCs) and solid oxide fuel cells (SOFCs) have received significant attention owing to their relatively higher efficiency. Nevertheless, PEMFCs only use precious metals anode (e.g., Pt) and to prevent poisoning of Pt catalysts by CO, high-purity (∼99.99%) H₂ is required. All these restrictions greatly increase the production cost. In FCs, the chemical energy that is not converted into electricity will partially be converted into heat, but the high operating temperature of SOFCs (typically between 600–1000 °C) is different from that in other types of FCs, and hence the waste heat can be well used in SOFCs. Therefore, some of the energy of SOFCs can be recovered by introducing the reaction products into the steam turbine. SOFC/gas turbine hybrid systems provide combined heat and power with efficiencies of up to 90%.^6,7 Given that oxides and base metals are sufficiently active under high operating temperature, the use of precious metals in the electrodes can be avoided. Briefly, SOFCs have several advantages over PEMFCs, including relatively inexpensive materials (e.g., metal-ceramic and metal supported),^9–13 high fuel flexibility (e.g., hydrocarbon, syngas, solid carbon, and ammonia)^14–21 and very high efficiency.

However, although SOFCs exhibit excellent electrocatalytic properties when hydrogen is employed as the fuel, it is difficult to store and transport H₂ due to its low volumetric energy density. Alternatively, as mentioned above, SOFCs offer excellent fuel flexibility and can use hydrocarbon fuels at high operating temperatures. By avoiding the energy losses associated with external reformers, directing the electrochemical oxidation of CH₄ to generate electricity can significantly improve the efficiency (5–7%)²² and accelerate the applications of SOFCs in transportation and power distribution equipment.^5,23 Presently, a well-developed methane-rich natural gas infrastructure is available, which can be readily and easily used in residential or office buildings. Therefore, hydrocarbon fuels are considered a potential alternative to hydrogen fuels in SOFCs. However, due to the deposition of carbon caused by the incomplete oxidation of hydrocarbons and sulfur poisoning of the catalyst caused by common pollutants in the fuel, the catalytic activity of the most investigated anode materials of SOFCs, i.e., Ni–YSZ, is dramatically affected during operation with hydrocarbon fuels.^10,24–28

To achieve the large-scale commercial applications of SOFCs, many studies have been devoted to investigating carbon-resistant anodes, and some breakthroughs have been achieved in this regard. In this review, we explain the causes and types of carbon deposits in hydrocarbon-fueled SOFCs and summarize the relevant characterization methods. In particular, the methods for reducing carbon deposition on the anode based on thermodynamics and reaction kinetics are extensively reviewed, and the recent advances in carbon-resistant anodes are summarized in detail according to three aspects, as follows: (1) bimetallic-cermet materials, (2) ceramic materials, and (3) anode reforming layer materials. Furthermore, the outlook for some exciting approaches is discussed to provide some suggestions to broaden the research vision.

1.1 Working principle of SOFCs

An SOFC consists of an oxygen electrode (cathode), a fuel electrode (anode) and a dense ionic conductor (solid electrolyte) sandwiched between them, as well as an external circuit connecting the anode and cathode to collect electric energy, as shown in Fig. 1. During the working process, the electrodes do not directly participate in the reaction, they only act as a catalyst and provide ion and electron transport. If hydrocarbons are used as fuel gas for SOFCs, and assuming that the fuel is completely oxidized during operation, then the following reactions will occur spontaneously at the cathode and anode, as follows:

Fig. 1 Schematic diagram of a hydrocarbon-fueled SOFC.

Cathode reaction:

O₂ + 4e⁻ → 2O²⁻ (1)

Anode reaction:Total reaction:

The cathode is in contact with oxygen, and molecular O₂ is reduced to oxide anions (O²⁻) using electrons from the anode. Subsequently, the O²⁻ ions travel through the electrolyte and enter the anode, where they oxidize fuels to produce water and carbon dioxide, finally releasing electrons at the anode. The resulting electrons pass through the external circuit and return back to the cathode. When fuel and oxidizer are abundant, the SOFC can continue to supply power to the external circuit.

The ideal cathode electrode is usually a hybrid conductor that conducts ions and electrons. This material should have high oxidation resistance and high catalytic activity in the cathode environment, which can easily dissociate molecular oxygen into O²⁻ ions. The most commonly used material for the cathode electrode in SOFCs is Sr-doped LaMnO₃ (LSM), but it exhibits relatively poor electrical performance at temperatures below 800 °C because there are fewer oxygen vacancies. Consequently, some alternative cathode materials have been developed,²⁹ including La_0.8Sr_0.2CoO₃ (LSC),³⁰ La_0.8Sr_0.2FeO₃ (LSF),³¹ La_0.8Sr_0.2Co_0.2Fe_0.8O₃ (LSCF),³² Ba_0.5Sr_0.5Co_0.8Fe_0.2O_3−δ (BSCF),³³ PrBa_0.5Sr_0.5Co_1.5Fe_0.5O_5+δ(PBSCF),³⁴ Ba_1−xPr_xCo_1−yFe_yO_3−δ(BPCF),³⁵ Sr_0.95Ce_0.05CoO_3−δ(SCCO),³⁶ and Ba_0.9Co_0.7Fe_0.2Nb_0.1O_3−δ(BCFN).³⁷

There are two charged substances in the electrochemical reaction of SOFCs, i.e., electrons and ions. In the case of electrons, their transfer from the region they produced to the region they consumed is relatively easy. In contrast, for ions, their transport is much more difficult, mainly because they are much larger and heavier than electrons, and thus an electrolyte is necessary to provide a path for ions to transport. Ions move through the electrolyte via a “jump” mechanism, specifically, ions “jump” from one position to another in the lattice of the electrolyte material and the jumping process only occurs in places such as vacancies and gap-filling ions, which is prompted by lattice defects. Therefore, the conductivity of the electrolyte is controlled by the defects in the lattice. Generally, the intrinsic defect concentration of the electrolyte is low, and thus the non-intrinsic defect concentration is introduced in the electrolyte lattice by doping. The most commonly used electrolyte material is yttria-stabilized zirconia (YSZ), which shows excellent stability in both highly reducing and highly oxidizing environments, but it has low ionic conductivity. Therefore, some other doped oxides, such as scandia-stabilized zirconia (ScSZ), gadolinium-doped ceria (GDC) and samarium-doped ceria (SDC), are being investigated because they have higher ionic conductivities and good compatibilities with fuel electrode materials.³⁸

The ionic conductivity of the electrolyte is also closely related to the temperature, and thus the appropriate operating temperature of SOFCs depends on whether sufficient ionic conductivity is achieved in the electrolyte at this temperature.³⁸ Here, even in the case of excellent electrolytes, their ionic conductivity is usually 4–8 orders of magnitude lower than the electronic conductivity of metals, and hence the area-specific resistance (ASR) of an SOFC is mainly determined by the resistance of the electrolyte. Consequently, ion transport can cause significant resistance loss, which can degrade the performance of the SOFC. Thus, to attenuate this effect, the electrolyte should technically be as thin as possible to shorten the ion conduction pathway. There are many commonly used methods to prepare thin electrolytes, such as tape casting,^39,40 spin coating,⁴¹ wet powder spraying⁴² and dip coating.^43–45 Among the processes employed for the fabrication of thin films, magnetron sputtering⁴⁶ and pulsed laser deposition⁴⁷ have emerged as unique technologies to obtain oxide films by selecting appropriate process parameters. Furthermore, the electrolyte must be dense enough to separate the oxygen and fuel, while being thin, and it must be an electronic insulator. In addition, given that the partial pressure of O₂ varies from ∼1 atm at the cathode to ∼10⁻²⁰ atm or less at the anode, the electrolyte needs to have appropriate mechanical strength.⁴⁸ When SOFCs operate at high temperatures, their materials will expand, producing cracks. Therefore, each component of SOFCs needs to possess a similar thermal expansion coefficient (TEC) to prevent the formation of cracks, and their TEC compatibilities are usually ensured by mixing with the electrolyte material.

Anode materials usually have high conductivity, good catalytic activity, and suitable adsorption and desorption of reactants. The ability of the anode layer to catalyze the oxidation of fuels has always been the focus of SOFC research. To maximize the reaction surface area, anodes with porous micro/nano-structure need to be prepared to achieve close contact between the gas (gas phase) and electronic (electron-conducting phase) and oxygen ion (ion-conducting phase), and these reaction sites are called the three-phase boundary (TPB). The TPB concept is an important guide for the optimization of anode and cathode microstructures. It should be noted that this simplified description of the TPB does not consider the fact that some oxides, which are often added to the anode, have mixed electronic and ionic conductivity (MEIC), where there addition of a MEIC oxide will expand the TPB.

Optimizing the thickness of the anode catalyst layer requires a fine balance between mass transport and catalytic activity, given that thinner layers facilitate better gas diffusion and catalyst utilization, whereas thicker layers contain a great catalyst load and provide more TPB regions. The most commonly used material for the anode electrode in SOFCs is a nickel and YSZ composite (Ni–YSZ), which is a ceramic and metal mixture. Nickel provides electronic conductivity and catalytic activity, while YSZ provides ionic conductivity, thermal expansion compatibility comparable to the electrolyte, and maintains the high porosity and large surface area of the anode structure.⁴⁹ However, in addition to the above-mentioned carbon deposition problems in Ni-based anodes, cell breakage caused by the sintering of nickel particles during high-temperature operation of SOFCs and significant volume changes during re-oxidation process will lead to the interruption of the partial electron conduction network and failure of the catalytic active sites in the anode. Furthermore, carbon deposition, nickel agglomeration and poor redox cycling all block the reaction sites at the TPB, and consequently the performance of Ni-based anode SOFCs degrades.^50,51 However, thus far, no anode catalyst has been able to match the competitive price of the Ni cermet anode with good catalytic performance and long-term stability. Fortunately, there numerous studies have been conducted to address the above-mentioned problems, where among them, the development of carbon-resistant anode catalysts is one of the most promising methods to achieve high SOFC efficiency.²⁴ Carbon-resistant anode materials will be described in detail later.

The performance of fuel cells can be summarized by a current–voltage characteristic diagram (I–V), where it is critical to maintain high voltages under high current loads of fuel cells. Unfortunately, due to the unavoidable losses, the actual voltage output of fuel cells is always lower than that predicted by thermodynamic theory. In general, there are three main voltage losses, which can be expressed as follows:

V = E_thermo − η_act − η_ohmic − η_conc (4)

where V represents the operating voltage of the SOFC; E_thermo represents the thermodynamic-predicted voltage of the SOFC; η_act represents the activation loss caused by the reaction kinetics; η_ohmic represents the ohmic loss caused by ion resistance and electronic resistance; and η_conc represents the concentration loss caused by mass transport.

The activation loss is the overpotential required to exceed the half reaction energy barrier. At low operation temperature (e.g., PEMFCs), the output voltage drops sharply due to activation loss. However, at high operating temperatures in SOFCs, this problem does not exist. The main factors that affect the voltage drop in SOFCs are ohmic loss and concentration loss.⁵²

1.2 Challenges in hydrocarbon-fueled SOFCs

As mentioned above, the high operating temperature of SOFCs is actually in the temperature zone where the direct internal reforming (DIR) reaction of hydrocarbon occurs, thus opening up the possibility of using hydrocarbon as fuel. Hydrocarbon fuels undergo a variety of basic reactions under the operating conditions of SOFCs, and the elementary reactions are usually divided into the following reactions:^14,53–57

methane steam reforming:

CH₄ + H₂O → CO + 3H₂ (5)

dry methane reforming:

CH₄ + CO₂ → 2CO + 2H₂ (6)

auto methane thermal reforming:

2CH₄ + O₂ + CO₂ → 3H₂ + 3CO + H₂O (7)

partial oxidation:direct methane oxidation:

CH₄ + 2O₂ → CO₂ + 2H₂O (9)

water–gas shift reaction:

CO + H₂O → CO₂ + H₂ (10)

Considering the commercial application of SOFCs, the problem of carbon deposition during their operation using hydrocarbon fuels on state-of-the-art anodes (nickel-based anodes) needs to be well addressed.

Basically, during methane-fueled SOFC operation, three types of reactions are considered as sources of carbon deposition, as follows:⁵⁸

Methane cracking:

CH₄ → C + 2H₂, ΔH = +19 kJ mol⁻¹ (11)

Reduction of carbon monoxide:

CO + H₂ → C + H₂O, ΔH = −131 kJ mol⁻¹ (12)

Boudouard reaction:

2CO → C + CO₂, ΔH = −172 kJ mol⁻¹ (13)

Carbon can be formed by the thermal cracking of hydrocarbons or the decomposition of CO.⁵⁹ Given that the Gibbs free energy of the Boudouard reaction is the largest, it is generally considered to be the main carbon-forming reaction in research. When any one or more of these reactions occur, the following carbon deposition formation processes will appear. A schematic of the process on the anode catalyst (e.g., Ni metal) is shown in Fig. 2, consisting of the following steps:^13,60 in stage 1, at the beginning of any one of the above-mentions reactions, tiny amounts of carbon begin to be deposited on the active catalytic site, but it does not cover the entire anode surface. In this process, the current density of the SOFC increases slightly, and hence the deposited carbon can improve the conductivity of the anode. In stage 2, the carbon deposited on the active catalytic site is interconnected into layers, which will cover the entire anode surface. Considering that the deposited carbon is porous, fuel gases are still allowed to pass through and react on the anode surface, but the slowly formed coverage of deposited carbon can reduce the reaction TPB region and hinder the reaction rate. Therefore, at this stage, the current and voltage will decrease as time progress. In stage 3, the deposited carbon forms a thick layer on the anode and prevents fuel gas from penetrating into the TBP region to react, severely impeding the reaction rate. Consequently, with time, the output power of the SOFC will slowly change to zero.

Fig. 2 Schematic of the carbon deposition processes.

In the process of carbon deposition, generally three types of carbon species exist^14,61,62 including amorphous carbon (Type I), crystalline graphite form (Type II) and injected carbon (Type III). As illustrated in Fig. 3, if the carbon cannot be removed from oxidized form, Type I carbon first appears on the surface of the catalyst active sites. The high-resolution transmission electron microscopy image (HR-TEM) of Type I carbon in the TEM grid is shown in Fig. 4(a).⁶³ As the reaction proceeds, a large amount of carbon continues to be added to the active sites, and the Type I carbon will go through two different pathways,⁶⁴ as follows: (1) Type I carbon automatically undergoes a graphitization reaction upward to become Type II carbon, depending on the various catalysts and Type II carbon will behave as varying forms of carbon materials, such as carbon flakes, carbon fibers and carbon nanotubes. The scanning electron microscopy (SEM) images of Type II carbon are shown in Fig. 4(b) and (c), where the deposited carbon appears as carbon nanofibers at the interface between Ni and the oxide particles on the anode, and Fig. 4(c) is a magnified image of Fig. 4(b) by ten times.⁶⁵ (2) Type I carbon injects downward into metal crystalline structure, causing the structure to form Type III carbon (metal carbides), resulting in morphological changes in the metal anode structure, causing it to expand and peel off.⁶⁶ As shown in Fig. 4(d), the SEM image of the cross-section of the anode demonstrates that a large number of carbon species grew inside the cell. Also, this is further proven in Fig. 4(e) and (f), where the size of the Ni particles is in the nanometer range and separated from the zirconia by a carbon layer, which confirms that Type III carbon was formed.⁶⁷ In addition, Type II carbon can be directly transformed into Type III carbon, as shown in Fig. 4(g), where type II carbon is slowly converted to Type III carbon, and some of the Ni particles were stripped off the anode after 7 h.⁶⁸ Given that Type III carbon is an irreversible structure and cannot be removed in the operating environment of the SOFC, this type of carbon species is extremely destructive to the catalytic performance and structure of the anode.

Fig. 3 Schematic of three types of carbon species.

Fig. 4 Characterization of different types of carbon species. (a) HR-TEM image of Type I carbon. (Reprinted from ref. 63 with permission from the American Chemical Society.) (b and c) SEM images of Type II carbon. (Reprinted from ref. 65 with permission from Elsevier Ltd.) (d–f) SEM and TEM images of Type III carbon. (Reprinted from ref. 67 with permission from Elsevier Ltd.) (g) TEM images of the formation process with Type III carbon. (Reprinted from ref. 68 with permission from The Electrochemical Society Ltd.)

Compared with Type II and Type III carbon, Type I carbon can be more easily removed by oxygen ions, and hence a better carbon tolerance anode catalyst should allow carbon species to form Type I than the others.^54,69

1.3 Characterization of carbon deposition

In the design of a favorable carbon-resistant anode, the carbon deposition phenomena must be studied in detail by using suitable characterization methods. In this case, SEM is the most direct method for characterization, clearly showing the morphological structure of the carbon deposition existing on the surface of the anode. Additionally, when combined with energy dispersive X-ray spectroscopy (EDS), the presence of carbon that is not visible to the naked eye (e.g., Type III carbon) can also be detected. However, whether it can be detected depends on the depth of the electron beams of EDS and the depth at which carbon deposition is present at the anode, thus it is best to scan the cross-section of the SOFC when performing SEM-EDS characterization.

Type III carbon can cause volume expansion of the metal layer and change the lattice parameters, resulting in an increase in the crystal plane spacing of the material, and this X-ray diffraction (XRD) can be used to detect the presence of Type III carbon. Bayer et al. used an ionized magnetron sputter deposition system to deposit metastable Ni₃C nanocomposite films in an amorphous carbon matrix and studied the changes in their composition and structure at both low annealing (300 °C) and high annealing (800 °C) temperatures. It can be seen from Fig. 5(a) that in the annealing process, the Ni₃C phase slowly changed to a metallic Ni phase. In addition, the shift in the Ni₃C diffraction peak to the right as the Ni diffraction peak also proved that the crystal plane spacing of the Ni metal decreased due to the decomposition of Ni₃C.⁷⁰ XRD can also be used to detect the existence of Type II carbon with slower and smaller scanning steps, but a detailed carbon database is required to match the detection data.⁷¹

Fig. 5 Characterization methods for carbon deposits. (a) XRD patterns showing the modification in the diffraction peak of Ni₃C nanocomposite films in an amorphous carbon matrix with temperature and time. (Reprinted from ref. 70 with permission from the American Chemical Society.) (b) Thorn-ring characteristics in HRTEM-SAED of amorphous carbon. (Reprinted from ref. 72 with permission from MYJoVE Corporation.) (c and d) Lattice pattern characteristics in HRTEM-SAED of crystalline carbon. (Reprinted from ref. 73 with permission from Elsevier Ltd and from ref. 70 with permission from the American Chemical Society.) (e) Raman spectra showing that nickel can catalyze the graphitization of amorphous carbon with an increase in the annealing temperature. (Reprinted from ref. 70 with permission from the American Chemical Society.) (f) Peak deconvolution of TPO curve can confirm the presence of amorphous carbon, carbon fiber and graphite. (Reprinted from ref. 76 with permission from the American Chemical Society.) (g) In situ FTIR spectra quantifying the composition of outlet gases from the SOFC. (Reprinted from ref. 77 with permission from Elsevier Ltd.) (h) SVUV-PIMS spectra of gas-phase components for humid CH₄ reacting over NiCu–0.025MgO–SDC and NiCu–SDC catalysts. (Reprinted from ref. 78 with permission from Elsevier Ltd.) (i) In situ TEM showing the sustainable graphite layer deposition phenomenon on NiO and MgO-modified NiO catalysts. (Reprinted from ref. 79 with permission from the American Chemical Society.).

HR-TEM and selected area electron diffraction (SAED) can determine the presence of all three types of deposited carbon. HRTEM-SAED can use electron diffraction analysis through selected microregions to identify the crystal structure of the material, which allows the detection of the type and crystallinity of the material. Amorphous materials (Type I carbon) will exhibit a halo pattern, as shown in Fig. 5(b), while amorphous carbon shows Thorn-ring characteristics in the HRTEM-SAED image.⁷² Alternatively, crystalline materials (Type II and Type III carbon) show a spot pattern in the SAED image. As illustrated in Fig. 5(c-i), carbon nanotubes were observed under low-magnification HR-TEM, and in the close-up image in Fig. 5(c-ii), the carbon nanotubes are highly crystalline; moreover, the insert picture in Fig. 5(c-ii) displays the lattice pattern of the crystalline structure of the carbon nanotube (Type II carbon).⁷³ In addition, in the study by Bayer et al., the HRTEM-SAED pattern demonstrated the successful deposition of Ni₃C (Type III carbon) in the amorphous carbon matrix, as shown in Fig. 5(d).⁷⁰ Further, by calibrating the diffraction spots and comparing them with the XRD standard card library, the detailed information of the crystalline materials can be obtained.^63,73,74

Carbon species display two main peaks at 1360 cm⁻¹ (D band) and 1575 cm⁻¹ (G band) in Raman spectroscopy. The G band is often attributed to the presence of crystalline graphite (Type II carbon),whereas the D band is attributed to amorphous species (Type I carbon). According to the enhancement in the signals of the D band and the G band in the Raman spectra in Fig. 5(e), Bayer et al. found that nickel could catalyze the graphitization of amorphous carbon at a high annealing temperature of 800 °C.⁷⁰ Also, the ratio of the intensity of the D and G band (I_d/I_g) is used to identify the degree of graphitization in the carbon deposits, where the magnitude of this value reflects the ability of the anode to resist carbon deposition, given that the presence of type I carbon has the least impact on the anode performance among the carbon species.⁷⁵

The majority of deposited carbon except Type III in the anode can be removed as gas via oxidation by oxygen.⁷¹ Therefore, temperature-programmed oxidation (TPO) can be used to semi-quantitatively detect the carbon content and type in the SOFC anode by fitting the test curve. As shown in Fig. 5(f), Sumi et al. performed peak deconvolution of the TPO curve of the Ni–ScSZ anode after exposure to 10% CH₄-N₂ for 1 h, and the results indicated that amorphous carbon (Type I carbon), carbon fiber and graphite (type II carbon) were all present in the anode. Furthermore, comparing the size of the different peak areas, it is possible to estimate the percentage of different carbon types, and the peak area of C_α indicates the presence of a small amount of carbon fiber.⁷⁶

In situ detection can reflect the composition of the reaction intermediates during operation, which is of instructive significance for understanding the carbon deposition process and the anti-carbon deposition principle of the designed anode. Chlipała et al. described that in situ Fourier transform infrared spectroscopy (FTIR) can be employed to quantify the composition of outlet gases from SOFCs. Fig. 5(g) shows the outlet gas concentrations of CH₄, CO₂, CO, and H₂ over time, and the yield and selectivity of CO and H₂ could be calculated immediately given that the alterations in the concentration of individual gases can be monitored in real time by in situ FTIR. Furthermore, non-equilibrium analysis of in situ FTIR in a biogas fuel SOFC showed that methane steam reforming (MSR) was the most beneficial direct internal reforming reaction due to the presence of observable H₂O in the reaction atmosphere, and CH₄ cracking appeared to be the main reaction in the reaction, leading to carbon deposition, given that the carbon activity coefficient of this process was orders of magnitude higher than other reactions.⁷⁷ Xie et al. used synchrotron-based vacuum ultraviolet photoionization mass spectrometry (SVUV-PIMS) to further explore the impact of MgO on CH₄-involved reactions. As shown in Fig. 5(h), when measured at 700 °C and acquired at the photon energy of 11 eV, the SVUV-PIMS spectra demonstrated improved signals for C₂H₄ and C₃H₆ compared to the Ni–Cu catalyst without MgO, suggesting that the introduction of MgO improved the selectivity for C₂₊ hydrocarbons and inhibited CH₄ cracking to carbon directly.⁷⁸

Sun et al. studied the carbon formation mechanism on Ni-based catalysts using environmental transmission electron microscopy (ETEM) over a wide temperature range in combination with molecular dynamics simulations and density functional theory (DFT) calculations. In situ TEM observation performed in a C₂H₂/H₂ atmosphere provided real-time evidence that Ni₃C is an intermediate phase that decomposes to graphitic carbon and metallic Ni, leading to the formation of carbon. The mechanisms of acetylene decomposition and evolution of the carbon atom configuration were revealed by molecular dynamics simulations, which verified the experimental results. The modification of MgO on NiO could effectively decrease the formation of graphitic layers, and thus enhance the catalytic performance of NiO. The black arrows in Fig. 5(i) indicate the growth direction of the deposited graphite layer, and compared to NiO, the graphite layer in the MgO-modified NiO was basically unchanged with time and thinner in thickness. These results provide insight into the origin of the carbon deposition and developing effective approaches to mitigate it.⁷⁹

Each characterization has its unique strengths and drawbacks, and thus in specific SOFC anode research, a variety of characterizations should be rationally used to analyze the carbon deposits in combination with the existing scientific research conditions.

1.4 Strategies to control deposition of carbon

To prevent carbon deposition, two general methods can be considered. One is based on thermodynamics, while the other by changing the reaction kinetics. From a thermodynamic point of view, the first strategy to control carbon deposition is to reduce the operating temperature of SOFCs. Sumi et al. reported that methane cracking reactions mainly occur above 800 °C⁵⁸ and Steele found that at the lower operating temperature of SOFCs (below 750 °C), the carbon produced by CH₄ cracking will not deposit on most oxides.⁸⁰ Low temperatures are not conducive to carbon deposition in SOFCs, while the conductivity of oxygen ions in the electrolyte also decreases greatly, and thus this method is not ideal given that lowering the operating temperature will lead to a lower SOFC performance.

According to the study by Sasaki et al. on balancing the C, H, and O contents in fuels and their relationship with carbon formation, carbon deposition will occur at equilibrium.⁸¹ Thus, to avoid using an SOFC operating environment in the carbon deposition zone, the second strategy from a thermodynamic point is to introduce O atoms (oxygen and carbon dioxide) in the hydrocarbon fuel or add steam to the fuel gas to further improve the O/C and H/C ratios in the fuel composition. Thus, mixture fuels avoid being in the carbon deposition region and inhibit carbon formation. Carbon-containing adsorbents are susceptible to oxidation by the chemically adsorbed oxygen species present on the surface of metals, and therefore when there is enough steam, the carbon can be removed at a faster rate. However, due to the dilution of the fuel by a large amount of steam and the increase in p(O₂) in the fuel, the SOFC efficiency and the open-circuit voltage will decrease, and the rate of carbon deposition is proportional to the mass of the hydrocarbon fuel, i.e., hydrocarbons with longer chains require more steam. In addition, steam enhances the formation of Ni(OH)₂, and excess steam may also cause thermal stress to damage the SOFC components. Ideally, the minimum amount of steam required should be generated in situ by an electrochemical reaction to avoid any additional loss in SOFC efficiency.

The reaction kinetic methods to avoid carbon deposition include increasing the carbon removal rate or reducing the carbon deposition rate by adjusting the anode catalyst composition. Based on DFT calculations, the anodic oxidation pathways of CH₄ fuel in nickel-based materials are as follows: (1) CH₄ begins adsorption on Ni and (2) the adsorbed CH₄ decomposes into *CH₃, *CH₂, *CH and *H. However, for the adsorbed *H, it further migrates and overflows on the oxygen atoms in the electrolyte (O_electrolyte) to form adsorbed *OH, and the adsorbed hydrogen further overflows on *OH to form adsorbed *H₂O, which then desorbs to form H₂O. However, for *CH, it will have two different oxidation pathways, depending on whether it occurs on Ni(111) or Ni(211). (3a) In the case of Ni(111), *CH continues to be oxidized to *CHO by spilling O_electrolyte, (4a) *CHO is broken down into *CO and *H, and (5) in the case of *H, it follows a similar path as above, and *CO desorbs to CO or further oxidized to CO₂ at high temperature. In contrast, for Ni(211), the step followed is the same as (3b), where *CH is decomposed into *C and *H. Also, *C will further become carbon deposits on adjacent Ni surfaces.⁸²

Carbon-resistant SOFC anodes should be designed with the intrinsic adsorption capacity of CO₂ and H₂O. Thus, the anodes adsorb CO₂, greatly improving the in situ CO₂ reforming of methane or reacting with the carbon deposits at the anode to form gas-phase CO (C + CO₂ → 2CO), thereby removing the deposited carbon. In the case of anodes that can adsorb H₂O, the adsorbed *H₂O is converted to *OH with the oxygen vacancy in the material , while *OH can react with *C to form *CHO. Subsequently, *CHO follows similar steps as (4a) and (5). Both pathways will mitigate the carbon deposits typically encountered by nickel-based anodes in hydrocarbon fuels.^83,84

Given that nickel acts as an excellent dehydrogenation catalyst, this reaction manifests itself as coke formation when nickel-based anodes are used for hydrocarbon-fueled SOFC.⁸⁵ As an alternative, non-nickel-based catalysts have been developed.

Most oxide ceramics do not catalyze the formation of carbon in the same way as Ni, and some oxide anodes of SOFCs can directly oxidize methane given they are inert to carbon formation. Steele et al. used Bi₂O₃–Pr₆O₁₁ as an anode and proved that the metal oxide anodes did not tend to catalyze CH₄ into C₂ compounds. Therefore, CH₄ could be completely oxidized to CO₂ and H₂O, reducing the formation of coke.⁸⁶ Some other non-nickel based ceramic material anodes were studied, such as Sr_0.88Y_0.08TiO_3−δ forming yttria-stabilized zirconia (YST–YSZ),⁸⁷ lanthanum-doped SrTiO₃ (La_xSr_1−xTiO₃),⁸⁸ La_0.75Sr_0.25Cr_0.5Mn_0.5O₃ (LSCM),⁵¹ La_4.0Sr_8.0Ti_11.0Mn_0.5Ga_0.5O_37.5 (LSTGM)⁸⁹ and Sr₂MgMoO_6−δ(SMMO).⁹⁰ However, although the direct electrochemical oxidation of methane was achieved on the ceramic anode, the conductivity of the electrons was insufficient to give acceptable performance compared to the traditional Ni/YSZ anodes.

Metals commonly used as dehydrogenation catalysts and for the manufacturing of carbon nanotube, such as Ni, Co and Fe, will result in the deposition of carbon at the metal sites. Although the solubility of carbon in Cu, Ag and Au is lower than that of Ni, Fe, Co and will not result in the formation of substantial Type II or Type III carbon,^91,92 some researchers believe that Cu is a poor catalyst for hydrocarbon hydrogenolysis, and it is this property that endows Cu cermet with excellent coking stability. However, this that leads to some limitations. Firstly, Cu does not exhibit an acceptable electrocatalytic performance compared to Ni; furthermore, it has a low melting point and sinters easily, resulting in relatively poor thermal stability and making it incompatible with many standard high-temperature SOFCs manufacturing technologies; finally, it is prone to interface reactions with other SOFC components. Accordingly, one way to improve the activity and stability of Cu-based anodes is to alloy Cu with a second metal with higher catalytic activity, and from another perspective, a way to avoid Ni-based anode coking is to alloy Ni with low-activity metals such as Fe and Cu to inhibit the carbon deposition rate. Consequently, the use of alloy metals seems to be a good approach to solve the trade-off between electrochemical properties and carbon coking tolerance.

2. The development of hydrocarbon-fueled anodes

The development of novel anodes focuses on carbon-resistant materials that meet the critical requirements of hydrocarbon-fueled SOFC. Based on the keyword “SOFC hydrocarbon anode” search in Scopus, the corresponding number of papers published annually from 1995 to 2022 showed an upward trend. Overall, several types of materials are employed in the study of SOFC anodes for hydrocarbon utilization, as follows: (1) bimetallic-cermet materials, (2) ceramic materials, and (3) anode reforming layer materials.

2.1 Bimetallic-cermet materials

Cermet is a porous structure of ceramic-metallic composites. Excellent SOFC anodes should have mixed ionic-electronic conductors to catalyze electrochemical reactions, and therefore O²⁻ conductive ceramic materials such as YSZ and CeO₂ can be added to cermet. In addition, the cermet will help match the TEC of the anode with the electrolyte and control metal sintering.

The bifunctional synergy of bimetallic catalysts can modify the electronic properties and surface structure of the electrode (minimizing Ni particles and increasing TPB length), and additional the alloy reduces the carbon–metal bonding energy, thus increasing the tendency of carbon to be oxidized by oxygen ions.⁹³ It is well known that nickel, as a highly active dehydrogenation catalyst, can produce hydrogen and coke from hydrocarbon.⁸⁵ In the presence of carbon, due to the interaction between the 2p electrons of carbon and the 3d electrons of nickel, they will lead to the formation of nickel carbides, which are difficult to remove. If another metal is added to form an Ni-based alloy, the d-band center of the Ni atom can be moved away from the Fermi level, thereby reducing the formation of nickel carbides kinetically.

2.1.1 Nickel-based bimetallic cermet. Many researchers verified that the inhibition of carbon deposition can be achieved when copper is added to the anode. One way to improve the catalytic properties of Cu and deactivate the tendency of Ni to form carbon is the use of Ni–Cu alloys. Additionally, the coupling of metal Ni–Cu alloy with CeO₂ can increase the oxidation rate of hydrocarbons.

Kim et al. and Lü et al. began to study Ni–Cu alloys around the same time. Kim et al. studied an Ni–Cu–CeO₂–YSZ anode-supported SOFC for the direct oxidation of methane. Ni–Cu alloy composites divided into 20% Ni and 80% Cu were exposed to dry methane for 500 h, showing a significant increase in power density over time. The open-circuit voltage (OCV) was approximately 1.0 V in both H₂ and CH₄ at 1073 K, which is marginally lower than the theoretical OCV at this temperature, and the peak power density of the cell reached 440 mW cm⁻² for H₂ and 330 mW cm⁻² for CH₄. The experimental results of this work had many implications in the development of anodes in SOFCs. Firstly, a small amount of carbon can significantly lead to an obvious improvement in the performance of SOFCs, given it connects the isolated metal particles inside the anode, enhancing the electronic conductivity of the anode. This is the same as stage 1 of carbon deposition discussed earlier. Secondly, SOFC anodes based on Cu–Ni alloys have different properties compared to Ni-based cermet or Cu-based cermet, where the catalytic properties of an alloy are more than just the sum of the properties of its two separate metals. Overall, they demonstrated that a direct oxidation SOFC made of Cu–Ni alloy cermet was feasible for CH₄ at 1073 K.⁹⁴ Lü et al. investigated the performance of Ni_0.7Cu_0.3–YSZ/Ni_0.3Cu_0.7–YSZ cermet as an anode for SOFCs under hydrogen and city coal gas fuel, and the XRD results proved that the anode was a composite material. The OCV of the cell exceeded 1.1 V in H₂, which was very close to the theoretical voltage, but the output power density was very low, i.e., only about 33–43 mW cm⁻² at 900 °C. However, when changing the fuel from H₂ to city coal gas, there was only a slight drop in the output current and power density, and the anode reacted quickly to the fuel conversion without significant decay. When pitch was present in the city coal gas, tar-like carbon was only deposited on the Al₂O₃ tube of the cell, not on the anode, and hence the Ni–Cu–YSZ anode displayed a self-cleaning function. However, they did not further test the life of these anodes.⁹⁵

Song et al. prepared NiCu–CeO₂-based anode materials for intermediate-temperature SOFCs using hydrocarbon fuels and investigated the promotion effect of Ni–Cu alloy on their electrochemical activity and stability. Specifically, they focused on the power density and short-term stability tests of the SOFCs with six different anodes in 97% CH₄/3% H₂O fuel at 700 °C. The peak power density of the cells with the Ni–CeO₂, Cu–CeO₂ and NiCu–CeO₂ anodes was 60, 28 and 80 mW cm⁻², while that of the cells with the Ni–Zr_0.1Ce_0.9O_2−δ (ZDC), Cu–ZDC and NiCu–ZDC anodes was 92, 32 and 87 mW cm⁻², respectively. Short-term stability tests for all the cells at 700 °C were conducted at a constant voltage of 0.6 V under 97% CH₄/3% H₂O. After the stability test, the carbon content on the surface of the Ni–CeO₂, Cu–CeO₂ and NiCu–CeO₂ anodes was 15.62, 2.19 and 4.93 wt%, while that of the Ni–ZDC, Cu–ZDC and Ni–Cu–ZDC anodes was 10.72, 1.93 and 2.88 wt%, respectively. These experimental data show that compared with the pure CeO₂-based anodes, Zr_0.1Ce_0.9O_2−δ exhibited good thermal stability and activity. Moreover, the stability of Ni_0.5Cu_0.5–ZDC was comparable to that of Cu–ZDC, and its high activity is comparable to that of Ni–ZDC.⁹⁶

In the study of Ni–Cu alloy anodes, it was found that the addition of a small amount of Cu also resulted in a slight improvement in the stability of an SOFC in hydrocarbon fuels. Akdeniz et al. altered the anode structure by impregnating Cu-nitrate solution and Ce-nitrate solution in the Ni–YSZ backbone to improve the carbon-resistance performance of the SOFC under direct methane fuel. By controlling the variables of molarity of metal-nitrate solution; the amount of impregnation in mL (Cu-nitrate: Ce-nitrate); the amount of pore former and anode functional layer composition, they determined the most appropriate proportions of Ni–Cu–CeO₂–YSZ for the carbon-resistant anode. Specifically, the optimal dosages and molarities of copper-nitrate and cerium-nitrate solution were determined to be 3 mL–1 M and 2 mL–1 M, respectively. The Ni–Cu–CeO₂–YSZ anode-optimized SOFC showed a voltage loss of only 6% after 5 h of operation in dry methane fuel, and the carbon resistance and lifetime of the cell were significantly higher than that of traditional nickel-based anodes.⁹⁷

Ni–Cu alloy was also used in BaZr_0.1Ce_0.7Y_0.1Yb_0.1O_3−δ (BZCYYb) and Pr_0.5Ba_0.5MnO_3−δ (PBMO) anodes. Li et al. prepared Cu–, Ni– and Ni–Cu–BZCYYb anodes by impregnating metal salt solutions in a BZCYYb scaffold. BZCYYb is not a good electronic conductor, and thus metal oxides were introduced to form a good electron path; however, BZCYYb can absorb H₂O on its surface to increase the local H₂O content for carbon removal. The impregnation of Ni or Ni–Cu significantly reduced the ohmic resistance and polarization resistance of the BZCYYb anodes when exposed to H₂ or CH₄, while Cu impregnation only reduced the ohmic resistance in H₂ and CH₄, and furthermore only the Ni–Cu–BZCYYb anode was resistant to carbon deposition in wet and dry CH₄.⁹⁸ Kim et al. also synthesized Ni–, Ni–Cu–PBMO anodes in a PBMO scaffold using the impregnation method, which were expected to provide synergistic effects due the high catalytic activity of nickel and the high carbon coking tolerance of the transition metals. Subsequently, at 700 °C in the C₃H₈ fuel test, it was observed by SEM that the Ni–Cu–PBMO anode had no carbon deposition compared to the Ni–PBMO anode, even though the peak power density decreased from 280 to 200 mW cm⁻². In addition, they demonstrated through DFT calculations that the binding energy of the Ni–Cu alloy to C species was lower than that of Ni and calculations showed that the Ni–Cu–PBMO anodes had excellent tolerance to carbon deposition.⁹⁹

A second alloy metal can be introduced in the anode through different methods, which will improve the performance of the SOFC in some ways. By using a microwave irradiation process, Islam et al. directly deposited Cu particles on the Ni–YSZ anode in just 15 without a calcination step. SEM analysis confirmed that the morphology of the deposited Cu particles was spherical and that the average size of the particles was less than 100 nm. The electrochemical performance of the electrolyte-supported Ni–Cu–YSZ anode was improved over time when operating at 1073 K in dry CH₄ fuel at a current density of 20 mA cm⁻², which was similar to the performance trend of the Ni–Cu–CeO₂–YSZ anode prepared by Kim et al. using the impregnation method.⁹⁴ In the study by Kim et al., the weight ratio of Cu to Ni was 4 : 1, and the maximum power density (MPD) of the cell with the impregnated Ni–Cu–CeO₂–YSZ anode in methane at 1073 K was 330 mW cm⁻². Also, in the study by Islam et al., the weight ratio of Cu to Ni was only 0.06 : 1, but the microwave-irradiated Ni–Cu–YSZ anode had an MPD of merely 49 mW cm⁻² in methane at 1073 K. The difference in power density was due to the difference in the electrolyte thickness (60 μm vs. 300 μm) and the presence of CeO₂. Specifically, the microwave radiation process can be employed to prepare Ni–Cu–YSZ anodes in less time than other methods such as impregnation and electrodeposition.¹⁰⁰

Park et al. prepared an Ni–Cu–YSZ anode by plating Cu on a porous Ni–YSZ cermet anode using an aqueous copper sulfate solution. In methane fuel at 700 °C, the power density of the cell with the Ni–Cu–YSZ anode (240 mW cm⁻²) was slightly lower than that of the cell with the Ni–YSZ anode (280 mW cm⁻²), possibly due to the increase in anode resistance from 1.81 Ω cm² to 1.93 cm² by copper plating. However, the Ni–Cu–YSZ anode cell demonstrated a stable performance over 200 h, while the Ni–YSZ anode cell degraded dramatically within 21 h due to carbon deposition.¹⁰¹ Kumar et al. prepared an Ni_0.9–Cu_0.1–YSZ_0.95–GDC_0.05 anode by co-tape casting, followed by co-firing, where the advantage of this method is that it maintains the integrity of the anode-electrolyte interface. The optimized anode had an excellent power density of 436 mW cm⁻² at 850 °C in methane fuel. Carburizing studies in methane showed that the Ni_0.9–Cu_0.1–YSZ_0.95–GDC_0.05 anode had a tendency to electrochemically oxidize deposited carbon and reduce carbon deposition by up to 50%. Therefore, this study showed that the new anode composition and the preparation route employed can serve as a precedent for the development of high-power density SOFCs for hydrocarbon fuels.¹⁰²

Co has better oxidation resistance than Ni and does not exhibit corrosiveness at a high overpotential or high p(O₂).¹⁰³ Cho et al. synthesized Ni_1−x–Co_x–Ce_0.8Gd_0.2O_1.9 (GDC) anodes using a glycine nitrate process. It was found that the polarization resistance in an H₂ atmosphere increased with an increase in the Co content, whereas in a CH₄ atmosphere the polarization resistance decreased with an increase in the Co content. In CH₄ fuel at 800 °C, the MPDs of single cells of Ni–GDC and Ni_0.8–Co_0.2–GDC were 150 and 250 mW cm⁻², respectively. Consequently, they considered that Ni–Co alloy reduced the electronic conductivity but still alleviated carbon deposition by inhibiting the formation of C–C bonds and/or reducing the thermodynamic driving force required for carbon nucleation.¹⁰⁴ Nicharee et al. prepared Ni–YSZ and Ni–Co–YSZ by impregnating YSZ with nitrates of nickel and cobalt, followed by calcination and hydrogen reduction, and the introduced Co could be completely dissolved in the Ni lattice. Ni–Co alloys improve their electrochemical performance under CH₄ fuel by reducing the resistance and anode overvoltage. The MPDs of the cells with the Ni–YSZ and Ni_0.85–Co_0.15–YSZ anodes were 94 and 136 mW cm⁻² at 750 °C and at 20% v/v CH₄ in He, respectively. The ohmic resistance and polarization resistance of the Ni–YSZ and Ni_0.85–Co_0.15–YSZ anodes were investigated at 750 °C for 60 h at a constant loaded current of 80 mA, and the ohmic resistance of Ni_0.85–Co_0.15–YSZ basically did not change, and the polarization resistance increased slightly, but the overall performance was more stable than Ni–YSZ. Therefore, Ni–Co alloy showed high efficiency in inhibiting carbon deposition and improved the electrochemical performance of SOFCs operating with CH₄ fuel.¹⁰⁵

Cesario et al. successfully synthesized Ni–, Ni–Co– and Ni–Cu-based alloys with GDC by using citric acid as a chelating agent through a one-step synthesis route and utilized them in the dry methane reforming (DMR) reaction of SOFC. XRD analysis combined with thermodynamic calculations revealed that the anode formed a complete solid solution of Ni–Co and Ni–Cu-based alloys. Surface area (S_BET) measurements suggested that the Ni–Co–GDC samples had a higher surface area, which favored the expansion of the reaction TPB. In addition, according to the H₂-TPR results, the Ni–Co–GDC anode exhibited higher reducing and stronger metal-ceramic interactions compared to the Ni–Cu–GDC anode. This tight Ni–Co cermet contact allowed better adsorption of CO₂ on the anode and reduced carbon formation by inverse Boudouard reaction (CO₂ + C ⇌ 2CO). Finally, the DRM results demonstrated that the Ni–Co–GDC anode showed higher CH₄ and CO₂ conversion than the Ni–Cu–GDC and Ni–GDC catalysts, which could mitigate carbon deposition on the surface of the anode.¹⁰⁶

Alloying Fe with Ni was also studied as a viable option. Kan et al. synthesized an Ni_0.9–Fe_0.1–GDC anode using a solid-state reaction and the MPD of the cell reached 340 mW cm⁻² at 650 °C in dry methane, while the cell with Ni-GDC had a maximum power of 300 mW cm⁻². Subsequently, the long-term stability of the cell with the Ni–GDC and Ni–Fe–GDC anodes was evaluated at a constant current density of 0.2 A cm⁻² in dry methane fuel, where the performance of the SOFCs with the Ni_0.9–Fe_0.1–GDC anode did not degrade for 50 h, while that with the Ni–GDC anode only operated for 12 h. The exit gas of SOFC was analyzed by gas chromatography during the long-term stability test, and the results showed that the Ni–GDC anode produced 41.1% H₂, 19.6% CO, 8.1% CO₂ and 22.3% H₂O after 10 h. The high concentrations of H₂ and CO indicated the incomplete combustion of methane (CH₄ + O²⁻ → CO + 2H₂ + 2e⁻) and the ratio of C to H in the outlet stream was 0.88 : 4, which was lower than the expected ratio of 1 : 4 (CH₄ + 4O²⁻ → CO₂ + 2H₂O + 8e⁻). In contrast, the outlet stream of the Ni_0.9–Fe_0.1–GDC anode contained 33.2% H₂, 11.2% CO, 19.5% CO₂ and 27.8% H₂O after 10 h, and the ratio of C to H in the outlet stream was 1 : 4, indicating that no carbon deposition occurred inside the SOFC. Due to the addition of Fe, the changes in the catalytic performance resulted in more complete oxidation of the methane fuel, which enhanced the long-term stability of the Ni–Fe–GDC catalysts.¹⁰⁷ Kim et al. investigated the catalytic activity of Ni–, Ni–Cu–, Ni–Fe–PBMO anodes under C₃H₈ fuels. The Ni and Ni–Fe catalysts had the best electrochemical performance, while at 700 °C in C₃H₈ fuel, the MPDs of the cells with Ni–PBMO, Ni–Cu–PBMO and Ni–Fe–PBMO anodes were 280, 200 and 300 mW cm⁻², respectively. However, the Ni–PBMO catalysts formed carbon fibers under dry C₃H₈ operating conditions, while the Ni–Cu–PBMO and Ni–Fe–PBMO catalysts exhibited excellent carbon deposition tolerance. Considering the catalytic activity and carbon deposition resistance, the Ni–Fe–PBMO catalyst is best-suited for SOFCs based on the YSZ electrolyte support.⁹⁹

Lee et al. and Park et al. also prepared Ni–Fe alloys, but unlike other researchers, they added a layer of Ni–Fe alloy on top of the cermet anode layer, not within it. Specially, Lee et al. prepared a common Ni–YSZ anode, a dual-layered anode of Ni–Fe alloy layer and Ni–YSZ anode layer, and a hybrid double-anode (HB Ni–Fe) SOFC with a porous Ni–Fe alloy layer added to the dual-layered anode. In methane fuel tests in SOFCs, the cells with the Ni–YSZ and double- and triple-layer anodes exhibited current densities of 0.45, 0.59, and 0.66 A cm⁻² at 0.8 V and 750 °C, respectively. In addition, according to the impedance analysis, the cell with Ni–Fe additional layers exhibited lower charge transfer resistance than that without additional layers, which may be due to the reduced coke formation because of the catalytic activity of the Ni–Fe layer in methane fuel. However, they also proposed that the morphology and process of the SOFC system could be further optimized.¹⁰⁸ Park et al. did similar work to that of Lee et al. in the same year, and still concluded that SOFCs with an additional Ni–Fe layer in methane fuel had a greater current output and lower charge transfer resistance than that without the Ni–Fe layer, and the porosity and sintering temperature of the additional Ni–Fe layer need to be optimized.¹⁰⁹

For alloy anode research, the morphology of the Ni–Fe alloy is also an influencing factor that affects the cell performance. In high-temperature environments, metal spherical particles often agglomerate, thus affecting the microstructure of the electrode and cell stability. Thus, to overcome this problem, Lee et al. employed electrospinning to make Ni–Fe alloy nanofiber anodes, which were compared to Ni–Fe spherical powder anodes. The XRD analysis showed that Ni and Fe formed a solid solution under reducing conditions, and the SEM image displayed that the Ni–Fe nanofibers retained a porous network with an average pore size of 30 μm even after high-temperature sintering. Furthermore, the cell with the Ni–Fe fiber anode exhibited an MPD of 1640 mW cm⁻² at 1173 K in hydrogen fuel, while that with the Ni–Fe spherical powder anode had only 1450 mW cm⁻². At all the operating temperatures, the polarization resistance of the Ni–Fe nanofiber anode was lower than that of the Ni–Fe powder anode. These results demonstrate that the highly porous structure of the Ni–Fe nanofiber anode structure can produce excellent charge transfer paths and enhance the gas diffusion performance, thereby enhancing the electrochemical performance and stability of the SOFC.¹¹⁰ Also, Although they did not test the SOFCs in hydrocarbon fuel, it is expected that the Ni–Fe nanofiber anode SOFCs can perform well in anti-carbon deposition research.

Among the transition metals, Mo is an ideal nickel-based alloy metal. Hua et al. reported the preparation of a novel Ni–Mo bimetallic alloy anode, which when fueled with CH₄–50 ppm H₂S, the MPD of the cell with the Ni–Mo–YSZ anode was 594 mW cm⁻² at 800 °C, and this bimetallic catalyst also provided a consistently stable power output over a long period of 120 h, making it a promising anode catalyst with excellent coke resistance and thermal stability.¹¹¹

Bkour et al. used an Ni–Mo–YSZ catalyst for the partial oxidation of isooctane, which was an alternative to the commonly used gasoline. The XRD results showed that the larger Mo atoms partially replaced Ni atoms to form Ni–Mo solid solutions. The Raman spectroscopy results also indicated that the presence of well-dispersed Mo=O species on the Ni surface was an active site related to the high coke resistance of the Ni–Mo–YSZ catalyst. Further, they used the Ni–Mo–YSZ catalyst as a reforming layer for the traditional Ni–YSZ anode SOFC, with a maximum power output of 452 mW cm⁻² at 750 °C in isooctane/air operation. In the 15 h performance stability test, the cell with the micro-reforming layer showed a lower degradation rate, while the cell without the reforming layer stopped working within 12 h. This improved stability suggests that the use of an Ni–Mo–YSZ reforming layer in isooctane/air operation mode has great benefits for the electrochemical performance of SOFCs. In addition, DFT-based calculations demonstrated that the d-band center of the Ni atoms located close to the Mo atoms shifted away from the Fermi level, thus reducing the electron interaction between carbon and nickel. Consequently, less carbon was adsorbed in the Ni–Mo system.¹¹² Moreover, they previously studied the catalyst capacity of an Ni–Mo–CZ catalyst in isooctane, and the carbon conversion rate reached 100% and the H₂ yield reached 75%. When Ni–Mo–CZ was used as the reforming layer on the traditional Ni–YSZ-based SOFC, it still showed an excellent voltage degradation rate of 4.8 mV h⁻¹ after 12 h operation at 750 °C in isooctane fuel.¹¹³ Similarly, Hou et al. developed a dual-functional Ni–Mo–CZ as an anode for isooctane-fueled SOFCs. The MPD of the cell with the Ni–CZ anode was 70 mW cm⁻² at 800 °C in isooctane/air, indicating its very poor anode activity, where the MPD tripled to 212 mW cm⁻² after the addition of 5 wt% Mo to the Ni–CZ anode. Also, the SOFC with the Ni–5Mo–CZ anode could run steadily for 30 h in an isooctane/air mixture, which has great potential in SOFCs by using complex liquid isooctane fuels.¹¹⁴

Majewski et al. investigated the carbon tolerance of Ni–Mo alloys in steam reforming and dry reforming of methane combining five different ceramic materials, which were YSZ, GDC, CSZ, Al₂O₃ and TiO₂. The choice of ceramic materials significantly affected the reaction kinetics of the Ni–Mo catalysts and the yield of ideal products for H₂ and CO during CH₄ reforming. Also, in the process of methane reforming, the presence of Mo₂C was also detected in the Ni–Mo–YSZ and Ni–Mo–Al₂O₃ catalysts, which could improve the catalytic activity of the steam reforming reactions. In general, the catalytic performance of the Ni–Mo-combined ceramic catalysts decreased in the order of Al₂O₃ > YSZ > TiO₂ > CSZ > GDC.¹¹⁵ These findings can guide the design of Ni–Mo bimetallic anodes for carbon-resistant SOFCs.

In addition to transition metals, Sn plays a major role in the research of hydrocarbon fuels. There are many reports in the literature on Ni–Sn alloys. Nikola et al. combined kinetic studies, isotope labeling experiments and DFT calculations to make outstanding contributions to the study of Ni–Sn alloy catalysts for hydrocarbon fuels. In particular, according to the results of DFT calculations, compared with Ni–YSZ, the Sn–Ni–YSZ catalysts (0.5–1 wt%, with respect to Ni) showed that the surface of the Ni–Sn alloy preferentially oxidized C atoms, promoting the formation of C–O bonds and further forming gas-phase CO and CO₂, rather than forming C–C bonds. Further catalyst characterization studies showed that the presence of Ni–Sn–YSZ had a great effect on the reduction of carbon deposition for methane steam reforming, and X-ray photoelectron spectroscopy (XPS) and EDS studies indicated that Ni–Sn surface alloys were thermodynamically preferred structures at the limit of low Sn concentrations.^116–118 Their study demonstrated that Ni–Sn–YSZ catalysts were more resistant to carbon deposition than Ni–YSZ catalysts in isooctane steam reforming.¹¹⁹

Kan et al. investigated the application of Ni–Sn–YSZ catalysts as anodes in methane-fueled SOFCs. an MPD of 390 mW cm⁻² was obtained at 800 °C in methane fuel using the Sn-doped Ni–YSZ cell, which was equivalent to the results obtained with the Ni–YSZ cell (400 mW cm⁻²). Furthermore, in a stability test at 800 °C with a steady applied current density of 0.2 A cm⁻², the Ni–YSZ anode cell stopped operating after 2.3 h, while the Ni–Sn–YSZ anode cell could run for 49.0 h. Moreover, graphitic carbon was deposited on the Ni–YSZ anode, while amorphous carbon was deposited on the Ni–Sn–YSZ anode. It was demonstrated that on the Sn-doped Ni–YSZ, the transition from amorphous carbon to graphite was hindered.¹²⁰ Subsequently, they added a functional layer between the native anode and the electrolyte to obtain a high power density even at intermediate temperatures, and studied the performance of the SOFC with the (Ni–Sn–YSZ|Ni–Sn–YSZ|YSZ|LSM + YSZ) configuration in comparison to that with the (Ni–YSZ|Ni–Sn–YSZ|YSZ|LSM + YSZ) configuration in wet methane fuel. The cells with the added functional layer showed a high-power density output at an operating temperature of 650 °C in methane fuel, and the Ni–YSZ and Sn-doped Ni–YSZ functional layers displayed similar power densities (390 mW cm⁻²vs. 410 mW cm⁻²), while the Sn-doped Ni–YSZ showed enhanced long-term stability (27 h vs. 137 h). Similar to the results of their previous studies, this good stability was ascribed to the much lower amorphous carbon deposition rate. As mentioned above that amorphous carbon (Type I carbon) can be easily removed, they retested the long-term stability of the Sn-doped Ni/YSZ functional layer cells and removed the carbon deposits on the cell every 100 h by high-temperature oxidative calcination, and consequently the cells could run stably for 300 h with no observed degradation. The Sn content measured by energy dispersive X-ray spectroscopy (EDX) prior to stability testing was 5.2 wt% in the functional layer and 17.2 wt% in the native anode. After 300 h of operation, the change in Sn content was negligible. The above-mentioned results confirmed that Sn was stable on the Ni surface despite the many hours of operation.¹²¹

Jiang et al. separately synthesized 1 wt% Cu- and Sn-impregnated traditional Ni/YSZ anodes. In the biogas fuel test, the MPD of the cells with Ni–, Ni–Cu–, Ni–Sn–YSZ anodes was 101, 85 and 272 mW cm⁻² at 750 °C. At a constant current density of 0.2 A cm⁻² at 750 °C, it was found that the stability of the pure Ni–YSZ anode SOFC deteriorated rapidly, and the operation stopped after 19 h. Also, the voltage of Ni–Cu–YSZ dropped to 0.025 V at a fast rate after 48 h test, while the cell with the Ni–Sn–YSZ anode remained unchanged throughout the test period, and the voltage reduction rate was only 2.98 × 10⁻⁴ V h⁻¹. Furthermore, they used SEM to characterize the anodes, and observe severe carbon deposition on the surface of the Ni–YSZ anode, whereas no observable fibrous carbon was found on the surface of the Ni–Sn–YSZ anode. After measuring the morphology of the anode surface, the elements on the anode surface were quantified in weight percentage through EDX, where the content of C element in the Ni–, Ni–Cu–, and Ni–Sn–YSZ anodes was 5.32, 3.61 and 4.5 wt%, respectively. Hence, the long-term stability of the SOFCs under biogas fuel could be significantly improved by adding 1 wt% Sn to the Ni/YSZ anode.¹²²

Singh et al. prepared (1 wt%) Sn–Ni–YSZ anodes at a high sintering temperature (1450 °C) and evaluated their direct utilization of CH₄ at 800 °C. They found that Sn enhanced the shrinkage of the Ni–YSZ anode during the sintering stage of the preparation process, and this shrinkage matched better with the electrolyte shrinkage, which may help to reduce the stresses during the fabrication process. However, Sn was not detected in the anode after the sintering stage, resulting in the presence of Sn having no significant effect on carbon formation on the Ni–YSZ anode.¹²³ Farrell et al. showed that the carbon deposition catalyzed by Ni in ethanol fuel of the Ni–YSZ anode could be reduced by alloying nickel with a small amount of Sn. The Ni–Sn alloys had a lower tendency to activate carbon deposition formation compared to monometallic Ni and did not significantly reduce the reaction rate.¹²⁴

For the same alloy, the properties of various ceramic materials will be different. Better results were observed when Ni–Sn alloys were used with GDC and SDC than YSZ. Yoon et al. synthesized (1.99 atom%) Sn–Ni–GDC anodes using the impregnation method. In terms of the electrochemical performance of CH₄ fuel at 650 °C, the cells with Ni–GDC and Ni–Sn–GDC anodes showed MPDs of 370 and 470 mW cm⁻² at 650 °C, respectively. Hydrophilic Sn doped on the surface of nickel was hydrated and hydroxylated by reaction with water generated from electrochemical oxidation of fuels, and the hydroxyl groups (*OH) contributed to the oxidation of the adjacent carbon and removed as CO or CO₂. Also, the Ni–Sn alloy prevented the enlargement of Ni particles, which not only hindered the sintering of Ni metal, but also expanded the TPB of the anode, This enabled O²⁻ to quickly reach the carbon deposition area and electrooxidize it, thereby converting it into carbon dioxide and releasing it from the anode. In addition, Sn: [Kr]4d¹⁰5s²5p² with an external electron configuration analogous to C: [He]2s²2p² (two valence electrons close to stable s-orbital) hindered the formation of nickel carbide. Therefore, with this anti-carbon deposition advantage, the power generation time of the CH₄-fueled SOFC with the Ni–Sn–GDC anode could last for more than 200 h without any performance degradation.¹²⁵ Myung et al. developed Sn-doped nano-sized Ni and GDC conjugated on a core GDC nano-composite anode (n-SNGG), where this structure had more uniformly distributed Sn because Sn was doped on nanoscale NiO embedded in the core GDC. Due to its excellent micro-junction, the cell with the n-SNGG anode showed an MPD of 930 mW cm⁻² at 650 °C in methane fuel and could operate stably for more than 40 h without attenuation in its performance.¹²⁶ Yang et al. reported that in humidified CH₄, the MPDs of the cell with (1 wt%) Sn–Ni–SDC anodes reached 640, 390, and 280 mW cm⁻² at 800 °C, 750 °C, and 700 °C, respectively. Also, after running at 700 °C at a constant voltage of 0.8 V in wet CH₄ for 230 h, the current density dropped by only 30%. High-angle annular dark-field scanning transmission electron microscopy (HAADF STEM) images showed that the addition of a small amount of Sn to the surface of Ni–SDC favored the formation of highly active Ni₃Sn, which was one of the highly active catalysts for hydrocarbon oxidation. Moreover, during the preparation of the cell, they found that the key process of oxidizing impregnated SnCl₂ to SnO₂ at temperatures above 800 °C was essential.¹²⁷ Li et al. studied the performance of single SOFCs (Ni–Sn–SDC|SDC|Ba_0.5Sr_0.5Co_0.8Fe_0.2O_3−δ) in the temperature range of 600–700 °C. The cell with the (5 molar%) Sn–Ni–SDC anode exhibited an MPD of about 600 mW cm⁻² at 700 °C with CH₄ as fuel. Also, at a constant output current of 300 mA at 600 °C with CH₄ as the fuel, its performance degraded by only 1.5% during 72 h of operation, while the Ni–SDC anode dropped by 3.7%. The addition of Sn significantly inhibited carbon deposition on the anode, thereby improving the stability of the single cells fueled by dry methane.¹²⁸

Arifin et al.⁴⁴ reported that the electrochemical performance of Ni–Sn–ScSZ in simulated biogas (CH₄/CO₂ = 2) was better than that of undoped Ni–ScSZ. When the fuel changed from hydrogen to biogas, an immediate decrease in OCV and MPD was observed, with a decrease of 86.5% and 33.3% for the cells with Ni–YSZ and Ni–Sn–ScSZ anodes, respectively. However, they also found higher levels of carbon deposition in Ni–Sn–ScSZ.¹²⁹

Previous work suggested that a high Sn loading (Sn/Ni ratio >1 wt%) adversely affects the performance of SOFCs. Thus, the loading amount of Sn needs to be adjusted to a suitable ratio, which can maximize the catalytic performance of Sn without hindering the Ni active site and TPB.

By varying the amount of Sn impregnated in the functional layer, Kan et al. found that although a small amount of Sn significantly improved the cell performance, this enhancement decreased with an increase in the Sn content. This was because Sn is preferentially retained on the surface of the anode and the presence of excess Sn will occupy a high percentage of the surface-active sites, thus weakening the anode response.¹²¹ The study by Singh et al. showed that the electrochemical performance of SOFCs in H₂ and CH₄ decreased when the Sn content increased from 1% to 5%.¹³⁰ Jang et al. investigated that the impregnation of 15 to 30 mg Sn on an anode with a surface area of 3.14 cm² resulted in the highest power density, whereas 60 mg Sn reduced the power density. To better understand the observed phenomena, they also conducted model experiments. The modeling study showed that the oversupplied Sn agglomerated, resulting in fewer reaction sites, and therefore lower power generation.¹³¹

Troskialina et al. impregnated Sn with mass ratios of with 0.00, 0.20, 0.38, 0.66 and 1.02 wt% Sn/Ni in Ni–YSZ catalysts, respectively. Repeated cells tests showed that the electrochemical performance of the SOFC was very sensitive to the amount of Sn doped in the Ni/YSZ anode, and among them, the cell with an Sn/Ni loading of 0.38 wt% had the highest performance in both H₂ and biogas fuels. According to the performance trend, it was predicted that when the Sn loading is > 1.02 wt%, the presence of Sn has a significant negative impact on the electrochemical oxidation of H₂, which consequently affects the application in hydrocarbon fuels. This finding can explain the results reported by Singh et al.,¹³⁰ because the cells studied by Singh et al. were under high Sn loads (1 and 5 wt%), and the anode was negatively affected by Sn on hydrogen electrochemical oxidation, resulting in the positive effect of Sn on the Ni–YSZ anode being ignored.

Similar to nickel, noble metals such as Ag, Pd, Ru, Rh and Pt are the commonly used metals in SOFCs. Although silver wires are commonly used as current collectors, silver itself is known to improve the performance of SOFCs.¹³² Wu et al. impregnated Ag with mass ratios of 0.00, 0.89, 1.59 and 2.48 wt% in Ni–YSZ catalysts, which were denoted as blank cell, cell A, B and C, respectively. For dry CH₄, the MPDs of the cells at 750 °C were 542, 262, 307 and 271 mW cm⁻², respectively. However, in the stability test, the blank cell could work for only 5 h, while the modified cells could run stably for longer time without performance degradation. The SEM results showed that for cell B, the carbon deposition was greatly reduced both inside and on the surface of the anode, and thus Ag greatly improved the carbon deposition resistance of the Ni–YSZ anodes fueled by dry CH₄, which is mainly due to the fact that the deposited carbon could be easily separated by Ag into isolated pieces.¹³³ Jiang et al. also found that doping the Ni–YSZ anode with 1 wt% Ag reduced the MPD of the cell when operating under the condition of 50% CH₄–25% CO₂–25% N₂, but the long-term stability could be greatly improved.¹²² Wu et al. further improved the relevant works about introducing Ag in the Ni–YSZ anode by the electroless plating method, instead of using the impregnation method. Also, it was found that the MPD of the SOFCs with the Ni–Ag–YSZ anode increased whether in H₂, dry CH₄ or dry C₂H₆ fuel. In particular, the Ni–YSZ cell only operated for a few minutes when working in dry C₂H₆, while the Ni–Ag–YSZ cell could run for up to more than 24 h.¹³⁴

Hibino et al. studied (0–10 wt%) Ni–Pd–SDC as an SOFC anode in a mixture of methane and air fuels between 450 °C and 550 °C. They found that methane was greatly activated on the anode by adding a small amount of Pd (0.145 mg cm⁻²), which significantly contributed to the oxidation of methane by oxygen to form hydrogen gas and carbon monoxide. This result led to a high MPD of 644 mW cm⁻² at 550 °C when using an SDC electrolyte with a thickness of 0.15 mm.¹³⁵ Babaei et al. investigated that the presence of Pd in Ni–GDC greatly promoted the electrooxidation of hydrocarbon, especially methanol and ethanol fuels. Specially, in the methanol oxidation reaction, the activation energy of the reaction decreased from 154 to 101 kJ mol⁻¹ on a 0.15 mg cm⁻² PdO-impregnated Ni–GDC anode. Also, they found that although carbon depositions were present in the working environment of all these fuels, no filament carbon fibers were formed in the reaction of methanol and ethanol compared to methane fuels. A possible reason for this was that the hydroxide group (i.e., OH) present in the molecular structure of the alcohols removed some of the carbon from the electrode surface.¹³⁶

Modafferi et al. synthesized an Ni–Ru–GDC anode via the hydrothermal method and studied its steam reforming (SR) and auto-thermal reforming (ATR) of propane. At reaction temperatures above 700 °C, high propane conversion and syngas (H₂ + CO) productivity were obtained. However, under SR conditions, the catalyst was inactivated due to the large deposition of filamentous carbon and amorphous carbon. In contrast, under ATR conditions, the formation of coke was completely inhibited because the presence of oxygen promoted the vaporization of the carbon residues.¹³⁷ In addition, Boaro et al. drew the same conclusion for an Ni–Rh–GDC anode.¹³⁸ Alternatively, Hibino et al. studied the use of Ni–Ru–GDC catalysts for SOFC anodes also, and when methane, ethane, and propane were used as the fuel, the cells with the Ni–Ru–GDC anode had MPDs of 750, 716, and 648 mW cm⁻² at 600 °C, respectively. The high electrochemical performance was due to the fact that Ru could efficiently promote the reforming reactions of H₂O and CO₂ produced by the electrochemically oxidized hydrocarbon fuels. In addition, in the stability test of the constant current of 600 mA cm⁻², the voltage of the cell did not change for almost 20 h.¹³⁹

Hussain et al. studied the performance of the Ni–Pt–GDC–Sr_0.94Ti_0.9Nb_0.1O₃ (STN) composite anode at a low temperature (400–600 °C), and the SOFC showed a 0.10 Ωcm² polarization resistance at 600 °C. The HR-TEM analysis displayed that nanocrystalline alloys of Ni–Pt were formed in the Ni–Pt–GDC electrocatalyst together with the coexistence of Ni and Pt nanoparticles as separate phases.¹⁴⁰

Pd, Pt, Rh and Ru are actually classified as Pt group metals and have superior reforming catalytic properties compared to Ni, Fe and Co; therefore, Pt group metals were used as electrodes in the early SOFC studies. However, due to the high cost of noble metals, there are few reports on the use of noble metals in bimetallic-anode research, which also limits the use of commercial SOFCs in the future.

The acidic sites on the anode are known to facilitate hydrocarbon cracking reactions. Thus, it is expected that the formation of carbon can be inhibited by adding alkali metals or alkaline earth metal oxides (such as Li, Na, CaO, BaO, and MgO) to the Ni-based cermet anode to improve its alkalinity. Moreover, the excellent hydrophilicity and adsorption of CO₂ by alkali metals or alkaline earth metal oxides can accelerate the removal rate of carbon depositions through in situ steam/CO₂ reforming hydrocarbon fuels and the adsorption hydroxyl groups (*OH) oxidized with adjacent carbon.

Liu et al. developed a simple chemical Li/Na insertion method to prepare an Li/Na-modified Ni–SDC anode (Li–Ni–SDC, Na–Ni–SDC), and found that the MPDs of the cells with the Li- and Na-modified Ni–SDC anodes reached about 212 and 231 mW cm⁻² at 800 °C for CH₄ fuel, respectively, which are much higher than that of the original Ni–SDC anode SOFC (105 mW cm⁻²). Moreover, they showed better long-term stability in 70 h, which was attributed to the unique ability of the alkaline metals to absorb H₂O and CO₂ under the SOFC operating conditions.¹⁴¹

Takeguchi et al. found that the addition of CaO with strong basicity to Ni–YSZ cermet modified the metal Ni to a slightly cationic state, which was effective in inhibiting carbon deposition without reducing the reforming activity.²⁴ Asamoto et al. also drew a comparable conclusion, namely, that the addition of CaO to the Ni–SDC anode slightly reduced the MPD; however, it enhanced the stability of the SOFC performance and inhibited carbon deposition on the anode. In methane fuel at 700 °C, the cell with the Ni–SDC anode exhibited an MPD of 40.3 mW cm⁻², while that with the Ni–CaO–SDC anode had an MPD of about 35.6 mW cm⁻². The reduced electrochemical performance of the SOFC was attributed to the dissolution of CaO in the ZrO₂ or CeO₂ lattice. However, after 30 min, there was 36.1 mg of carbon deposited on the Ni–SDC anode, compared to only 31.0 mg on Ni–CaO–SDC.¹⁴²

La Rosa et al. investigated the effect of BaO doping in the Ni_0.53Cu_0.47–GDC anode, and for the cells doped or not doped with BaO, similar MPDs of 284 and 310 mW cm⁻² were obtained in dry methane fuel at 750 °C, respectively. Alternatively, the study using XRD to characterize the cells after running in dry methane for 200 h showed that the BaO-doped cells did not have significant carbon deposition, while the undoped anodes displayed the presence of carbon deposits.¹⁴³ Islam et al. conducted a more in-depth study by incorporating BaO in the SOFC Ni–YSZ anode via two different methods, i.e., impregnation and microwave irradiation. When the impregnation method was employed, BaO was stable on Ni, but diffused to YSZ and interacted with it to block O²⁻ transport, resulting in morphological changes and volume expansion of YSZ. However, the selectively deposited BaO only on Ni by using microwave irradiation minimized the interaction between BaO and YSZ. Furthermore, the microwave-prepared anodes possessed comparable electrochemical performance with the impregnated anodes in CH₄, but showed lower carbon accumulation. The increase in carbon tolerance was attributed to the discrete BaO/Ni interface, which preferentially adsorbed H₂O, and subsequently gasified the adjacent carbon deposits.¹⁴⁴ McIntyre et al. used in situ vibration Raman spectroscopy in their study and found that the infiltration of 1% BaO significantly reduced the carbon accumulated on the Ni–YSZ cermet anode in CH₄ fuel at 730 °C. This can be attributed to the following reasons: (1) BaO occupied the catalytic site of Ni, slowing down the surface reaction at the anode and (2) the presence of BaO changed the binding energy of carbon and the Ni catalysts. They further studied the ability of different gas-phase reformers to remove carbon, and found that H₂O was the most effective reformer for carbon removal, followed by O₂, and then CO₂. However, H₂O and CO₂ only partially oxidized the anode, while prolonged exposure to O₂ completely oxidized nickel to nickel oxide.¹⁴⁵ Itagaki et al. noticed that the addition of 0.2 wt% BaO to Ni-SDC anodes significantly reduced the carbon accumulated in CH₄ fuel. This is because the addition of BaO further improved the dispersion of Ni particles, which increased the electrochemical reaction field of CH₄.¹⁴⁶

Hu believed that NiO in an NiO–MgO solid solution was more difficult to be reduced than pure NiO, which is mainly because MgO isolates NiO and inhibits the formation of Ni–Ni bonds during reduction, which contributes to the formation of very small Ni particles to inhibit carbon deposition. Also, due to the surface alkalinity of the NiO–MgO solid solution, it demonstrated excellent activity and selectivity for CO₂ reforming of methane.¹⁴⁷ Zanganeh et al. also considered that for Ni_xMg_1−xO solid-solution catalysts, the high dispersion of reduced nickel species, the alkalinity of the carrier surface, and the nickel-support interaction endowed the catalyst with very high resistance to carbon formation.¹⁴⁸

Yang et al. obtained an excellent MPD of 714 mW cm⁻² in methane fuel at 800 °C by impregnating 2.5 wt% MgO in an Ni–SDC porous anode, and the SOFC could run for 330 h with a power attenuation about 12.5%. Furthermore, through material characterization and theoretical calculations, they believed that the addition of a small amount of MgO to nickel cermet will increase the Lewis alkalinity of the anode, thereby enhancing the ability of the catalyst to chemically adsorb H₂O and CO₂ generated from SOFC operations, and hence reduce coking. Also, HAADF-STEM showed that small Ni particles were located on the surface of the MgO nanoplates.¹⁴⁹ Xie et al. prepared a nickel-based anode modified with MgO nanolayers via the in situ reduction of an Ni_0.9−xCu_0.1Mg_xO solid solution, which could greatly eliminate sample inhomogeneity and additional costs due to the manufacturing methods. The MPD of the SOFC using the Ni–Cu–0.025MgO–SDC anode reached 670 mW cm⁻² in methane fuel at 700 °C, which is about 10% higher than that using the Ni–Cu–SDC anode. Combined with synchrotron vacuum ultraviolet photoionization mass spectra, they deduced that MgO improved the selectivity for C₂₊ hydrocarbons and directly inhibited the cracking of CH₄, and due to the high Lewis alkalinity of MgO, steam adsorption and dissociation were promoted, thereby producing more CO by indirect means.⁷⁸

2.1.2 Non-nickel-based bimetallic cermet. Similar to nickel, Co is also known as a dehydrogenation catalyst. Anodes based on Co–Cu bimetallic mixtures exhibited excellent anti-carbon formation stability when directly utilizing methane and had improved performance compared to pure copper SOFCs. Lee et al. synthesized a Co–Cu–CeO₂–YSZ anode using the impregnation method, which exhibited an MPD of 250 mW cm⁻² in CH₄ fuel at 800 °C and could operate stably for 500 h under a humid CH₄ atmosphere.¹⁵⁰ However, due to the sintering of the Cu phase, there was a discontinuity in the Co conductive network of the anode. Gross et al. improved the conductive network by electrodepositing Co onto reduced Cu to fabricate Co–Cu electrodes. Given that Cu easily diffused through the Co membrane and formed a monolayer of Cu on Co, this anode had exceptional resistance to carbon formation in hydrocarbon fuels.¹⁵¹ Fuerte et al. synthesized Co–Cu–CeO₂ anodes via the inverse microemulsion method and found that replacing copper with 20 at% cobalt could achieve higher stability to carbon formation. Moreover, the bimetallic Co–Cu anode showed high catalytic activity for methane oxidation, which was conducive to charge transfer at the TPB.¹⁵²

Yan et al. fabricated a Co–W–YSZ anode by infiltrating cobalt nitrate and ammonium metatungstate in YSZ, and found that the formed Co₃W and Co₇W₆ intermetallic nanoparticles had high thermal stability up to 900 °C and showed good anti-coking in methane. Exposure of the Co–YSZ anode to CH₄ at 800 °C for 5 h resulted in carbon deposition of 7 wt% on the SOFC, compared to less than 1.5 wt% on the Co–W–YSZ anode. The carbon removal mechanism of this anode was manifested by the formation of tungsten acid on the surface of the Co–W alloy during SOFC operation, which inhibited carbon deposition in methane.¹⁵³

Kaur et al. explored the use of Cu–Fe alloy. They prepared Cu–Fe–CeO₂–YSZ anodes with different Cu/Fe molar ratios (1 : 0, 3 : 1, and 1 : 1) by dipping and used them for SOFC research in n-C₄H₁₀ fuel. The XRD results showed the formation of cubic phases of Cu–Fe metals. With an increase in the Fe loading, an improvement the performance was observed, where the single cell with a Cu : Fe molar ratio of 1 : 1 reached an MPD of 240 mW cm⁻² at 800 °C in n-C₄H₁₀. They found that the output power dropped only slightly over the operation for 50 h, and TGA and SEM showed that the carbon formed at the anode was non-graphite.¹⁵⁴ Further, they used the Cu–Fe–CeO₂–YSZ anode for CH₄ fuel research, and the MPDs of the cell with the Cu–CeO₂–YSZ and Cu–Fe–CeO₂–YSZ anodes were 70 and 90 mW cm⁻² in CH₄ fuel at 800 °C, respectively, and the MPD increased with an increase in the Cu–Fe metal loading. During a 46 h stability test of the Cu–Fe–CeO₂–YSZ anode, the performance only dropped from 125 mW cm⁻² to 100 mW cm⁻², which showed the good stability of the SOFC in methane.¹⁵⁵Table 1 presents a summary some typical bimetallic-cermet anode materials and their electrochemical performances.

Table 1 Summary of electrochemical performance of SOFCs with typical bimetallic-cermet anode materials

No.	Anode	Temperature (°C)	Fuel	MPD (mW cm^−2)	Long term (h)	Ref.
1	Ni–Cu–CeO2–YSZ	800	CH4	330	500+	94
2	Ni–Cu–YSZ	900	City coal gas	43	—	95
3	Ni–Cu–ZDC	700	CH4	87	48+	96
4	Ni–Cu–CeO2–YSZ	700	CH4	120	5+	97
5	Ni–Cu–BZCYYb	750	CH4	106	130+	98
6	Ni–Cu–PBMO	700	C3H8	200	—	99
7	Ni–Cu–YSZ	800	CH4	49	25+	100
8	Ni–Cu–YSZ	700	CH4	240	200+	101
9	Ni–Cu–YSZ–GDC	850	CH4	436	—	102
10	Ni–Co–GDC	800	CH4	250	—	104
11	Ni–Co–YSZ	750	20% v/v CH4 in He	136	60+	105
12	Ni–Fe–GDC	650	CH4	340	50+	107
13	Ni–Fe–PBMO	700	C3H8	300	—	99
14	Ni–Mo–YSZ	800	CH4–50 ppm H2S	594	120+	111
15	Ni–Mo–CZ	800	Isooctane	212	30+	114
16	Ni–Sn–YSZ	800	CH4	390	49+	120
17	Ni–Sn–YSZ functional layer	650	CH4	410	137+	121
18	Ni–Sn–YSZ	750	Biogas	272	50+	122
19	Ni–Sn–GDC	650	CH4	470	200+	125
20	Sn doped Ni/GDC–GDC	650	CH4	930	40+	126
21	Ni–Sn–SDC	700	CH4	280	230+	127
22	Ni–Sn–SDC	700	CH4	600	72+	128
23	Ni–Ag–YSZ	750	CH4	407	100+	133
24	Ni–Ag–YSZ	750	CH4	251	24+	134
25	Ni–Ag–YSZ	750	C2H6	379	24+	134
26	Ni–Pd–SDC	550	CH4	644	—	135
27	Ni–Ru–GDC	600	CH4	750	20+	139
28	Ni–Ru–GDC	600	C2H6	716	20+	139
29	Ni–Ru–GDC	600	C3H8	648	20+	139
30	Ni–Li–SDC	800	CH4	212	70+	141
31	Ni–Na–SDC	800	CH4	231	70+	141
32	Ni–CaO–SDC	700	CH4	35.6	0.5+	142
33	Ni–Cu–BaO–SDC	750	CH4	284	200+	143
34	Ni–MgO–SDC	800	CH4	714	330+	149
35	Ni–Cu–MgO–SDC	700	CH4	670	—	78
36	Co–Cu–CeO2–YSZ	800	CH4	250	500+	150
37	Cu–Fe–CeO2–YSZ	800	n-C4H10	240	50+	154
38	Cu–Fe–CeO2–YSZ	800	CH4	90	46+	155

---

Considering the efficient catalytic ability of Ni on hydrocarbon fuels, the research on carbon-resistant alloys mainly focuses on Ni-based bimetallic cermet at present. For the second metal species in nickel-based alloys, this review mainly summarizes Cu, Co, Fe, Mo, Sn, Ag, Pd, and Ru metals, as well as alkaline earth metal oxides such as CaO, BaO, and MgO. It is worth noting that in the Ellingham diagram drawn according to the change in enthalpy of metal oxide formation at different temperatures, the metal alloys that can be reduced by hydrogen at the SOFC operating temperature are limited, but there are more types of metals that can be reduced by CO than hydrogen.¹⁵⁶ Thus, future alloying studies should consider the synthesis of a wider variety of bimetallic cermets through non-hydrogen reduction pathways.

For the bimetallic-cermet preparation process, the producing method can not only affect the mass of metals and the fabrication time, but also determine the most important property, which is the cermet structure. In addition to the traditional alloy preparation methods, such as ball milling with firing and impregnation method, plating, electrodeposition, microwave irradiation process and co-tape casting can also be considered. In addition to alkane fuel, alcohol fuel should also be studied because the hydroxyl group in the molecular structure of alcohol fuels is beneficial to resist carbon deposition. However, compared with alkane fuels, the molecules in alcohol fuels require more oxygen ions to react, and thus higher requirements are demanded on the charge transfer ability of the anode materials.

For existing hydrogen reduction systems, the alloys of Mo, Sn, and Ru with Ni and the solid solution of MgO with Ni demonstrate better electrochemical performances, and thus these systems should be given more attention. Moreover, researchers have also established a relatively mature theoretical system for the excellent stability and carbon-resistant mechanism of the above-mentioned systems in hydrocarbon fuels. Briefly, the excellent stability and carbon-resistant mechanism of bimetallic cermet anode can be summarized as follows: (I) hydrophilic metal M and H₂O generated by the electrochemical oxidation of fuel are conducive to methane steam reforming, and upon further hydration to hydroxylation reaction, the hydroxyl group (*OH) helps the oxidation of adjacent carbon, and is removed in the form of CO or CO₂. (II) The outstanding adsorption of metal M to CO₂ accelerates the in situ dry methane reforming and the combination of adsorbed CO₂ with deposited carbon to form gas-phase CO. (III) The surface of the metal alloy preferentially oxidizes the C atom, promoting the formation of C–O bonds instead of forming C–C bonds. (IV) Carbon deposited on alloy-modified catalysts is more likely to form disordered amorphous carbon. (V) Ni-based alloys prevent an increase in Ni particles, thereby increasing the active sites and expanding the reaction TPB region. (VI) The external electron configuration of metal M is similar to C: [He]2s²2p² (two valence electrons close to stable s-orbitals), and hence the presence of Ni-based alloys hinders the formation of nickel carbide.

Although these bimetallic-cermet anodes show good potential for hydrocarbon catalytic performance, during the actual application environment of SOFCs, they will undergo multiple redox cycles and temperature alteration processes, which will reduce the mechanical strength of the bimetallic-cermet and even cause cracks in the cells. Alternatively, metals or alloys formed by exsolving from perovskite materials of the anode can improve the stability of the cell, which will be discussed in Section 2.2.

2.2 Ceramic materials

Ceramic oxide anodes, such as SrTiO₃, La_4.0Sr_8.0Ti_11.0Mn_0.5Ga_0.5O_37.5, La_0.75Sr_0.25Cr_0.5Mn_0.5O₃ and Sr₂MgMoO_6−δ mentioned above, are known for their slow carbon deposition, high temperature and redox stability. Unlike metals, they are not catalysts for C–H bond breaking, but contribute to the direct electrochemical oxidation of hydrocarbon fuels. However, the main disadvantages of oxide anodes are their low catalytic activity and electronic conductivity. Typically, they are doped or used for in situ exsolution of cations to form highly reactive metals or alloy nanoparticles to achieve the desired properties.

2.2.1 Cerium-based materials. Among the rare earth family members, cerium (Ce) is the most abundant element with a content of 66.5 ppm in the Earth's crust. CeO₂ has attracted significant interest given that it plays a vital role in emerging technologies for environmental and energy-related applications.¹⁵⁷ The excellent catalytic activity of CeO₂-based materials stems from the charge migration associated with reversible CeO₂–Ce₂O₃ transitions and oxygen vacancy formation. The process of storing and releasing oxygen from cerium-based oxides is maintained by the Ce⁴⁺/Ce³⁺ redox couple. The addition of other metal oxides to the cubic fluorite phase CeO₂ makes it easier to reduce Ce⁴⁺; thus, effectively increasing the concentration of Ce³⁺. Also, due to the high carbon deposition resistance, dry hydrocarbon fuels can be supply directly to the anode.^158,159

Yun et al. studied the Ni–Sm_0.2Ce_0.8O_2−δ (SDC) anode as an alternative anode in methane fuel and compared it to the traditional Ni–YSZ anode. The mixed ion and electron conductive properties of SDC under reduction conditions enlarged the reaction TPB, resulting in a 20–25% improvement in the performance of the SOFC with the Ni–SDC anode compared to that with the Ni–YSZ anode. In addition, the performance of the SDC-coated Ni–SDC anode under 0.1 A cm⁻² at 850 °C for 180 h did not show any significant degradation, indicating that SDC caused the oxidation of methane and inhibited carbon deposition by electrochemically oxidizing the carbon produced from methane cracking on Ni. In contrast, Ni–SDC had lower performance when switching from H₂ to CH₄. This study, based on the most widely used SOFC Ni–YSZ anode demonstrated the advantages of CeO₂-based materials in hydrocarbon fuels.¹⁶⁰

Putna et al. investigated Rh supported on SDC as an anode, demonstrating that it was possible to oxidize CH₄ electrochemically in an SOFC without prior steam reforming, but the power density of the cell was very poor (less than 7 mW cm⁻² at 800 °C).¹⁶¹ Given that Cu–YSZ anodes, unless cerium dioxide is added, perform poorly in methane,¹⁶² Lu et al. did similar work to study the effects of adding CeO₂ by testing Cu–SDC and Au–SDC anodes in n-butane fuels. The MPDs for these two cells made were 122 and 132 mW cm⁻² at 650 °C with H₂ fuel, respectively, and both around 50 mW cm⁻² at 650 °C with n-butane fuel. These similar performances showed that the two metals were only electronic conductors and had no catalytic effect in the direct-oxidation. The addition of CeO₂ suggested that it played an important role in improving the anode performance by increasing its catalytic activity or mixed ion conductivity.¹⁶³

Gd-doped ceria (GDC) has also been studied by many researchers because of its excellent performance in various SOFC fuels. Marina et al. studied a Gd-doped ceria (GDC) anode with a 10–15 μm thick Au–GDC composite layer as the current-collector. It was found that the MPD of the cell was 178.5 mW cm⁻² at 1000 °C with 4.9% CH₄ + 2% H₂ + 3% H₂O + 90.1% N₂ fuel. Moreover, after 1000 h of operation at 1000 °C with a low partial pressure of CH₄ and H₂O of 10 and 3 kPa, respectively, no carbon depositions were found on the GDC anode. The cell could last several rapid thermal cycles in the temperature range of 200–1000 °C and underwent a complete redox cycle without degradation.¹⁶⁴ Wisniewski et al. applied Ir-loaded-Ce_0.9Gd_0.1O_2−x (Ir–GDC) in CH₄/CO₂ ratios between 2 and 0.66 at 600–800 °C. They found that the Ir–GDC catalysts showed considerable stability for DRM, and no carbon formation was observed except under highly reductive conditions of CH₄:CO₂ = 2 at 800 °C.¹⁶⁵ Han et al. studied the performance of Ni-GDC anodes in H₂, CH₄, C₃H₈, NH₃ or simulated underground coal gasification (UCG) gases. The composition of the simulated UCG gas at room temperature was 66% H₂ + 19% CO + 15% CH₄. Among them, the MPD of the cells in different fuel gases was reduced in the order of H₂ > simulated UCG gas > NH₃ > CH₄ > C₃H₈. The Ni-GDC anodes showed good stability in other fuel gases except for C₃H₈ after 18 h at the constant voltage of 0.6 V at 600 °C. Among them, in CH₄ and C₃H₈ fuels, the coking results of the cells and the long-term stability were not as bad as predicted in the function of balance composition (moles%) with the fuel and the temperature between 100–900 °C. These results were attributed to two possible causes, where one was that the slow kinetics at 600 °C led to less pyrolysis of CH₄ and C₃H₈, resulting in less coke formation. The other was due to the rapid migration of oxide ions. The relatively abundant supply of O²⁻ reacted with the deposited solid carbon to help suppress carbon deposition at low operating temperatures.¹⁶⁶

Using an Ni–YSZ/yttria-doped ceria (YDC) anode, Murray et al. pioneered the work of direct electrochemical oxidation of methane at low operating temperature (generated power densities of up to 370 mW cm⁻² at 650 °C). The amount of carbon deposited increased with an increase in temperature when the temperature was higher than 700 °C, but at a given temperature, less carbon deposition occurred on Ni–YSZ–YDC than on the Ni–YSZ anode.¹⁶

Ahn et al. examined the effect on the SOFC anode performance of replacing CeO₂ with a Ce_0.6Zr_0.4O₂ (ZDC) solid solution in Cu–YSZ composites. Before Cu was added, CeO₂ was calcined at 723 and 1273 K, respectively, resulting in a significant reduction in the anode performance with the MPD decreasing from 280 to 170 mW cm⁻² in humidified (3% H₂O) H₂ fuel at 973 K. Conversely, the MPD of the cell decreased from 240 to 210 mW cm⁻² when CeO₂ was replaced with Ce_0.6Zr_0.4O₂. Compared with CeO₂, Ce_0.6Zr_0.4O₂ improved the thermal stability of the anode, which was ascribed to the improved reducibility of the Ce_0.6Zr_0.4O₂ solid solution. Although Cu–ZDC–YSZ has not been tested in hydrocarbon fuels, we expect that it will also demonstrate good thermal stability.¹⁶⁷ Similar work was performed by Song et al. using 10% ZDC, which verifies our hypothesis. Compared with pure CeO₂-based anodes, the Zr_0.1Ce_0.9O_2−δ anode exhibited good thermal stability and activity in 97% CH₄ + 3% H₂O fuel at 700 °C. The MPD of the cell with the Cu–ZDC anode was 32 mW cm⁻¹, while the cell with the Cu–CeO₂ anode showed a peak power density of only 28 mW cm⁻². Short-term stability tests for all the cells at 700 °C were conducted at a constant voltage of 0.6 V under 97% CH₄/3% H₂O. After the stability test, the carbon content on the anode surfaces of Cu–ZDC was 1.93 wt%, while that on Cu–CeO₂ reached 2.19 wt%⁹⁶.

A co-doped CeO₂-based anode with Y and Yb was first reported by Zhou et al. for use in direct methane SOFCs. This novel ionic conductor, Ce_0.8Y_0.1Yb_0.1O_1.9, exhibited higher ionic conductivity between 400–750 °C than single-doped CeO₂ (Ce_0.8Y_0.2O_1.9 and Ce_0.8Yb_0.2O_1.9). The two rare earth elements worked synergistically to improve the ionic conductivity and reforming or oxidation of hydrocarbon fuels. The MPD of the Ni–CYYb configuration cell was 94 mW cm⁻² at 800 °C in dry CH₄ fuel. After the test at constant current density of 200 mA cm⁻² at 750 °C for 120 h, SEM–EDS characterization showed that there was essentially no carbon deposition on the anode. In addition, the stability of the Ni–CYYb SOFC anode in H₂ contaminated with 5 ppm H₂S was also studied, and they found that it could catalyze the conversion of H₂S to SO₂. All these results indicate the outstanding potential of this new material as an SOFC anode.¹⁶⁸Table 2 presents a summary of some typical cerium-based anode materials and their electrochemical performances.

Table 2 Summary of electrochemical performance of SOFCs with different cerium-based anode materials

No.	Anode	Temperature (°C)	Fuel	MPD (mW cm^−2)	Long term (h)	Ref.
1	SDC-coated Ni–SDC	850	CH4	248	180+	160
2	Rh–SDC	800	CH4	7	—	161
3	Au–SDC	650	n-butane	50	—	163
4	Cu–SDC	650	n-butane	50	—	163
5	Au–GDC	1000	4.9% CH4 + 2% H2 + 3% H2O + 90.1% N2	179	1000+	164
6	Ni–GDC	600	66% H2 + 19% CO+ 15% CH4	537	48+	166
7	Ni–YSZ–YDC	650	CH4	370	—	16
8	Cu–ZDC	700	CH4	32	48+	96
9	Ni–CYYb	800	CH4	94	120+	168

---

The process of storing and releasing oxygen from cerium-based oxides is maintained by the Ce⁴⁺/Ce³⁺ redox couple and doping suitable metal ions in the cubic fluorite phase CeO₂ can reduce the concentration of Ce⁴⁺, thereby increasing the oxygen vacancy concentration for excellent catalytic activity. The research on cerium-based oxide doping is currently mature, mainly including samarium-doped ceria (SDC), gadolinium-doped ceria (GDC), yttria–doped ceria (YDC), zirconia–doped ceria (ZDC), and Ce_0.8Y_0.1Yb_0.1O_1.9. For the ceramic composition of the anode, the bi-ceramic composition (e.g., Ni–YSZ–YDC) is beneficial for the expansion of the TPB when the structure of the cell is reasonable. Also, in addition to the common cerium-based materials, the above-mentioned BaZr_0.1Ce_0.7Y_0.1Yb_0.1O_3−δ, Pr_0.5Ba_0.5MnO_3−δ and in situ exsolution of perovskite materials are also popular research areas for ceramic materials and may have better properties than cerium-based oxides.

2.2.2 In situ exsolution of perovskite materials. Metal oxides that conform to the chemical formula of ABO₃ and cubic structure are commonly referred to as perovskite oxides. In the conventional perovskite lattice, the smaller B-site ions have a coordination number of 6 and occupy the corner of the cube to form the BO₆ octahedron, and the larger A-site ions have a coordination number of 12 and are arranged in the center of the cube, as shown in Fig. 6(a). Perovskite is a substance with great potential, in which the A and B sites can be doped with many different elements, causing changes in its ion and electron conductance, and thus the A–A′ and B–B′ sites are endowed with catalytic properties. Thus, the ions doped with extra valences into the A and B sites may cause the material to produce extra electrons, making it an n-type semiconductor, and the ion doping with a lower state will produce oxygen vacancies and cause a change in the O²⁻ conductivity.

Fig. 6 (a) Unit cells of perovskite-type oxides and typical cations occupying the A and B sites.¹⁸⁶ (Reprinted from ref. 186 with permission from Multidisciplinary Digital Publishing Institute.) (b) Comparison of heterogeneous electrodes manufactured by infiltration and exudation. (c) Exsolution of Ni at B-site on the surface of perovskite. Reproduced with permission.¹⁸⁷ (Reprinted from ref. 187 with permission from the American Chemical Society.) (d) Average CH₄ conversion rate of Ni/LSM n = 1 (prepared by exsolution), Ni/LSMO n = 1 (prepared by impregnation) and Ni/YSZ at an extended reaction time (reduction preprocessing condition: in diluted H₂ for 4 h, T = 850 °C. (Reprinted from ref. 174 with permission from Elsevier Ltd.) (e) Schematic diagram of segregation at the A-site and its driving force in LaMnO₃ perovskite system.¹⁸³ (Reprinted from ref. 183 with permission from the American Chemical Society.).

Different cations and defects will make perovskites exhibit a variety of properties due to their structure characteristics.¹⁶⁹ The atoms on the surface of perovskites usually have lower coordination numbers than their inner atoms, and thus the perovskites have higher surface-free energies, which enable the rearrangement of their surface atoms and the separation of certain cations from the matrix to the surface to generate metal nanoparticles. Accordingly, the composition, microstructure of perovskite oxide, and the related properties of reforming, oxidation, and electrocatalysis of fuel will also be changed. When two different materials come into contact, a heterogeneous interface will be formed, and thus the exsolved metal nanoparticles and the perovskite body will become a metal-oxide heterogeneous structure.¹⁷⁰ Heterostructures usually have a small band gap and their internal electric fields accelerate the ion diffusion dynamics, resulting in excellent electrical conductivity. In addition, compared with the weak interaction between the parent and nanoparticles produced by impregnation, the robust interaction between the parent phase and the exsolution particles enhances the structural stability of fuel cells under the working conditions,¹⁷¹ as shown in Fig. 6(b). Therefore, the construction of heterogeneous interfaces by in situ exsolution can optimize the catalytic performance of the materials. Over the years, in situ exsolution of perovskite has become an effective method to prepare highly active materials.

Usually, catalytic nanoparticles dispersed on the outer surface of porous materials are prepared via the physical vapor deposition method or chemical impregnation method and attached to the surface of the carrier material by a high-temperature calcination process. However, the impregnation method restricts the size, uniform dispersion, and firmness of the deposited species;¹⁷² in addition, they tend to aggregate at high temperatures because of the weak mutual effect between the nanoparticles and bulk phase, and their performance will inevitably degrade with the extension of the working time.¹⁷³Fig. 6(d) shows the in situ exsolution material in La_1.5Sr_1.5Mn_1.5Ni_0.5O_7±δ (Ni/LSM n = 1), the impregnated material Ni/La_0.5Sr_1.5MnO_4±δ (Ni/LSMO) n = 1) and the average conversion of CH₄ at 850 °C for the traditional anode Ni/YSZ constructed by Sebastián et al.¹⁷⁴ The conversion rate of Ni after in situ exsolution was 9.71 ± 0.50 mol%, which was higher than that of the impregnated material (7.89 ± 0.20 mol%), and the hydrogen yield was also higher. After 30 h, the CH₄ conversion rate of the impregnated material decreased significantly, which for Ni/YSZ was almost zero. However, the conversion rate of Ni/LSM did not change much in the exsolution system, and the researchers believed that this is due to the sintering or condensation of the impregnated particles, resulting in the uneven distribution of nickel on the material surface, reducing catalyst activity, and forming carbon deposits on the material surface. Recently, researchers have found a more advanced and effective way to dissolve nanoparticles, i.e., in situ exsolution. Compared with the impregnation and deposition methods, in situ exsolution is cheaper and is a simpler process; moreover, the nanoparticles constructed in this way do not easily agglomerate at high temperatures and have reversibility to avoid catalyst agglomeration through reoxidation because they are more evenly distributed and smaller. On account of the strong mutual action between the nanoparticles and matrix, the problem of particle coarsening is effectively alleviated, the life of the catalyst will be greatly improved,¹⁷² and the stability of the cell will also be better.

In general, the segregation of ABO₃ perovskite oxide cations can be divided into A-site cations (such as Sr) and B-site cations, i.e., exsolution elements. A-site cation segregation is a common phenomenon in perovskite oxidation cathode,¹⁷⁵ and Sr is commonly used as an A-site dopant, and thus it is usually called “Sr segregation”. B-site cationic segregation refers to the doping of transition metals or noble metals at the B-site, and the doped metals are subsequently induced to exsolve out of the lattice in situ as active nanoparticles in a reducing atmosphere.¹⁷⁶ Cathode materials usually work in a strong oxidizing atmosphere at high temperatures, which promotes the separation of the A-site cations from the surface, while the B-site is more prone to exsolve in a reducing atmosphere. For instance, if the B sites of perovskite oxides are doped with Pd,¹⁷⁷ Ru,¹⁷⁸ Pt,¹⁷⁹ Co,¹⁸⁰ Ni,¹⁸¹ and other catalytic active noble metals or transition metals, they will exsolve from the perovskite skeleton structure to form nanoparticles in a high-temperature reduction atmosphere, embedding on the perovskite phase and forming a heterogeneous structure.¹⁷⁵

Niania et al. proposed the mechanism of A-site cation (Sr) surface segregation in the La_0.6Sr_0.4Co_0.2Fe_0.8O_3−δ perovskite material based on in situ surface characterization results, where Sr segregates on the exposed grain surface and forms a monolayer, and more strontium will form Sr-based particles over time. Strontium nucleates at the grain boundaries with an increasing segregation rate and the resulting defects can also be heterogeneous nucleation sites, and eventually strontium agglomerates due to surface migration.¹⁸² Lee et al. suggested that the foremost driving force of cation segregation in perovskite-based oxides is the elastic and electrostatic interactions between the dopants in the perovskite oxides and the surrounding lattice.¹⁸³ These researchers demonstrated that the purpose of A-site dopant cation segregation was to decrease the elastic energy to the minimum, which is caused by the mismatching between the size of the dopant and the host cation, and decrease the electrostatic energy produced by the interaction between the surface and the dopant charged defects near the surface, as shown in Fig. 6(e). Neagu et al. observed the annealing of La_0.43Ca_0.37Ti_0.94Ni_0.06O₃ at 900 °C under high vacuum and proposed the mechanism of the in situ exsolution of the B-cations (Ni) from the perovskite, as shown in Fig. 6(c). Sr segregation at the A-site is considered to be one of the important reasons for the chemical instability of the perovskite surface and the corresponding performance degradation of the SOC electrode,¹⁸⁴ for example, the SrO_x segregation layer on SrTi_1−xFe_xO₃(STF) could reduce the oxygen reduction kinetics of the electrode, and thus affect the cell performance.¹⁸⁵ In contrast, the exsolution of the B-cations usually improves the electrochemical performance of the electrode.

The classical nucleation theory indicates that the main reason for exsolution and nucleation in the reduction process is the over-saturation of the cations in the main lattice. In the case of perovskites composed of AB_1−xM_xO₃, the metal element M is more easily reduced to metal from cations and exsolved than other transition metals B. As mentioned earlier, the Gibbs free energy variation (ΔG) in the process from Mⁿ⁺ to M is the dominating driving force for the exsolution. Gao et al. proposed that exsolution is the result of four physical processes, i.e., diffusion, reduction, nucleation, and growth, as shown in Fig. 7(a). The Ni²⁺ ions exsolved from the matrix to the surface, then were reduced to zero-valent Ni, and the reduced Ni became nanoparticles, which increased in size with time. After nucleation, the growth rate of the particle size slowed down, indicating that ΔG began to increase after reaching the minimum value.¹⁸⁸Fig. 7(b) shows the ΔG of each metal element oxide reduced to metal in La_0.4Sr_0.4Sc_0.9Ni_0.1O_3−δ. The ΔG > 0 of the La, Sr, and Sc metal oxides show that they are thermodynamically stable in the reduction conditions; nevertheless, the ΔG < 0 of Ni shows that it is easier to be reduced, which is conducive to exsolution.

Fig. 7 (a) Schematic diagram of the exsolution process. (b) Sc, La, Sr, and Ni oxides are reduced in H₂ according to the Gibbs free energy of the corresponding metal. (Reprinted from ref. 188 with permission from Elsevier Ltd.) (c, d), (f, g) and (i, g) are (001)-oriented film, (110)-oriented film, and (111)-oriented LSTN film AFM and SEM top view images, respectively. The scales of (d, g and j) and their inset are 1 μm and 200 nm, respectively. (e, h and k) Surface schematic of an LSTN film exsolved with Ni nanoparticles. AFM imaging area = 1 μm². The thickness of LSTN film is about 300 nm. AFM height and transverse diameter distribution of representative exsolved particles. (l) (111)-oriented films and (m) (001)-oriented films show that Ni particles exsolved from LSTN and embedded in the films, the lower normalized interface area in the (001)-oriented films mean a higher interface energy. (Reprinted from ref. 191 with permission from the American Chemical Society.).

Oh et al. used the COMSOL Multiphysics software to calculate the elastic field and free energy in the exsolution process, indicating that the interaction between the surface free energy and strain energy is the driving force for exsolution.¹⁸⁹ In addition, the surface and interface energy of the nucleation particles and the local strain field in the matrix may further affect the distribution of grain size. In the initial study of the surface effect, the researchers believed that the (110) face was most suitable for B-site cation exsolution.¹⁹⁰ However, Kim et al. came to a different conclusion that other planes such as the (001) and (111) planes also exhibited strong exsolution tendencies.¹⁹¹ La_0.2Sr_0.7Ti_0.9Ni_0.1O_3−δ (LSTN) films in the (001), (110), and (111) orientation were deposited on the corresponding single-crystal SrTiO₃ substrates in the (001), (110), and (111) orientation using pulsed laser deposition (PLD). The exsolution of the nickel nanoparticles on the three oriented LSTN films is shown in Fig. 7(c)–(k). On the film in the (111) orientation, the largest number of particles (195 μm⁻²), the smallest particle size (20 ± 7 nm), the strongest embedding, and the shortest distance between particles were observed, which was the most uniform oriented film. The smaller average particle size indicates higher catalytic activity, and without a strong shielding effect or diffusion effect due to the large distance. Fig. 7(l) and (m) shows that the surface interface energy of the (111) orientation was the lowest, and because the normalized projected area of the partial (001)-oriented film connected to the LSTN matrix is smaller than the corresponding area of the (111)-oriented film, the interface energy (γ_int) of the (001) oriented film was higher. On account of the higher interface energy, the assembly nuclear energy barrier and critical nucleation size increased and the nucleation rate and particle population density decreased.

Early studies focused on perovskite oxides with a stoichiometric ratio of A/B = 1, which exsolved cations through the inherent reduction driving force once used by researchers to prepare “smart catalysts” for controlling automobile emissions.¹⁹² For example, Nishihata et al. found that the Pd-perovskite catalyst had higher activity than the traditional Pd–Al₂O₃ catalyst.¹⁹³ The method for the preparation of this “smart catalyst” has been widely used in the preparation of in situ exsolution perovskite oxide SOFC anodes lately.¹⁹⁴ However, the intrinsic reducibility of perovskite can only drive limited reducible cations to exsolution¹⁷⁶ and the types of metal elements that can be exsolved in situ are limited to noble metals such as Ru, Pt, Pd, and Ni. Fig. 6(d) shows that the average CH₄ conversion of Ni/YSZ was 14.00 ± 0.44 mol%, which was close to the reaction equilibrium value 15 mol%, while the average conversion of the exsolution was only 9.71 ± 0.50 mol%. Obviously, the electrocatalytic activity of the nano-catalyst exsolved by inherent reducibility is insufficient and its electrochemical performance is still limited compared with the traditional nickel-based catalyst. Thus, recently, to improve the catalytic activity of in situ exsolved perovskite, many researchers have focused on investigating A-site defects (i.e., perovskite oxides with stoichiometric ratio of A/B < 1) using cheap and stable transition metals to regulate the exsolution.¹⁹⁵ The A-site defect increases the tendency of the B-site cation to exsolve from the lattice. Also, the exsolution restores the perovskite to a stable structure locally, making the whole structure “defect-free”, thus conforming to the chemometrics of ABO₃. This mechanism can be explained by the point defect reaction proposed by Neagu et al.,¹⁷² as follows:

A large number of A-site defects produced by doping will make formula (14) move to the left, limiting the number of inherent Schottky defects, thus reducing the number of . Simultaneously, formula (15) moves to the right under a certain p(O₂) to counteract the changes in formula (14), which promotes the reduction of lattice oxygen and the doped metal to in the reducing gas. Afterward, the metal atoms spontaneously detach from the oxide lattice, leaving cationic vacancies , as shown in formula (16). Another metal exsolution path is shown in formulas (17) and (18). The doped metal and lattice O generate MO, and then the metal will exsolve in the reduction process. The presence of vacancies destroys the stability of the perovskite lattice, partially causing the B-site cations to spontaneously exsolve to reconstruct the stoichiometry of the bulk phase. In addition, perovskites with A, O-site defects can be considered as B-site excess, which explains the tendency of B-site exsolution from another perspective. At the atomic scale, the process of removing oxygen from perovskites lacking A-site cations can be seen as moving the B-site dopant locally from the main perovskite skeleton to the outside to form an initial exsolution state, as shown in Fig. 8(a).¹⁷² The A-site-occupied unit cell remains unaffected, while the B site in the A-site missing unit cell is excess and will be discharged.

Fig. 8 (a) Schematic diagram of B-site cation exsolution in A-site defect perovskite, where O-site is represented by silver balls; B site is represented by yellow highlighted gray balls; black big spheres represent A site; and red color indicates that V_A is an A-site vacancy. By reducing the oxygen in the A-site defect unit, some B-sites are partially isolated from the perovskite matrix and enter the initial BO_n exsolution, as showed in the red atom group in the right figure. The effect of non-stoichiometry in stoichiometry materials and A-site-deficient perovskite exsolution showed by SEM. (b) Oxygen stoichiometry and A-site-deficient La_0.52Sr_0.28Ni_0.06Ti_0.94O₃ after reduction at 930 °C for 20 h, and Ni was exsolved in 5% H₂/Ar. (c) La_0.3Sr_0.7Ni_0.06Ti_0.94O_3.09 sample with excess stoichiometric O with A site stoichiometric did not exsolve after reduction at 930 °C for 20 h under 5% H₂/Ar conditions. (Reprinted from ref. 172 with permission from Macmillan Publishers Ltd.)

Neagu also studied the exsolution of Ni in two samples containing 6% Ni at the B-site with different oxygen vacancy concentrations at the A-site under the same reduction conditions.¹⁷² The A-site oxygen-deficient material La_0.52Sr_0.28Ni_0.06Ti_0.94O₃ generated a large number of metal Ni nanoparticles, which exsolved to the perovskite surface, as shown in Fig. 8(b). In contrast, no particle growth was observed in La_0.3Sr_0.7Ni_0.06Ti_0.94O_3.09, which is stoichiometric and oxygen-excess, as shown in Fig. 8(c).

This experiment showed that the B-site atoms had a greater tendency to undergo exsolution in A-site-deficient perovskites than that in oxygen-rich perovskites, and A-site defects help to form oxygen vacancies with high fluidity, an increase in which can improve the ionic conductivity.¹⁹⁶ When the composition of the perovskite oxide is changed to A/B < 1, more cations can segregate on the surface of the perovskite bulk, which may be a more efficient method to construct perovskite structures with well-dispersed nanoparticles. Therefore, it is possible to control the surface modification of perovskites by constructing a non-stoichiometric system, which cannot be achieved by traditional-component perovskites.¹⁷² Sun et al. used the in situ exsolution method to prepare nickel-doped and strontium-doped lanthanum chromate (LSC) with A-site defects as SOFC anodes.¹⁷⁶ The electrochemical performance of the perovskite anodes promoted by the A-site defects in the SOFC was evaluated for the first time and it was found that the A-site defects indeed facilitated the in situ exsolution of Ni on the nickel-doped anodes (La_0.7Sr_0.3)CrO₃(LSCNi). The cell with the A-site-defect anode showed an MPD of 460 mW cm⁻² in 5000 ppm H₂S-H₂, while that of the cell using the stoichiometric LSCNi anode was only 135 mW cm⁻². Dueñas et al. compared the stoichiometric perovskite oxide La_0.70Sr_0.3Cr_0.85Ni_0.15O_3−δ(L70SCrN) with 5% A-site deficient La_0.65Sr_0.3Cr_0.85Ni_0.15O_3−δ(L65SCrN) and found that the latter had higher exsolution ability. The researchers calculated that the net oxygen consumption of L70SCrN was 0.27 (mol O mol⁻¹), while the net oxygen consumption of L65SCrN was 0.36 (mol O mol⁻¹).¹⁹⁷ The lack of A site enhanced its reducibility, which is conducive to the exsolution of metal Ni. Zhu et al. prepared Sr_0.95(Ti_0.3Fe_0.63Ni_0.07)O₃(STFN), which is A-site-deficient, by substituting Ni for Fe in the SrTi_0.3Fe_0.7O₃(STF) perovskite anode.¹⁹⁸ The MPD of the cell with the STFN anode that exsolved Ni–Fe nanoparticles was 1300 mW cm⁻² at 850 °C, while the MPD of the cell with the STF anode was 1100 mW cm⁻². In summary, researchers assume that the presence of A-site defects in perovskite means there is the driving force to trigger the external exsolution of the B-site cations, and the produced nanoparticles have a more uniform surface distribution and coverage with better catalytic performance.

In addition, some reports demonstrated that the particle size, coverage or distribution of in situ exsolution nanoparticles on the perovskite surface can be adjusted by changing factors such as the reduction temperature, working time, and number of oxygen vacancies. Kobsiriphat et al. found that the volume of B-site Ru nanoparticles increased with an increase in temperature, which exsolved from La_0.8Sr_0.2Cr_1−xRu_xO_3−δ. After annealing in H₂ at 600 °C, 700 °C and 800 °C for 15 min, the average diameter of the Ru nanoparticles increased from 1 nm at 600 °C (Fig. 9(a)) to 3–4 nm at 800 °C (Fig. 9(c)), and the grain density also decreased with an increase in annealing temperature.¹⁹⁹ He et al. also found that temperature changes can adjust the exsolution of the nanostructures. These researchers doped cerium dioxide in titanate at 1400 °C and found that up to one-third of the doped cerium was exsolved from the A-site of La_0.8Ce_0.1Ni_0.4Ti_0.6O_3−δ at 1300 °C to form a lamellar structure on the grain boundary, as shown in Fig. 9(d). When the temperature was lower than 1300 °C, cerium oxide cubes of different sizes appeared on the perovskite surface. Fig. 9(j) shows that when La_0.8Ce_0.1Ni_0.4Ti_0.6O_3−δ was annealed at 1300–900 °C, the cubic exsolution tendency of CeO₂ increased with temperature from the bottom to top, and the lower temperature could make more CeO₂ cubic phase grow with a smaller size. When annealed at 1200 °C, the particle diameter was about 220 nm, and there was about one particle per μm² on the perovskite surface (Fig. 9(e)). When the annealing temperature was reduced to 1100 °C, a large number of CeO₂ particles was exsolved on the surface of titanate. Compared with 1200 °C, the ceria particles were smaller (130 nm) and their density was higher (about 9 particles per μm²), as shown in Fig. 9(f). Although the exsolution of the ceria cubes could be observed under scanning electron microscopy at 1100 °C, it was possible that XRD could not detect the diffraction peak of CeO₂ because the amount of ceria exsolution at low temperature was lower than the detection limit (<1 wt%). Further lowering the annealing temperature resulted in a smaller cerium dioxide cubic phase, with particle sizes of about 55 and 40 nm at 1000 °C and 900 °C, and densities of about 20 and 10 grains per μm², as shown in Fig. 9(g) and (h), respectively. Fig. 9(i) shows that no CeO₂ cubes exsolved from the surface of the perovskite after annealing at 800 °C, possibly due to the insufficient thermal energy to support cerium cation diffusion at this low temperature. Therefore, the temperature is considered a key factor affecting the exsolution and separation of CeO₂ in this perovskite material.¹⁷³

Fig. 9 (a–c) TEM images of La_0.8Sr_0.2Cr_1−xRu_xO_3−δ annealed in H₂ at 600 °C, 700 °C, and 800 °C for 15 min, respectively. The average diameter of ruthenium nanoparticles increased from 1 nm at 600 °C to 3–4 nm at 800 °C, and the grain density decreased with an increase in the annealing temperature. (Reprinted from ref. 199 with permission from Elsevier Ltd.) (d)–(h) Backscattered electron micrograph of the surface of the La_0.8Ce_0.1Ni_0.4Ti_0.6O_3−δ particle surface annealed in the air for 2 h. (d) 1300 °C, (e) 1200 °C, (f) 1100 °C, (g) 1000 °C, (h) 900 °C and (i) 800 °C. (j) Schematic showing the tendency of CeO₂ cubic exsolution with an increase in temperature (from bottom to top). (Reprinted from ref. 173 with permission from Elsevier Ltd.) (k–m) In situ TEM images of PBMCo in 0.5 Pa H₂, at 770–850 °C, with one image per 40 °C. (n) In situ XRD patterns of PBMCo at room temperature to 800 °C in 10% H₂–N₂. (o) I–V–P curves of the cell with PBMCo anode in the range of 800–900 °C in H₂. (Reprinted from ref. 200 with permission from the American Chemical Society.)

Sun et al. studied the in situ exsolution of Co nanoparticles from a double perovskite-structured crystal, Pr_0.5Ba_0.5MnO_x(PBMCo). Fig. 9(k)–(m) show that there was no evidence of exsolution on the perovskite surface at 770 °C, and Co did not segregate until the temperature increased to 810 °C. When the temperature further increased to 850 °C, large particles appeared on the edge of the perovskite bulk, which adhered to it, and at a certain temperature, these exsolved particles became clearer than before. As shown in Fig. 9(n), the in situ TEM image is consistent with the results of in situ XRD, and a significant Co diffraction peak appeared at 800 °C, indicating that the exsolution of Co mainly occurred above 800 °C, and the exsolution rate was also high. These researchers also characterized the performance of the catalyst as an SOFC anode in H₂, and Fig. 9(o) shows the I–V–P curves at different temperatures. As the test temperature increased, more and more Co nanoparticles exsolved, and their MPD also increased, which was 520 mW cm⁻² at 800 °C and reached 1100 mW cm⁻² at 900 °C.²⁰⁰

The running time also has an influence on the performance of the cell with an exsolution nanoparticle anode. Madsen et al. tested the electrochemical performance of the cell loaded with La_0.8Sr_0.2Cr_0.82Ru_0.18O_3−δ–GDC as the anode at 800 °C and working in wet H₂ over time. The MPD over 15 min was about 250 mW cm⁻² and stabilized at 390 mW cm⁻² over time to 100 h. The results were similar to the tests at other temperatures, where the MPD of the cell increased from 120 mW cm⁻² at 15 min to 300 mW cm⁻² at 100 h at 750 °C.²⁰¹

Neagu et al. introduced the concept of oxygen-deficient δ by reducing perovskite (A_1−αBO₃) with the A-site deficient. δ not only means the number of oxygen vacancies , but also indicates the degree of reduction, which means that the amount of oxygen removed per formula unit of perovskite is quantitatively equivalent to .¹⁷² It is also related to the number of reduced B-site ions , which has a great influence on exsolution. The formation of defects can be expressed by the Kröger-Vink formula, as follows:

Given the effects of δ on solubility and conductivity, Neagu et al. studied the changes in oxygen vacancies (δ) and conductivity (σ) with the quantity of Ga doping (x) after the reduction of La_0.4Sr_0.4Gx_0.4Ga_xO_3−x/2−δ (0 ≤ x ≤ 0.15), and proposed that the oxide ion mobility of the perovskite oxide n-type conductor bulk could be increased by increasing the number of oxygen vacancies, thereby improving the conductivity.¹⁶⁹ It is known that both ion conduction and electron conduction contribute to the conductivity of the ion-electron mixed conductor. When the B-site cation is partially substituted with other low-valent elements (e.g., Ga element), oxide ion vacancies are formed to keep the electrical neutrality of the perovskite structure, as shown in formula (19), while the oxide ion conduction in most solid electrolytes is achieved by jumping from the lattice position to an adjacent vacancy (oxygen vacancy) due to the activation energy.²⁰² Therefore, after replacing Ti⁴⁺ with Ga³⁺, the oxygen generated by compensation doping can support ion conduction together with the vacancies produced in the course of the reduction process. The most likely reason for the occurrence of electron conduction is that electrons jump from B⁽ⁿ⁻¹⁾⁺ to Bⁿ⁺ through oxygen bridges, where the higher the carrier (e⁻) concentration and mobility μ_e⁻, the higher the conductivity σ_e⁻, as follows:

The experimental results showed that both oxygen vacancy δ and conductivity σ increased significantly with an increase in the Ga doping amount. It has been reported that the electronic conductivity of LSGM under reduction conditions is very small, and thus the reduction of Ga has no positive contribution to conductivity. The increase in conductivity is mainly affected by the increase in the concentration of oxygen vacancies, and the change in conductivity is largely determined by the concentration of oxygen vacancies. Fig. 14(b) and (c) show that the effect of the doping number of δ and σ has a similar trend. Under the condition of 1000 °C and 5% H₂/Ar, the δ of the doped sample increased from 0.035 to 0.060. With the Ga doping amount gradually increasing, the conductivity increased initially, and then decreased, but the overall value is larger than that of undoped Ga samples. When the doping amount x = 0.05, the conductivity reached the maximum value of 50 S cm⁻¹. According to formula (19), it can be seen that δ is proportional to the concentration of free electrons, while formula (20) shows that σ is proportional to the concentration of electrons, indicating that σ is proportional to δ, that is, under certain conditions, the conductivity will also increase with an increase in the concentration of oxygen vacancies. The number of oxygen vacancies increases with an increase in the doping amount x, but the conductivity increases first, and then decreases with an increase in x, showing that the increase of oxygen vacancies is not beneficial to redox activity all the time. In addition, for systems with α = 0.2 in the A-site stoichiometry (1 − α), when the value of δ is less than or approximately equal to 0.05, no or only a little exsolution was produced, as shown in Fig. 10(a). Given that δ depends on p(O₂), temperature, and doped elements, adjusting these factors can almost control the B-site cations to exsolve from any A-site defect system after reduction.

Fig. 10 (a) Reduction degree δ (represented by oxygen atoms per formula unit of perovskite, at/f.u.) as a function of the dopant under different situations, indicating in the exsolution region with different temperatures, the degree of exsolution during reduction in a dry or wetted (d or w) reducing gas of 3% H₂O or 5% H₂O/Ar. (Reprinted from ref. 172 with permission from Macmillan Publishers Ltd.) (b) In different cases, the relationship between oxygen vacancy and Ga doping was measured by the increase in weight during oxidation after the first reduction (■), after the second reduction (▲), after annealing at 1300 °C (●). The sample was pellets with 60–65% density. The reduction step was in a flow of 5% H₂/Ar at 1000 °C for 20 h. (c) Electrical conductivity (σ) as a function of Ga content (x) in 60–65% La_0.4Sr_0.4Ga_xTi_1−xO_3−x/2−δ pellets was measured and decreased in 5% H₂/Ar at 1000 °C (20 h). (□) represents the conductivity of the sample after reaching the measurement temperature of 880 °C, and (■) represents the conductivity after 12 h at this temperature. (Reprinted from ref. 169 with permission from the American Chemical Society.) (d) High-resolution XPS spectra of O1s on the anode surfaces of D-LBM, D-LBMFC-1, and D-LBMFC-2, where O1s can be divided into three types: lattice oxygen O_latt (529.1 eV), adsorbed oxygen load O_ad (531.2 eV) and adsorbed water (532.9 eV). (e) I–V–P curves of wet CH₄-loaded single cells with different anodes. (Reprinted from ref. 192 with permission from the American Chemical Society.) (f) Relationship between conductivity and reciprocal temperature of La_0.5Sr_0.5Ti_0.75Ni_0.25O₃: samples prepared in (■) air and in (◆) Ar/H₂ (98/2); after 2 h reduction in Ar/H₂(98/2) at 1200 °C. (△) was reduced at 1200 °C for 2 h, and then oxidized at 800 °C for 5 h in air. (○) was reduced at 800 °C for 48 h, and then tested in wet Ar/H₂(98/2) (pH₂O = 0.025 atm). (Reprinted from ref. 203 with permission from Elsevier Ltd.)

La_0.5Ba_0.5MnO_3−δ with co-doped Co–Fe was synthesized by Hou et al. using the Pechini method, and after reduction, the core–shell nanoparticles consisted of Co_0.94Fe_0.06 alloy exsolved from the surface, which significantly improved the catalytic performance of the anode. These researchers proposed that when the reduction temperature is lower, the oxygen is more unstable, resulting in the stronger redox ability of the sample and easier formation of oxygen vacancies. Surface oxygen is mainly adsorbed on the oxygen vacancies and easily desorbed at high temperatures. The XPS results of O1s (Fig. 10(d) and Table 3) showed that the content of lattice oxygen in the sample decreased, and the content of adsorbed oxygen increased, with the gradually increase in Fe and Co in the sample that had been reduced. The high content of adsorbed oxygen indicates that there is a large number of oxygen vacancies in the material, which could increase the high oxygen ionic conductivity. These researchers also tested the performance of the cells loaded with La_0.5Ba_0.5MnO_3−δ (LBM), La_0.5Ba_0.5Mn_0.9Fe_0.05Co_0.05O_3−δ (LBMFC-1) and La_0.5Ba_0.5Mn_0.8Fe_0.1Co_0.1O_3−δ (LBMFC-2) at 850 °C using wet methane as the fuel. The MPD of the cell with the LBMFC-2 anode was 503 mW cm⁻², as shown in Fig. 10(e). The addition of Co and Fe introduced more oxygen vacancies for perovskite, accelerating the anodic kinetics for electrocatalytic reactions.¹⁹² Arrivé et al. used La_0.5Sr_0.5Ti_0.75Ni_0.25O₃ titanate as the anode of an SOFC, which could make Ni nanoparticles in situ exsolve after reduction. The conductivity of the sample tested under the SOFC anode conditions after reduction at 800 °C for 48 h (symbol ○ in Fig. 10(f)) was nearly an order of magnitude higher than that of the unreduced sample prepared under the same conditions (symbol ◆ in Fig. 10(f)). This is because the pre-reduced sample at 800 °C may have a higher oxygen vacancy concentration than the unreduced sample, resulting in higher ionic conductivity.²⁰³

Table 3 The content of various oxygen species on the surface of the reduced sample based on the deconvolution results of O1s XPS peak¹⁹²

Anode	Lattice oxygen (%)	Adsorbed oxygen (%)	Adsorbed H2O (%)
D-LBM	49.31	42.81	7.88
D-LBMFC-1	47.32	44.46	8.22
D-LBMFC-2	37.01	58.60	4.39

---

For most anodes loaded with Ni nanoparticles, it is still an inevitable problem that the Ni metal particles will catalyze the formation of carbon in a hydrocarbon-containing environment for a long time, but the well-distributed of metal particles on the catalyst can reduce the formation of coking, and therefore the perovskite oxide anode prepared by in situ exsolution is considered to have a better anti-coking performance.⁹³ Oh et al. loaded Ni nanoparticles on La_0.4Sr_0.4TiO₃ by in situ exsolution, nickel nitrate aqueous solution impregnation, and vapor deposition, respectively. Characterization methods such as atomic force microscopy were used to study the exsolution state of the Ni nanoparticles, and the anti-coking performance of the three samples was compared. It was revealed that the in situ exsolution metal particles embedded on the surface of perovskite were the main reason for the in situ exsolved metal particles having a better anti-coking performance.¹⁹⁰Fig. 11(a) shows a TEM image of the exsolution particles at the (110) end. It can be seen that 30% of the Ni nanoparticles was submerged below the perovskite surface, and a clear comparison with the Ni loaded by impregnation and exsolution is shown in Fig. 11(c). The researchers etched the sample with concentrated HNO₃, and it can be observed from Fig. 11(d) and (e) that the pits left were very similar to the density and degree of size distribution of the particles before, confirming that the nanoparticles were embedded in the perovskite surface. Fig. 11(f) shows the AFM image of the etched surface, indicating that they were indeed embedded deep in the perovskite. In contrast, the nanoparticles embedded in perovskites by traditional deposition methods were almost invisible on similar perovskites with A-site defects, indicating that the interaction between the deposited metal and the perovskite phase was much lower.

Fig. 11 (a) TEM image of aged nickel particles dissolved on (110) as-grown surface (dark field) (3% H₂O/5% H₂/Ar, 930 °C, 60 h). (b) TEM image of the metal-perovskite interface (bright field). (c) Schematic diagram of the particle-substrate interface of deposited and dissolved nickel particles. (d) SEM image of particles before etching. (e) SEM image the particles after etching in HNO₃, illustrated by histograms of particle and pore sizes determined by image analysis. (f) Three-dimensional atomic force microscopy image similar to a hole in (e). SEM images of vapor-deposition nickel particles from La_0.4Sr_0.4TiO₃ (g) before aging and (h) after aging (H₂, 650 °C, 24 h and 800 °C, 6 h, respectively). SEM images of Ni particles precipitated from La_0.52Sr_0.28Ni_0.06Ti_0.94O₃ (5% H₂/Ar, 900 °C, 12 h) before aging (i) and after aging (j) (5% H₂/Ar, 900 °C, 70 h). (k) After coking test, Ni particles of about 20 nm were formed on La_0.4Sr_0.4TiO₃, and the growth of carbon fibers was obvious. (l) Growth of considerable carbon fibers from Ni particles (30–100 nm) grown on La_0.4Sr_0.4TiO₃ by vapor deposition. (m) Diagram of the probable carbon fiber growth mechanism based on the literature. (n) After coking test, exsolution of La_0.52Sr_0.28Ni_0.06Ti_0.94O₃ (5% H₂, 880 °C, 6 h) to form Ni particles of about 25 nm and only limited carbon fibers grew. (o) Pseudo-color micrographs depicting the side details of the sample in (n). (p) Side-view micrographs and pseudo-color micrograph detail illustrations of different areas in material (n). (q) After coking, about 60 nm Ni particles were formed on La_0.52Sr_0.28Ni_0.06Ti_0.94O₃ (5% H₂/Ar, 1000 °C, 6 h) by exsolution and only limited coking occurred. (r) Sample and (q) pseudo-color top-view micrograph. (s) Pseudo-color photographs of different parts of g and matching secondary electron micrographs show the empty pits and the particles next to them. The coking test was performed at 800 °C in 20% CH₄/H₂ for 4 h. In the pseudo-color micrograph, green represents perovskite, red represents carbon, and yellow represents nickel metal.¹⁹⁰ (Reprinted from ref. 190 with permission from Macmillan Publishers Ltd.)

Compared with Ni particles just deposited on the perovskite surface, the in situ exsolution Ni particles showed a relatively low tendency of agglomeration and coking. Fig. 11(g)–(j) compare the thermal stability of the deposited particles with the exsolution particles with the same size. The former rapidly agglomerated at 800 °C or lower temperatures, as shown in Fig. 11(g) and (h), while the latter remained stable at 900 °C. The number of Ni particles loaded on the in situ exsolution sample was about twice that of the deposited sample, as shown in Fig. 11(j). To test the anti-coking stability, the samples were placed in 20% CH₄/H₂ at 800 °C for about 4 h. The Ni particles (20 nm) prepared by the impregnation method severely coked and produced carbon fibers, as shown in Fig. 11(k). The Ni particles with a size of 30–100 nm prepared by vapor deposition also produced a large number of micron-sized carbon fibers, as shown in Fig. 11(l). As shown in Fig. 11(n) and (q), the in situ exsolution Ni metal particles of the same size had very little carbon fiber growth because the robust mutual effect between the nanoparticles and the perovskite support prevented carbon fiber growth and particle removal. The growth mechanism of the carbon fibers in the case of in situ exsolution can be explained as the Ni particles adhered to the substrate, and the carbon fibers could only grow on the substrate next to the Ni particles simultaneously, as shown in Fig. 11(m), but the particles occasionally lifted, as showed in region 1 in Fig. 11(p).

Managutti et al. prepared a perovskite in situ exsolution system with Ni-doped (Pr_0.5Ba_0.5)_1−x/2Mn_1−x/2Ni_x/2O_3−δ (S-PBMNx) as an SOFC anode, where x = 0, 0.05, 0.1 and 0.2. The CH₄ conversions of PBMN₀ and PBMN_0.05 (R-PBMN₀ and R-PBMN_0.05) were less than 1% after reduction and in situ exsolution of Ni nanoparticles, while R-PBMN_0.2 with the highest Ni stoichiometry had the best performance in DRM, and the CH₄ and CO₂ conversions were 11% and 32%, respectively. Moreover, R-PBMN_0.2 had almost no carbon deposition (just 0.017 g g_cat⁻¹ h⁻¹) after continuous operation for 5 h at 800 °C. Researchers have also found that bimetallic alloy exsolution anodes of SOFCs may have better performance than single metal exsolution anodes of SOFCs. Iron has been widely studied as a metal promoter because of its ability to effectively enhance the catalytic performance.²⁰⁴ Sun et al. prepared A-site-deficient LaSrCrO₃ (A-LSC) as an anode material doped with Ni, Fe, and Ni-Fe, respectively. By introducing A-site defects under a reducing atmosphere, the in situ exsolution method promoted the uniform distribution of nano-Ni, Fe, and Ni–Fe alloys. The Ni–Fe composites obviously changed the reduction mode of Ni and Fe oxides, and adding Fe changed the reduction mode of the material. Fe and Ni may exsolve on (La_0.7Sr_0.3)(Cr_0.85Ni_0.12Fe_0.03)O_3−x (LSCNi–Fe) firstly, and then form alloys. The presence of the alloy promotes the generation of oxygen vacancies, which is conducive to the exsolution of Ni and Fe. The original reduction behavior of Ni includes two stages, i.e., Ni³⁺ reduction to Ni²⁺, and then Ni²⁺ reduction to metal Ni. The reduction behavior of Fe includes three stages, i.e., Fe³⁺ reduction to Fe²⁺, Fe²⁺ reduction to metallic Fe, and Fe³⁺ reduction to metallic Fe. The performance tests of the cell showed that compared with the cell with the single metal anode catalyst, that with the bimetallic anode catalyst of nano-Ni–Fe alloy with high catalytic activity exhibited a significantly improved electrochemical performance in 5000 ppm H₂S/syngas and the cell was tolerant to carbon deposition, with an MPD of 560 mW cm⁻². This was higher than the 450 mW cm⁻² of (La_0.7Sr_0.3)(Cr_0.85Ni_0.15)O_3−x and 360 mW cm⁻² of (La_0.7Sr_0.3)(Cr_0.85Fe_0.15)O_3−x, which were loaded with Ni and Fe, respectively. Based on DFT calculations, the researchers claimed that Ni–Fe alloys can clear away carbon deposition on the electrocatalyst surface, unlike the individual metal Ni, which promotes the formation of C–C bonds, and this conclusion is consistent with the experimental results, indicating that the addition of Fe significantly promotes the removal of carbon deposition.²⁰⁵

Ma et al. studied Ni-doped La_0.6Sr_0.4FeO_3−δ (LSFN), which could exsolve Ni–Fe alloy nanoparticles, as an anode for hydrocarbon-fueled SOFCs. Fig. 12(a)–(c) show the high-resolution TEM images of the LSFN powder after calcination and reduction for 2 h under a humid H₂ atmosphere at 750 °C, where it can be seen that small nanoparticles with an average diameter of about 20 nm were uniformly dispersed on the sample surface. Also, the particles in Fig. 12(b) have clear lattice fringes, proving that the nanoparticles were single crystals. As shown in Fig. 12(d) and (e), the LSFN symmetric cell operated in wet C₃H₈ and CH₄ at 750 °C, with an open circuit voltage of 1.18 and 1.0 V and MPDs of 400 and 230 mW cm⁻², but the MPDs of the LSF symmetric cell were just 80 and 240 mW cm⁻², respectively. Fig. 12(f) shows that LSFN had good long-term stability for hydrocarbon fuel oxidation compared with LSF. After more than 36 runs, only a slight power reduction and a small amplitude fluctuation were observed. It showed excellent stability in wet C₃H₈ fuel at 750 °C, while the output current of LSF showed a large fluctuation trend.²⁰⁶ Tang et al. prepared an anode material, i.e., (La_0.4Sr_0.4)(Ti_0.85Ru_0.07Ni_0.08)O_3−δ (L_0.4STRN), with good catalytic performance and stability by doping Ru in La_0.4Sr_0.4Ti_0.85Ni_0.15O_3−δ (L_0.4STN). It was found that the doped Ru could avoid the agglomeration of the exsolution Ni nanoparticles, generating a large number of defects in the perovskite, and thus increasing the amount of oxygen vacancies. The researchers also prepared single cells loaded with L_0.4STN and L_0.4STRN (R-L_0.4STN and R-L_0.4STRN) after reduction. The MPD of the R-L_0.4STN|SDC|LSGM|LSCF cell was 445 mW cm⁻², using wet CH₄ as the fuel at 800 °C, which was better than that of the R-L_0.4STN cell of 202 mW cm⁻². In the stability test, the cell with R-L_0.4STRN retained a voltage of about 0.78 V, while that of the R-L_0.4STN cell decreased significantly at a rate of 0.379% h⁻¹. The cell with R-L_0.4STRN exhibited excellent carbon deposition resistance, which may be due to the synergistic effect of the well-distributed Ni nanoparticles and doped Ru.²⁰⁷

Fig. 12 (a) Low-magnification and (b and c) high-magnification TEM images of La_0.6Sr_0.4Fe_0.9Ni_0.1O_3−δ powder after reduction. (Inset: Fast Fourier transform pattern of selected region.) (d and e) I–V–P curves of cells with different anodes: LSFN (d) and LSF (e). (f) Current density of the cell with different anodes tested at a constant voltage of 0.7 V in different electrodes and humidified C₃H₈ as a function of time. (Reprinted from ref. 206 with permission from Elsevier Ltd.)

Although many studies have constructed heterostructures via exsolution to improve the catalytic ability of perovskites, most of the exsolution species are limited to reducible metal cations, such as nickel, cobalt, iron and noble metals. He et al. reported an improved exsolution method for the co-exsolution of active oxides and metal nanoparticles from a perovskite material, i.e., La_0.8Ce_0.1Ni_0.4Ti_0.6O_3−δ (LCeNT), and constructed a metal cerium titanate heterostructure, CeO₂–Ni@LCeNT, with stable activity to enhance the catalytic performance. After the air annealing treatment of LCeNT, CeO₂ cubes grew on the support, and its morphology could be adjusted by changing the annealing temperature, while the nickel nanoparticles continued to exsolve during the subsequent reduction process. The MPD of the cell with the CeO₂-Ni@LCeNT anode reached 642 mW cm⁻² at 900 °C in CH₄ (3% H₂O). When the stability test was carried out in wet CH₄ at 850 °C to evaluate its carbon deposition resistance, it was found that the current density of the cell decreased slightly before the conversion of CH₄ reached equilibrium after running in CH₄ (3% H₂O) for about 1 h. In the subsequent 19 h running time, the current density stability did not deteriorate significantly. When the working time was delayed to 105 h with the same cell, the current density of the cell was maintained at about 1 A cm⁻², which proved that the CeO₂–Ni@LCeNT cell had outstanding stability. The SEM and STEM tests showed that there was no fibrous carbon on the anode, confirming that CeO₂–Ni@LCeNT has an excellent anti-coking performance. The in situ exsolution of CeO₂ and Ni nanoparticles embedded in the LCeNT substrate to form a strong particle–substrate interaction and a strong structure is the main reason for the excellent anti-coking performance.¹⁷³Table 4 presents a summary of some typical in situ exsolution perovskite materials and their electrochemical performances.

Table 4 Summary of electrochemical performance of SOFCs with different in situ exsolution perovskite anode materials

No.	Anode	Temperature (°C)	Fuel	MPD (mW cm^−2)	Long term (h)	Ref.
1	(La0.7Sr0.3)(Cr0.85Ni0.12Fe0.03)O3−x	850	5000 ppm H2S-syngas	400	—	205
2	Ni-doped La0.6Sr0.4FeO3−δ	750	C3H8	400	35+	206
3	Ni-doped La0.6Sr0.4FeO3−δ	750	CH4	230	—	206
4	La0.8Sr0.2FeO3	750	C3H8	240	30+	206
5	La0.8Sr0.2FeO3	750	CH4	80	—	206
6	(La0.4Sr0.4)(Ti0.85Ru0.07Ni0.08)O3−δ	650	CH4	∼100	25+	207
7	(La0.4Sr0.4)(Ti0.85Ru0.15)O3−δ	800	CH4	202	—	207
8	La0.8Ce0.1Ni0.4Ti0.6O3−δ	900	CH4	306	—	173
9	Pr0.5Ba0.5MnOx	900	Syngas	900	—	200
10	La0.5Ba0.5Mn0.8Fe0.1Co0.1O3−δ	850	CH4	503	200+	192
11	La0.5Ba0.5MnO3−δ	850	CH4	226	—	192

---

In summary, we discussed the in-depth mechanism of in situ exsolution of perovskite materials and their application advantages in SOFC anodes. Due to the higher surface free energy of perovskites and the interaction with strain energy, transition metal or noble metal doped at the B-site of perovskite can be exsolved in situ out of the lattice as active nanoparticles in a reducing atmosphere. Although the impregnation method is the traditional and effective method for constructing metal nanoparticles on the surface of the anode, it requires multiple cycle processes, and the amount of metal loading is restricted by the volume of open pores in the anode. Moreover, the metal nanoparticles formed by the impregnation technique are unstable at high temperatures because they are in an unrestricted environment. In contrast, the metal nanoparticles formed by in situ exsolution are more uniformly distributed and smaller than that formed by the impregnation method, and the precipitated metal particles are semi-embedded in the perovskite matrix, resulting in strong particle–substrate interactions. Hence, the in situ exsolved nanoparticles are not easy to agglomerate at high temperatures. Moreover, for hydrocarbon SOFCs, the Ni particles on the anode prepared by the impregnation method are severely coked and produce a large number of carbon fibers, and even if Ni particles with a particle size of 30–100 nm are prepared by the vapor deposition technique, they also produce a large number of micron-sized carbon fibers, while the Ni metal particles of the same size exsolved in situ have only a small amount of carbon fiber. This is attributed to the powerful interaction between the exsolved nanoparticles and the perovskite matrix, which prevents the growth of carbon fibers and metal nanoparticle lifting.¹⁹⁰

The early research about the stoichiometric ratio of perovskite oxides mainly focused on A/B = 1, where the exsolution of the B-site cations relied on the inherent reduction driving force. In contrast, A-site-deficient perovskites (i.e., perovskite oxides with stoichiometric ratio of A/B < 1) increase the tendency of the B-site cations to exsolve from the lattice with a wider distribution and better coverage on the surface of perovskites, given that the exsolution restores the perovskite oxides to a stable structure locally, making the stoichiometry of the perovskite conform to ABO₃. In addition, A-site-deficient perovskites contribute to the formation of highly mobile oxygen vacancies. There are many intrinsic and extrinsic factors that affect the nucleation and growth of exsolved nanoparticles. The intrinsic factors include surface properties (surface and interfacial energy of nucleated particles), local strain fields in the perovskite matrix, A-site defects and concentration of oxygen vacancies. External factors include processing time and temperature. Thus, by exploring the above-mentioned influencing factors, we can further understand the mechanism of in situ exsolution of perovskite to achieve more artificial regulation to improve the electrochemical performance of the exsolution system.

The development of in situ exsolution is restricted by the fact that most exsolved species are limited to reducible metals such as Ni, Co, Fe and precious metals, limiting the selection of active species. Thus, increasing the category of species from in situ exsolution is the focus of future research. Also, many researchers have focused on the application of SOFC anodes with in situ exsolution for H₂ fuel instead of hydrocarbon fuels, but the anti-carbon deposition advantages of the mosaic structure constructed by in situ exsolution cannot be ignored, and thus more studies should be carried out on hydrocarbon fuels in the future. Furthermore, most of the operating temperatures of cells are mainly at 750 °C or even higher, and the construction of in situ exsolution systems at a lower temperature remains to be solved.

Bimetallic alloys may have better properties than single metals for in situ exsolution metal component systems, and thus the study of bimetallic-cermet may provide more guidance for the in situ exsolution of perovskite materials. In addition, among the various matrix materials, the cell with a Pr_0.5Ba_0.5MnO_x anode exhibited an outstanding performance, which may deserve more attention and exploration. We believe that the excellent heterojunction structure will greatly improve the catalytic performance and stability of hydrocarbon SOFCs.

2.3 Anode reforming layer materials

In addition to directly changing the composition of the fuel cell anode, considering the cost, simplified system, anti-coking, and selectivity of materials, the assembly of a catalytic reforming layer with high catalytic ability for hydrocarbon fuels on the cell maybe an effective way to improve the anti-carbon deposition stability of direct hydrocarbon fuel SOFCs.

Noble metals have good catalytic performance for the reforming reaction. Wang et al. proved that a 3–7 wt% Ru–Al₂O₃ composite had great catalytic ability for the partial oxidation of methane and CO₂/H₂O reforming reaction. H₂-TPR and TEM results showed that there was a robust interaction between RuO_x and Al₂O₃ in the catalysts, which may be why the 3 wt% Ru–Al₂O₃ catalyst exhibited good catalytic stability. In addition, the carbon deposited on Ru–Al₂O₃ was less than that deposited on Ni–Al₂O₃, indicating that it is easier to eliminate the potential carbon deposited on Ru–Al₂O₃ catalyst.²⁰⁸

By adding an Ru–CeO₂ catalyst layer to Ni–YSZ, Zhan et al. compared the cell performance test results with the test results of the conventional Ni–YSZ anode-supported fuel cell using a propane-air mixture. Fig. 13(a) shows the temperature dependence of propane conversion in the cell operated in 10.7% C₃H₈-air fuel with an Ru–CeO₂ catalyst layer on the YSZ anode. The propane conversion increased from 20% at 450 °C to 80% at 500 °C, and Ru–CeO₂ could catalyze the C₃H₈ partial oxidation at T > 500 °C. The conversion of propane on the Ni–YSZ anode reached 80% at 650 °C, which is 150 °C higher than that of the Ru–CeO₂ catalyst. This indicates that the Ru–CeO₂ layer is a more effective low-temperature catalyst than Ni–YSZ for propane partial oxidation.²⁰⁹ Generally, it is believed that the electrochemical reaction in SOFCs usually occurs in the TPB. Thus, to promote gas diffusion, prolong the length of the TPB and form a continuous ion electron conduction path, Sun et al. prepared a fuel cell with a flower-like mesoporous cerium oxide microsphere/ruthenium nanoparticle catalyst layer by loading Ru on CeO₂ using the impregnation-reduction method. As shown in Fig. 13(b) and (c), most of the CeO₂ particles in the flower-like mesoporous CeO₂–Ru microspheres were monodispersed spherical particles with a hollow structure. It can also be seen in the high-magnification SEM and TEM image in Fig. 13(d) and (e), respectively, that many petal-shaped nanosheets formed an open porous flower-like structure, and the microspheres were hollow structures, which facilitated the diffusion of the fuel. At 600 °C, the MPD of the cell was 654 mW cm⁻², which was much higher than the of the cell without a catalyst layer of 81 mW cm⁻², as shown in Fig. 13(f).²¹⁰ Similarly, Sun et al. synthesized a Pd layer consisting of 3D nanoflowers with a diameter of 1.0–2.0 μm as an anodic catalytic reforming layer on an Ni–YSZ anode by galvanic replacement reaction. The MPDs of the Pd–Ni–YSZ single cell with catalytic reforming layer and single cell with just the Ni–YSZ anode were 196 and 153 mW cm⁻², respectively, when ethanol was used as the fuel, indicating that the Pd catalytic reforming layer increased the MPD by 43 mW cm⁻². Furthermore, the voltage of the cell with the catalytic reforming layer decreased significantly within 59 h, while the cell loaded with the Ni–YSZ anode had a sharp drop in voltage after 6 h, which was because the Pd nanoparticles could act as a barrier layer to prevent diffusion, and thus significantly suppress carbon deposition on the nickel-based anode surface.²¹¹

Fig. 13 (a) Relationship between propane conversion and temperature of Ni–YSZ anode-supported Ru–CeO₂ catalyst layer with 10.7% C₃H₈–18.7% O₂–70.6% Ar as the fuel. Single-cell Ru–CeO₂|PSZ|Ni–YSZ|YSZ|LSCF-GDC, LSCF. (Reprinted from ref. 209 with permission from Elsevier Ltd.) (b and d) SEM images and (c and e) TEM images of the flower-like mesoporous CeO₂ microspheres. (f) Voltage and power density versus current density for the anode-supported SOFCs with and without catalyst layers, tested in 5% iso-octan–9% air–3% H₂O–83% CO₂ at 100 mL min⁻¹ in the anode and ambient air in the cathode at 600 and °C. (Reprinted from ref. 210 with permission from Elsevier Ltd.)

Wang et al. prepared mesoporous cerium oxide Ce_0.8Sm_0.22O_1.9 (SDC) with a high specific surface area as the catalytic layer of CH₄-fueled SOFCs. After calcination at 700 °C for 3 h, the SDC powder still had a mesoporous structure with a high surface area of 77 m² g⁻¹ and a large pore volume of 0.2276 cm³g⁻¹, which is conducive to free gas diffusion. After impregnation of the mesoporous SDC catalyst layer with Ru, the supported cell exhibited an excellent MPD of about 462 mW cm⁻² at 650 °C.²¹² Zhao et al. applied a reforming layer composed of Ni loaded on La-doped ceria such as La₂Ce₂O₇ (LDC) and La_1.95Sm_0.05Ce₂O₇ (LSDC) on an Ni–Ce_0.8Sm_0.2O_2−x anode. In wet CH₄ at 650 °C, the MPD of the conventional cells was only 580 ± 20 mW cm⁻², but after the addition of the Ni-LDC and LSDC layers, the MPD of both increased to 699 ± 20 mW cm⁻² and 639 ± 20 mW cm⁻², respectively. In the stability test at 650 °C and 0.2 A cm⁻² in wet methane, the voltage of the conventional cell decreased significantly within 10 h, while the cell with any one of the catalyst layers was stable for 26 h.²¹³

When used in monometallic catalysts, the catalytic activity of metals such as iron, cobalt, and copper is relatively low, but the addition of other metals to form new substances or bimetallic catalysts may result in catalysts with good performance, and among them, cobalt is the most appropriate metal for changing nickel-based catalysts because the electronic interaction between nickel and cobalt contribute to improving the dispersion of metal alloys and carbon deposition resistance.²¹⁴

Due to the low reforming performance of Cu-based catalysts, a second phase catalyst needs to be added when hydrocarbons are fed. Therefore, Jin et al. synthesized spinel oxide Cu_1.3Mn_1.7O₄ as an anode inner reforming layer for an Ni–SDC anode-supported SOFC for direct operation in methane, and after the in situ reduction of CH₄, well dispersed nano-Cu metal mesh was obtained in the MnO matrix. The MPDs of the cells reached 242 and 311 mW cm⁻² at 650 °C and 700 °C, respectively. The power output of the cell increased by about 30% compared with the cell without an internal reforming layer under the same test conditions. For the cells with a Cu_1.3Mn_1.7O₄ internal reforming layer, the cell voltage became more stable for 40 h. In contrast, the cell without a Cu_1.3Mn_1.7O₄ catalyst layer rapidly failed after working for 16 h. In the case of the CuB₂O₄ (B = Fe, Mn, Cr, Ga, Al, etc.) spinel oxides after reduction, the stability of the spinel oxides, metal oxides and the strong interaction between the metal oxides and Cu contributed to the outstanding catalyst durability.²¹⁵

Because Ru_xCo_1−x supported on a CoO catalyst has high activity for dissociating H₂O/CO₂ and Ni has catalytic activity for CH₄, Chen et al. dispersed Ru and Ni on CeO₂ to form a catalyst. It is expected that Ru and Ni can catalyze H₂O and CH₄, respectively, and thus the coupling of the two dissociated substances can play a role in the synergistic activation of H₂O and CH₄. Based on this, these researchers prepared a Ce_0.90Ni_0.05Ru_0.05O₂ (CNR) reforming catalyst supported on a BaZr_0.1Ce_0.7Y_0.1Yb_0.1O_3−δ (BZCYYb) anode. The single cell composed of the of PrBa_0.5Sr_0.5Co_1.5Fe_0.5O_5+δ (PBSCF) cathode and SDC electrolyte ran directly on nearly dry (containing only 3.5% H₂O) methane at 500 °C, and its MPD was 370 mW cm⁻². This is obviously higher than the 142 mW cm⁻² reported as the highest value for a proton ceramic fuel cell fueled with steam-rich CH₄ and 78% higher than the fuel cell without catalyst layer, and the cell showed no signs of coking after about 550 h continuous operation.²¹⁶ However, catalyst materials doped with noble metals such as Ru,²¹⁷ Pd, and Pt²¹⁸ are limited by their high cost, and thus cannot be applied widely.

Therefore, researchers began to avoid using noble metals, such as preparing nanostructured materials with other metals or adding non-noble metal elements in the anode catalytic reforming layer. Kim et al. innovatively coated a cerium oxide layer modified with metal nano-catalysts by cathodic electrochemical deposition (CELD). The first CELD was performed in 0.05 M Ce + Sm nitrate solution with 5% Sm content to obtain the SDC layer. A cathodic potential of 0.65 V per saturated calomel electrode (SCE) was applied to the nickel electrode to reduce the NO₃ anions or dissolved O₂ to produce OH⁻, with solid (Ce, Sm)(OH)₃ on the electrode surface. The low deposition rate of 0.65 V per SCE could produce a more nanostructured cerium oxide morphology, ensuring that the catalytic base had the best electrocatalytic activity. Through drying or heat treatment in air, these deposits were oxidized to obtain an oxide coating, which became a petal-like SDC nanostructure composed of high-density, randomly oriented nanosheets. Ni(OH)₂ was obtained by the second CELD in the same way, and finally the deposited layer was in situ transformed into SDC nanostructures and nickel metal particles were dispersed on it in a reduced H₂ atmosphere. The catalytic performance of the SDC nanostructures modified by 12 at% Ni nanoparticles was tested by gas chromatography (GC) at 550–700 °C for CH₄ steam reforming (3% H₂O/CH₄). The results were expressed as the ratio of generated H₂ concentration to the theoretically predicted H₂ concentration. The ratios of the anodes after two-step CELD at two temperatures reached 54.1% and 79.2%, while that of the uncoated samples was 3.5% and 12.4%, respectively. The bare nickel catalyst layer had a higher EA value of 113.25 kJ mol⁻¹, while that of the CELD anode layer was just 35.75 kJ mol⁻¹, which was lower than the EA value of the typical Ni–CeO₂ (63–88 kJ mol⁻¹), indicating that the reforming layer significantly promoted CH₄ steam reforming. Also, the cell loaded with the reforming layer could operate stably at 650 °C for 90 h.²¹⁹ Hua et al. introduced Ni–Sn bimetallic nanoparticles in the traditional nickel-based anode to form a porous and flexible nickel foam-loaded Ni–Sn/Al₂O₃ nanocluster reforming catalytic layer. The MPD of the cell was as high as 0.946 W cm⁻² at 850 °C in a CH₄–CO₂–200 ppm H₂S atmosphere. In CH₄–CO₂ containing 200 ppm H₂S, the conversion of CH₄ was still as high as 95.6%. The output voltage of the cell with the conventional Ni–YSZ anode without this catalytic reforming layer decreased continuously during the 48 h stability test, while the CH₄ conversion and CO selectivity of the cell loaded with this reforming layer remained at about 95% during the 48 h test.²²⁰

Zhao et al. deposited Ni–Mo–Ce_0.5Zr_0.5O_2−δ (Ni–Mo–CZ) on the top of a conventional Ni–YSZ anode as an internal catalytic reforming layer. The Ni–Mo–CZ layer effectively increased the gasoline conversion and increased the H₂ and CO yields of Ni–YSZ at 750 °C. The gasoline conversion of the Ni–YSZ anode without the Ni–Mo–CZ catalyst was just 40%, and the yields of H₂ and CO were 5% and 8%, respectively. When the nanoscale Ni–Mo–CZ catalyst was added, the gasoline conversion rate increased from 40% to 57%, and the H₂ and CO yield increased to 17% and 22%, respectively. However, the thicker catalyst layer resulted in higher mechanical stress after sintering, and thus only a limited amount of catalyst (about 20 mg) could be loaded on the traditional button cell, which prevents the improvement in the cell performance. Therefore, the researchers produced a tubular single-cell device to load a larger amount of catalyst (about 80 mg), and the gasoline conversion rate increased to 78%, the CO yield increased to 55%, and the H₂ yield reached 40%. In addition, the MPD of the tubular single cell without the Ni–Mo–CZ catalyst was 144 mW cm⁻², which increased to 254 mW cm⁻² by filling with catalyst. The single cell without the catalyst was damaged after running for 6 h, while the stability of the button cell with the catalyst layer decreased after running for about 10 h. However, even after 15 h operation, the tubular cell with catalyst did not show a significant performance degradation, which was 42% longer than the button cell life. Thus, the above-mentioned test results show that the cell with Ni–Mo–CZ had a good catalytic reforming performance using gasoline as fuel, and for tubular single cell, one of the advantages is that more catalyst can be added to the reforming layer, further improving the cell performance.²²¹Table 5 presents a summary of some typical anode reforming layer materials and their electrochemical performances.

Table 5 Summary of electrochemical performance of SOFCs with different anode reforming layer materials

No.	Catalyst layer	Temperature (°C)	Fuel	MPD (mW cm^−2)	Long term (h)	Ref.
1	Ru–Al2O3	750	CH4 : O2 = 4 : 1	581	6+	208
2	Ru–Al2O3	750	CH4 : H2O = 2 : 1	489	—	208
3	Ru–Al2O3	750	CH4 : CO2 = 2 : 1	478	—	208
4	Ru–CeO2	750	10.7% C3H8 +18.7% O2 + 70.6% Ar	480	—	209
5	Flowerlike CeO2–Ru	600	5% isooctane + 9% air + 3% H2O + 83% CO2	654	—	210
6	3D flower-like Pd	750	Ethanol	196	—	211
7	Ru–SDC	650	CH4	462	—	212
8	Ni–La2Ce2O7	650	CH4	699	26+	213
9	Ni–La1.95Sm0.05Ce2O7	650	CH4	639	26+	213
10	Cu1.3Mn1.7O4	650	CH4	242	60+	215
11	Ce0.90Ni0.05Ru0.05O2	500	CH4	370	380+	216
12	Ni NP–SDC	650	CH4	—	6+	219
13	Ni–Sn–Al2O3	850	200 ppm H2S –CH4 + CO2	946	48+	220
14	Ni–Mo–CZ	750	Gasoline/air	254	15+	221

---

In summary, the addition of a reforming layer with high catalytic activity has been demonstrated to be an effective way to improve the performance and long-term stability of hydrocarbon-based SOFCs. To prepare the reforming layer, in addition to finding a catalyst that can efficiently catalyze the reforming reaction, it is also necessary to optimize the microstructure and solve the gas diffusion problem caused by the reforming layer, such as preparing the catalyst in a porous form, which is an important direction to overcome the shortcomings of the reforming layer. Furthermore, it is noteworthy to match the thermal expansion coefficient of the reforming layer with the anode material, where an inappropriate thermal expansion coefficient will cause the reforming layer to fall off from the cell in the high-temperature operating environment.

3. The kinetic anode models

Considering the possibility of competing or co-existing reforming reactions on SOFC anodes, studying the mechanism of reforming hydrocarbon fuel is the first step in realizing the direct application of hydrocarbon fuel. Therefore, the modeling research based on the first-principles method is conducive to the design of SOFCs suitable for various fuels before the actual experiments, as well as the interpretation of experimental phenomena, guiding the subsequent experiments to a certain extent. In this part, we summarize some of the reaction mechanisms that have been discovered and give an in-depth description of the causes and solutions of anode carbon deposition.

Yang et al. calculated the process by DFT that the adsorbed *H₂O is converted to *OH, while *OH can react with the *C near the Ni/BaO interface to form *CHO, which is then electrochemically oxidized at the TPB. It was found that water associated with the Ni/BaO interface plays a critical role in increasing the tolerance to coking and inactivation. They introduced Ni, YSZ and BaO powder samples into wet argon (3 vol% H₂O), respectively, and found that the weight of the Ni and YSZ powders remained almost unchanged based on the thermogravimetric analysis results, but a large degree of weight alteration was observed in the BaO powder test, indicating the good water absorption of BaO. Similarly, after introducing Ni–BaO samples into wet argon, peaks around 1600 cm⁻¹ were observed in the Raman spectrum, corresponding to the bending mode of water, which also confirmed the water absorption capacity of BaO. Furthermore, the cell with or without the BaO-modified anode was introduced into dry C₃H₈ fuel gas at a constant current density of 500 mA cm⁻² at 750 °C, and it was found that the voltage of the cell without the BaO-modified anode fell to 0 V in 0.5 h, while the cell with the BaO-modified anode was stable at 0.8 V for more than 100 h. This suggested that water-mediated carbon removal reactions were more advantageous at the Ni/BaO interface. Also, it was found that water dissociation occurred at BaO and carbon generation occurred at the Ni site of Ni/BaO boundary. The results showed that the performance and properties of resistance to carbon deposition were largely dependent on the direct involvement of the Ni/BaO interface.²²²

Based on the principle that noble metal reactions increase the activity of water–gas shift reactions and formaldehyde oxidation at low temperature, Yang et al. chose the relatively stable Ni(111) surface to model and used the VASP software to perform ab initio calculations. Combined with the PAW method and CI-NEB method, the dissociation adsorption of H₂O on MgO, and the reaction between surface diffusion and carbon deposition of intermediate materials at the Ni/MgO interface were calculated by simplifying the reaction mechanism. The results of thermogravimetric analysis and temperature-programmed desorption of CO₂ measurements indicated that the MgO nanoparticles on the surface of Ni ceramics were beneficial to capture H₂O and CO₂, significantly increased the number of *OH substances for Ni stably, and thus enhanced the coke resistance. The cell with the MgO-modified anode could operate in CH₄ (3 v% H₂O) fuel at 800 °C for more than 325 h. Also, this principle can be directly applied to various fields such as direct methanol fuel cells and other catalytic reactions.¹⁴⁹

Ahmed et al. used mathematical models of finite size to analyze the temperature, current, voltage and internal distribution of the components in SOFC anodes, considering the three reactions, i.e., water–gas shift, methane steam reforming and reverse methanation (RMTN). Based on the fact that Xu et al. found that the activation energy of RMTN is very close to the activation energy of methane disappearance,²²³ they believed RMTN should be considered in addition to the water–gas shift reaction when kinetic rate expressions are used to simulate the methane steam reforming reaction because the lack of simulation calculations about RMTN reaction leads to prediction bias for characteristic variables such as current density of SOFC systems. This study found that from the perspective of SOFC degradation, a fast kinetic rate did not necessarily lead to the optimal temperature distribution and current density; on the contrary, the use of an anode with a slightly lower kinetic rate could improve the operational stability.²²⁴

The reactive force field molecular dynamics simulation (ReaxFF-MD) model can correctly describe a series of complex physicochemical processes in the TPB of SOFCs, and Duin et al. used ReaxFF-MD to study the transport performance of the protons in Y-doped BaZrO₃ and oxygen ions in YSZ, and verified the feasibility of the model in SOFC applications through experiments.^225,226 Also, Lu et al. used ReaxFF-MD to clearly display the process of coke formation on the anode and simulated the number of C and O atoms in Ni crystals over time, as shown in Fig. 14, explaining the increase in the amount of methane in the early stages of methanol and methane absorption. The results showed that alkyl groups are more likely to stick to the surfaces during adsorption. Initially, the C in Ni(100) provides the driving force, causing changes in the formation of new carbon cluster phases in the later stage, resulting in Ni(100) damage. Finally, more and more C is absorbed on the surface, leading to the minimization of the reforming process. Thus, doping the surface of nickel with other elements can increase the steric hindrance of the carbon adsorption process.²²⁷

Kim et al. used FIB-SEM tomography and finite element analysis to perform three-dimensional reconstruction of reduced and re-oxidized Ni–YSZ anode functional layers, analyzed their morphology and stress distribution, and explained the very common stack fault phenomenon. As showed in Fig. 15, by virtually heating at 500–900 °C to obtain the thermal stress curves, they analyzed the thermal stress distribution of the re-oxidized anodes. According to the simulation results, it can be seen that the thermal stress is concentrated at the NiO–YSZ interface, and the stress at the NiO interface is higher than that at the YSZ interface. This means that there is a high probability of rupture in the NiO region.²²⁸

Fig. 14 Process of coke formation on the anode and change in the number of C and O atoms with time in the Ni crystal. (Reprinted from ref. 227 with permission from Elsevier Ltd.)

Fig. 15 Stress distribution diagram under different heating conditions. (Reprinted from ref. 228 with permission from Elsevier Ltd.)

Guo et al. found that an all porous solid oxide fuel cell (AP-SOFC) exhibited excellent resistance to carbon deposition in experimental tests,^229,230 and further Xu et al. developed a 2D model of AP-SOFC to study the effects of key parameters such as O/C ratio and electrochemical reaction rate on carbon deposition in SOFCs. They found that the process of diffusion of O²⁻ from the cathode to the anode improved the O/C ratio of the anode, which can effectively inhibit methane coking and carbon deposition. In addition, the data obtained through design experiments showed that in different CH₄ molar fraction ratios of fuel gases, the CH₄ conversion efficiency of the anode-supported AP-SOFC was above 97%, and the outlet syngas was abundant, thus, AP-SOFCs can co-generate electricity and syngas without carbon deposition. The anode support of AP-SOFC can obtain a greater MPD than the electrolyte support (1607 vs. 146 mW cm⁻²), and for the same cell, the fuel composition and fuel flow rate affect the O/C ratio of the anode surface, which in turn affects the electrochemical performance.²³¹

Combined with Raman spectroscopy, Bkour et al. found that well-dispersed MoO_x substances, especially Mo=O bonds, on Ni nanoparticles can promote the coke resistance and catalyst stability of Ni–Mo–YSZ catalyst samples. Calculations based on DFT showed that compared with Ni–YSZ catalysts, the Mo-doped Ni surface sites on the YSZ support are thermodynamically feasible and can enhance the carbon tolerance by increasing the activation barriers of C–H bond cleavage and C–C coupling reactions, as show in Fig. 16. In addition, for the projected density of states (PDOS) of the Ni 3d orbitals in Ni(111) with Mo, the d-band center of the Ni atom is located near the Mo atom away from the Fermi level, thus reducing the electron interaction between carbon and nickel. Therefore, less carbon will be adsorbed in the Ni–Mo–YSZ catalyst system.¹¹² The DFT calculation result of the Ni–Mo–YSZ catalyst was similar to the study of the Ni–Sn–YSZ catalyst,¹¹⁶ where the activation energy barrier of C–O was much lower than that of C–C on the surface of the Ni–Sn alloy calculated by DFT, and thus the C atom will be preferentially oxidized rather than form carbon deposits.

Fig. 16 In Ni–Mo/YSZ and Ni/YSZ model systems (A) transition states of C–C coupling paths and (B) their corresponding configurations of initial state (IS), transition state (TS) and final state (FS) in partial representation. The purple, silver, red, brown, light green and dark green spheres represent Mo, Ni, O, C, Zr, and Y atoms, respectively. (Reprinted from ref. 112 with permission from Elsevier Ltd.)

Audasso et al. proposed a two-dimensional modeling tool to simulate the performance of SOFCs under direct internal methane reforming and introduced a coefficient matrix (σ) to represent the distribution of active sites. The same method was also used to determine the optimized catalyst distributions. For the DIR process, the code SIMFC (SIMulation of Fuel Cell) considered two different approaches proposed in the literature, i.e., thermodynamic equilibrium and reaction kinetics. The SOFC crosscurrent configuration supported by a single cell anode was also considered, and (i) how to simulate the effect of reforming catalyst inactivation on SOFC performance and (ii) how local control could be used to detect optimized catalyst distributions to provide minimal thermal stress to the cell. In addition, the σ coefficient matrix was appropriately adjusted to obtain the optimal distribution that is easily achievable in the manufacturing process. Studies found that steep temperature distributions lead to higher mechanical stress, which leads to creep formation and ultimately leads to the degradation of the overall anode.²³² This model is reliable for guiding SOFC direct internal reforming studies, and because of the simplicity of the calculation, it can be integrated into more complex simulation software such as Aspen Plus.

Su et al. synthesized H₂-reduced Sr₂ZnMoO₆ (R-SZMO) as an SOFC anode material. Due to the inherent anion-Frenkel defect pairs in the material, it showed good catalytic capacity and coking resistance. According to DFT calculations, from the perspective of local structure, with the help of the electroactive O site on the optimal SZMO(110) surface, the optimal adsorption of key intermediates methyl, methylene and formaldehyde was located near O and Mo sites, showing excellent catalytic capacity for methane. Meanwhile, high methane catalytic capacity triggers high energy output and carbon deposition resistance. When CH₄ was used as the fuel gas, compared with other typical double perovskite anode materials (Sr₂CoMoO₆, Sr₂MgMoO₆, and La_0.75Sr_0.25Cr_0.5Mn_0.5O₃), the cell with the Sr₂ZnMoO₆ anode exhibited the highest output power and could operate stably in CH₄ for 110 h without significant degradation.²³³

Overall, molecular dynamics, finite element simulation, supplementary mathematical models and other simulation methods are comprehensively used to simulate and calculate the causes and solutions of anode carbon deposits in SOFCs. They lay a good theoretical foundation for further understanding the anode reaction kinetics and the design of carbon-resistant anodes. In the future, the reaction mechanism and reaction kinetics of hydrocarbon fuels on carbon-resistant anodes should be studied in depth, and the theoretical calculation or simulation model can be verified by corresponding experiments under certain conditions.

4. Future anode research direction

4.1 Multilayer anode

It should be noted that the electrochemical reactions in high-performance electrodes mainly occur in the range of about 10 μm from the electrolyte, and consequently the anode can be divided into two layers, i.e., a thinner functional layer for electrocatalysis and a thicker conductive layer for current harvesting. Consequently, by considering the desired characteristics of each layer separately, the catalytic performance of the anode can be maximized.⁶¹

Ye et al. successfully prepared a dual-layer structure containing a Cu–CeO₂–YSZ support anode and Ni–ScSZ functional anode by acid leaching of nickel and wet impregnation method. In the case of ethanol fuels, each ethanol molecule needed to react with six times the oxygen ions compared to each hydrogen molecule, which meant more charge transfer was required. Therefore, the performance of the Cu–CeO₂–ScSZ composite anode mainly depended on the conductivity and connectivity of Cu with CeO₂ particles in the porous matrix for ethanol fuels. The anode with a greater copper loading provided further conductive paths for electrons as well as connected additional highly efficient CeO₂ catalytic active reaction sites. The additional anode functional layer significantly improved the effective reaction region because the limited inherently porosity of the original single-layer anode inhibited the total loading of copper and cerium dioxide. The MPDs of the Cu–CeO₂–YSZ|Ni–ScSZ|ScSZ|PCM single cell reached 604 and 408 mW cm⁻² in hydrogen and ethanol steam at 800 °C, which were higher than 372 and 222 mW cm⁻² of the Cu–CeO₂–ScSZ anode in their previous work, respectively.²³⁴ In addition, they found that the double-layer SOFC anode could stably output power density for more than 50 h in ethanol steam in 0.7 V at 800 °C. Also, the cell did not exhibit the phenomenon of layer-to-layer peeling in the 50 h test environment at high temperature given that this dual-layer anode structure was fabricated by multi-layer tape casting and co-sintering processes. However, they also emphasized that the long-term stability of the anode in hydrocarbon fuels still needs further study. The great electrochemical performance improvement proved the possibility of enhancing the properties of Cu–CeO₂ anodes by optimizing the microstructure of dual-layer anodes.²³⁵

Hao et al. prepared an La_0.2Sr_0.7TiO₃–Ni–YSZ functional gradient anode (FGA) support SOFC using the co-tape casting method and sintered it using the field-assisted sintering technique (FAST). Specially, LST–Ni–YSZ FGA was a composite anode that possessed a variable anode component content and porosity from the anode to the electrolyte. They confirmed that the FGA structure was acquired and maintained excellent stability after the FAST process through SEM images, successfully avoiding the deformation and delamination that occur after conventional sintering. This SOFC anode exhibited the MPD of 600 mW cm⁻² at 750 °C in methane fuel, and after 70 h of operation, no carbon deposits were detected through Raman spectroscopy. The authors emphasized that unlike the traditional sintering processes, distortion and delamination between the layers of the anode could be avoided due to the use of FAST technology to sinter LST–Ni–YSZ FGA. Thus, the phenomenon of interlayer stripping did not occur during the 70 h test.²³⁶

Wang et al. prepared an Ni–Fe/Ni–YSZ dual-layered anode by using the tape casting/lamination/sintering method. The electrochemical impedance measurement and relaxation time distribution analysis indicated that the cell with a double-layer straight hole path allowed rapid gas phase transport, thereby reducing concentration polarization, and also decreasing activation polarization by enhancing the accessibility of the electrochemical reaction sites. Furthermore, after eight redox cycles, the dual-layered anode support SOFC maintained its structure, while the Ni–YSZ single-layer anode support SOFC was damaged after one redox cycle. The results demonstrated that the dual-layered anode exhibited good long-term stability, and they also noticed that for the same cell, according to the cross-sectional SEM/EDX images, the re-oxidation stage involved the separation of the alloy phase during a redox cycle.²³⁷

The construction of multilayer anodes contributes to the diffusion and reforming reaction of hydrocarbon fuel, which can realize a greater loading of active catalytic component and appropriate porosity. However, it is worth noting that multilayer anodes may undergo interlayer stripping during cell testing, Therefore, this problem needs to be considered when constructing multilayer anodes.

4.2 Bimetallic-perovskite

The different carbon deposition resistances observed between different nickel cermets indicate the importance of the anode ceramic composition, and the TPB region of the different ceramics is also important in terms of carbon deposition resistance.

The novel A-site layered PBMO exhibited a superior performance and remarkable stability under the reducing conditions of SOFCs. Sengodan et al. annealed Pr_0.5Ba_0.5MnO_3−δin situ at 800 °C under a hydrogen atmosphere to make layered perovskite PrBaMn₂O_5+δ (PBMO) and used it as an SOFC anode for the first time. They further impregnated the Co_0.5Fe_0.5 catalyst solution in the PBMO skeleton, and the Co–Fe–PBMO anode SOFC delivered MPDs of 1770, 1320 and 570 mW cm⁻² in humidified H₂, C₃H₈ and CH₄ at 850 °C, respectively. They also systematically studied that when the pure H₂ was changed to H₂ containing 30 ppm H₂S, the cells with layered PBMO with Co–Fe catalyst did not show significant degradation of the voltage during 50 h of operation, indicating that the layered PBMO anode had superior sulfur resistance. In addition, the Co–Fe–PBMO anode cell was subjected to a constant current load of 0.2 A cm⁻² at 700 °C and operated in C₃H₈ for more than 500 h without significant voltage degradation.²³⁸ Li et al. infiltrated Pr_0.5Ba_0.5MnO_3−δ nanoparticles in the YSZ backbone, and then subjected it to redox cycles to form nanoparticle-based alloy Co–Mo–PBMO₅–YSZ anodes. Also, when tested in simulated biogas, the Co–Mo–PBMO₅–YSZ anode cells exhibited an MPD of 1100 mW cm⁻², while that of the Co–YSZ anode cell was only 800 mW cm⁻². Moreover, the Co–Mo–PBMO₅–YSZ anode cell showed no performance degradation in the 30 h stability test, while the Co–YSZ cell manifested a poor performance due to coking.²³⁹ These experimental results showed that the layered perovskite PBMO anode exhibited good redox stability and excellent resistance to coking in hydrocarbon fuels and sulfur pollution, which established new design principles for the optimization and stabilization of nanoparticle-based catalyst SOFC.

The excellent catalytic activity and anti-sintering of alloy-ceramic catalysts are closely related to the strong interaction between the alloy and support. We believe that bimetallic materials combining multiple types of perovskites may be the focus of future research.

4.3 Application of single-atom catalyst anode in intermediate-temperature SOFCs

Designing single-atom catalysts (SACs) by managing the catalytic sites at the atomic scale can result in high-quality activity and minimize the use of metals, and SACs have unique selectivity compared to nanoparticle catalysts due to the lack of collection sites. The presence of high surface energies in the unsaturated and/or low-coordination active sites of SACs can lower the energy barrier and enhance the charge transfer.^240–242

Tang et al. reported the synergistic effect of two groups of single-atom sites (Ni₁ and Ru₁) anchored on the surface of CeO₂ nanorods, which were highly active for reforming methane with CO₂ to generate syngas (CO + H₂). By comparing the three groups of catalysts of Ce_0.95Ni_0.025Ru_0.025O₂, Ce_0.95Ru_0.05O₂ and Ce_0.95Ni_0.05O₂, they found that Ce_0.95Ni_0.025Ru_0.025O₂ catalysts had the lowest apparent activation barrier and the highest CO and H₂ conversion rates. Computational studies had found that the synergistic effect of Ni₁ and Ru₁ sites was manifested in the complementary functions of (1) Ni₁ (activated CH₄) and Ru₁ (activated CO₂), and (2) Ni₁ site activated CH₄ and generated H atoms, and then coupled H atoms at Ru₁ sites to form H₂. And they also found that the single-atom sites anchored on the surface of CeO₂ remained monodisperse during catalysis up to 600 °C, which could basically meet the application of single-atom catalysts in hydrocarbon SOFC.²⁴³

Similarly, Wu et al. successfully modulated oxygen vacancies (O_V) at different concentrations of CeO₂ by replacing the Ce⁴⁺ cations with smaller cations M (M = Mg, Co, Zn). Also, Ni₁/MCeO₂ SAC was used for methane dry reforming reactions, which maintained a highly active state in the catalytic processes up to 800 °C. The Ni SACs exhibited extremely high activity in the activation of C–H bonds; however, carbon deposited on Ni caused their inactivation, but the in situ adsorbed oxygen (O_ad) species adjacent to the Ni SACs were efficient reactants to eliminate carbon deposition, which could react with the deposited carbon to form gas-phase CO and CO₂, thereby removing the carbon deposits. Also, O_ad species were typically generated on the oxygen vacancies (O_Vs), specifically, O_V activated CO₂ to produce CO and O_ad species, and thus the issue of carbon deposition can be solved by building enough O_V around the Ni SACs. During the methane dry reforming reaction, carbon deposition decreased by 50% as the O_V concentration increased from 21.9% to 30.8%. The synergistic combination of the effective C–H activation function of the Ni SAC and the CO₂ activation function of O_V led to a highly active and efficient carbon removal process for methane conversion in the O_V–SAC-catalyzed system.²⁴⁴

For fuel cells, SACs with high ORR performance are mainly used in proton exchange membrane fuel cells (PEMFC) and anion exchange membrane fuel cells (AEMFC). However, few researchers have applied single-atom catalysts in SOFCs, which is mainly because their high operating temperature leads to severe single-atom aggregation.²⁴⁵ However, it is worth noting that reducible oxide supports, such as Fe₂O₃, CeO₂, and TiO₂, are particularly suitable for stabilizing single atoms due to the strong covalent metal support interactions (CMSIs), and high-temperature thermally stable single-atom catalysts can be obtained through this interaction.²⁴⁶ Li et al. linked SACs with reversible proton-conducting solid oxide cells (R-PSOCs) and selectively anchored Pt atoms to position B in Pr₄Ni₃O_10+δ by a CMSI-based mechanism. The obtained Pt single-atom catalyst could withstand a high temperature of 700 °C and was highly active against oxygen reduction and oxygen evolution after 10% H₂O–air and 10% H₂O–10% CO₂-air for 100 h, respectively. They also assembled an Ni–BZCYYb1711|BZCYYb1711|Pt–Pr₄Ni₃O_10+δ cell, and its electrochemical performance was improved by almost 100% compared with Ni–BZCYYb1711|BZCYYb1711| Pr₄Ni₃O_10+δ, and this catalyst preparation method is suitable for industrial-scale production.²⁴⁷

Therefore, we believe that single-atom anode SOFC catalysts based on the CMSI mechanism and combined with a large number of mature single-atom catalyst systems in hydrocarbon research will have broad prospects and commercial value in the future.

5. Conclusions and perspectives

In summary, the process of carbon deposition with the classification of carbon species, carbon detection methods, and strategies to solve anode carburizing of SOFCs were discussed. The use of alternative anodes is more promising than other strategies such as reducing the SOFC operating temperature and adjusting the O/C and H/C ratios of fuels. Therefore, a comprehensive overview of several decades of development and recent trends of carbon-resistant anode materials were provided, which were classified as follows: (1) bimetallic-cermet materials, (2) ceramic materials, and (3) anode reforming layer. Also, we briefly summarized the theoretical calculations and simulation models of anodes for hydrocarbon SOFCs, aiming to promote the research and development of alternative anodes. It can be seen that the commercialization of hydrocarbon SOFCs has slowed down due to many major technical challenges, where among them, the carbon deposition generated at the anodes during the operation in hydrocarbon fuel seriously decreases the catalytic ability and durability of the anodes, and damages the structure of the entire cell. Thus, to overcome these challenges to obtain efficient and stable carbon-resistant anodes, several research directions are proposed, as follows:

(1) A large number of studies proved that bimetallic-cermet or solid solution of metal-alkaline earth metal oxide ceramic anodes have excellent electrochemical performances and long-term stability in hydrocarbon fuels. Future cermet anode research should consider the synthesis of more types with bimetallic alloys and solid solution of metal-alkaline metal oxides by non-hydrogen reduction pathways, such as reduction by carbon or carbon monoxide, given that hydrogen limits the reduction of some metal oxides.

(2) The catalytic activity and anti-coking of bimetallic-cermet are related to the interaction between the alloy and carrier ceramics, and for the ceramic components, the following two directions can be considered: (I) constructing bi-ceramic components (e.g., Ni–YSZ–YDC) and (II) using or synthesizing ceramics that can change their structure under a high-temperature reducing atmosphere, then forming a stable structure containing high conductivity and substantial oxygen vacancies (e.g., Pr_0.5Ba_0.5MnO_3−δ).

(3) The heterojunction structure anode formed by in situ exsolution in perovskite materials will greatly improve the catalytic performance and stability of hydrocarbon SOFCs. (I) A-site-deficient perovskites (i.e., perovskite oxides with stoichiometric ratio of A/B < 1) increase the tendency of the B-site cations to exsolve from the lattice with wider distribution and better coverage on the surface of the perovskite, which is a more efficient method for constructing heterojunction structure anodes. (II) Moreover, combined with sufficient bimetallic-cermet research, the study of alloy metal nanoparticles with high catalytic activity in situ exsolved from perovskite materials should be valued.

(4) In addition to improving the alternative anode material composition, the anode microstructure should be optimized to obtain a greater loading of active catalytic component and appropriate porosity. The construction of a multilayer anodes contributes to the diffusion and reforming reaction of hydrocarbon fuel.

(5) The comprehensive use of molecular dynamics, finite element simulation, supplementary mathematical models and other simulation methods is very important to understand the reaction process of hydrocarbon fuel on the catalyst and provide theoretical guides for the lifecycle assessment and large-scale fabrication of SOFCs. Hence, advanced in situ characterization methods integrated with theoretical calculation and other novel material design strategies such as single-atom design ideas are helpful to effectively optimize the composition and microstructure of carbon-resistant anodes.

(6) Combined with the practical application environment of hydrocarbon-based fuel cells, the research on alternative anodes should be optimized and integrated with the current advanced SOFC cathodes and electrolytes. Also, in the actual long-term operation of the cells, the atmosphere of the anodes will change from single CH₄ to CH₄, CO, CO₂, H₂, H₂O and other gas mixing systems, and hence it is more meaningful to study the effects of the fuel composition.

From the perspective of global energy strategy, it is imperative to vigorously promote the development of the natural gas industry. Presently, a well-developed methane-rich natural gas infrastructure is available, which can be readily and easily employed in residential or office buildings, significantly reducing the additional cost of using hydrocarbon-fuel SOFCs, accelerating the application of SOFCs in transportation and power distribution equipment. The above-mentioned research directions can promote the process of direct electrooxidation of hydrocarbon fuels to form CO₂ and H₂O or the process of indirect oxidation of syngas (CO + H₂) through reforming reaction to form CO₂ and H₂O, and play an essential role in improving the electrochemical performance, durability and cost-effectiveness of SOFCs. The optimized anode structures also make it easy to extend them to various cell structures, such as planar, tubular and honeycomb types, thereby speeding up the commercial application of hydrocarbon-fueled SOFCs.

To further facilitate the application of SOFCs, in addition to accelerating the development of multi-fuel compatible anodes, it is also necessary: (I) to develop efficient and green methods for the fabrication of SOFCs, e.g., 3D printing combined with theoretical model optimization. (II) Developing hybrid power systems based on SOFCs to improve their energy efficiency and economic viability, e.g., integrated coal gasification fuel cell combined cycle (IGFC), given that some of the heat energy of SOFCs can be recovered by introducing the reaction products into the steam turbine, where SOFC/gas turbine hybrid systems provide combined heat and power with efficiencies of up to 90%. (III) Reducing the operating temperature of SOFCs. Reducing the working temperature can improve the rapid start-up of the cell, alleviate the mismatch of the thermal expansion coefficient of different material systems and reduce the heat loss of the cell stack. Although the core influencing factor of the intermediate-to-low temperature SOFC is the ionic conductivity of the electrolyte, anode materials with high catalytic activity at intermediate-to-low temperature, matching electrolytes with high ionic conductivity and perovskite-based cathode materials with high activity and redox reaction stability will be a more suitable cell system. It is expected that hydrocarbon-fueled SOFCs will be widely used in stationary power generation and automobile applications.

Abbreviations

FCs	Fuel cells
PEMFCs	Proton-exchange membrane fuel cells
SOFCs	Solid oxide fuel cells
LSM	Sr-doped LaMnO3
LSC	La0.8Sr0.2CoO3
LSF	La0.8Sr0.2FeO3
LSCF	La0.8Sr0.2Co0.2Fe0.8O3
BSCF	Ba0.5Sr0.5Co0.8Fe0.2O3−δ
PBSCF	PrBa0.5Sr0.5Co1.5Fe0.5O5+δ
BPCF	Ba1−xPrxCo1−yFeyO3−δ
SCCO	Sr0.95Ce0.05CoO3−δ
BCFN	Ba0.9Co0.7Fe0.2Nb0.1O3−δ
YSZ	Yttria-stabilized zirconia
ScSZ	Scandia-stabilized zirconia
GDC	Gadolinium-doped ceria
SDC	Samaria- doped ceria
ASR	Area specific resistance
TEC	Thermal expansion coefficient
TPB	Three-phase boundary
MEIC	Mixed electronic and ionic conductivity
DIR	Direct internal reforming
HR-TEM	High-resolution transmission electron microscopy image
SEM	Scanning electron microscope
EDS	Energy dispersive X-ray spectroscopy
XRD	X-ray diffraction
SAED	Selected area electron diffraction
TPO	Temperature programmed oxidation
FTIR	Fourier transform infrared spectroscopy
MSR	Methane steam reforming
SVUV-PIMS	Synchrotron-based vacuum ultraviolet photoionization mass spectrometry
ETEM	Environmental transmission electron microscopy
DFT	Density functional theory
YST–YSZ	Sr0.88Y0.08TiO3−δ formed yttria-stabilized zirconia
LSCM	La0.75Sr0.25Cr0.5Mn0.5O3
LSTGM	La4.0Sr8.0Ti11.0Mn0.5Ga0.5O37.5
SMMO	Sr2MgMoO6−δ
OCV	Open-circuit voltage
ZDC	Zr0.1Ce0.9O2−δ
BZCYYb	BaZr0.1Ce0.7Y0.1Yb0.1O3−δ
PBMO	Pr0.5Ba0.5MnO3−δ
MPD	Maximum power density
DRM	Dry methane reforming
CZ	Ce0.5Zr0.5O2
CSZ	Ceria-stabilized zirconia
XPS	X-ray photoelectron spectroscopy
EDX	Energy dispersive X-ray spectroscopy
n-SNGG	Sn-doped Ni/gadolinium-doped ceria–gadolinium-doped ceria
HAADF STEM	High-angle annular dark-field scanning transmission electron microscopy
SR	Steam reforming
ATR	Auto-thermal reforming
STN	Sr0.94Ti0.9Nb0.1O3
UCG	Underground coal gasification
YDC	Yttria-doped ceria
CYYb	Ce0.8Y0.1Yb0.1O1.9
LSMO	La0.5Sr1.5MnO4±δ
STF	SrTi1−xFexO3
LSTN	La0.2Sr0.7Ti0.9Ni0.1O3−δ
LSC	Strontium-doped lanthanum chromate
LSCNi	Nickel-doped (La0.7Sr0.3)CrO3
L70SCrN	La0.70Sr0.3Cr0.85Ni0.15O3−δ
L65SCrN	La0.65Sr0.3Cr0.85Ni0.15O3−δ
STFN	Sr0.95(Ti0.3Fe0.63Ni0.07)O3
PBMCo	Pr0.5Ba0.5MnOx
LBM	La0.5Ba0.5MnO3−δ
LBMFC-1	La0.5Ba0.5Mn0.9Fe0.05Co0.05O3−δ
LBMFC–2	La0.5Ba0.5Mn0.8Fe0.1Co0.1O3−δ
PBMNx	(Pr0.5Ba0.5)1−x/2Mn1−x/2Nix/2O3−δ
A-LSC	A-site-deficient LaSrCrO3
LSCNi–Fe	(La0.7Sr0.3)(Cr0.85Ni0.12Fe0.03)O3−x
LSFN	Ni-doped La0.6Sr0.4FeO3−δ
L0.4STRN	(La0.4Sr0.4)(Ti0.85Ru0.07Ni0.08)O3−δ
L0.4STN	L0.4STN La0.4Sr0.4Ti0.85Ni0.15O3−δ
LSGM	Lanthanum gallate with strontium doping
LCeNT	La0.8Ce0.1Ni0.4Ti0.6O3−δ
GDC–PBCC	Gd0.2Ce0.8O1.9–PrBa0.8Ca0.2Co2O5+δ
LDC	La2Ce2O7
LSDC	La1.95Sm0.05Ce2O7
CNR	Ce0.90Ni0.05Ru0.05O2
CELD	Cathodic electrochemical deposition
SCE	Saturated calomel electrode
GC	Gas chromatography
NiMo–CZ	NiMo–Ce0.5Zr0.5O2−δ
PAW	Projector-augmented wave
CI-NEB	Climbing image-nudged elastic band
RMTN	Reverse methanation
ReaxFF-MD	Reactive force field molecular dynamics simulation
FIB-SEM	Focused ion beam-scanning electron microscope
AP-SOFC	All porous solid oxide fuel cell
PDOS	Projected density of states
IS	Initial state
TS	Transition state
FS	Final state
SIMFC	SIMulation of fuel cell
R-SZMO	H2-reduced Sr2ZnMoO6
FGA	Functional gradient anode
FAST	Field assisted sintering technique
SACs	Single-atom catalysts
AEMFC	Anion-exchange membrane fuel cells
CMSIs	Covalent metal support interactions
R-PSOCs	Reversible proton-conducting solid oxide cells

Latin symbols

I	Current
V	Voltage
p	Gas pressure
ΔG	Gibbs free energy variation
γ int	Interface energy
P	Cell power
δ	Oxygen vacancy
σ	Conductivity
e^−	Carrier concentration
μ	Mobility

Subscript

E thermo	Thermodynamic predicted voltage
η act	Activation loss
η ohmic	Ohmic loss
η conc	Concentration loss
Id	Intensity of D band
Ig	Intensity of G band
S BET	Surface area
gcat	The mass of catalyst
Oad	Adsorbed oxygen
OVs	Oxygen vacancies

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work was supported by 21C Innovation Laboratory, Contemporary Amperex Technology Ltd by project No. 21C-OP-202212, Foundation of Key Laboratory of Advanced Energy Materials Chemistry (Ministry of Education), Nankai University, the Foundation of State Key Laboratory of High-efficiency Utilization of Coal and Green Chemical Engineering (Grant No. 2022-K15), China University of Mining & Technology (Beijing), Beijing National Laboratory for Condensed Matter Physics, and the National Natural Science Foundation of China (No. 51172275, 51672029 and 51372271).

References

1. S. R. Bull, Renewable energy today and tomorrow, Proc. IEEE, 2001, 89, 1216–1226.
2. J. P. Barton and D. G. Infield, Energy storage and its use with tntermittent renewable energy, IEEE Trans. Energy Convers., 2004, 19, 441–448.
3. J. Qiao, Y. Liu, F. Hong and J. Zhang, A review of catalysts for the electroreduction of carbon dioxide to produce low-carbon fuels, Chem. Soc. Rev., 2014, 43, 631–675.
4. J. S. Sawyer, Man-made carbon dioxide and the “greenhouse” Effect, Nature, 1972, 239, 23–26.
5. B. C. H. Steele and A. Heinzel, Materials for fuel-cell technologies, Nature, 2001, 414, 345–352.
6. R. O'Hayre, S.-W. Cha, W. Colella and F. B. Prinz, Fuel Cell Fundamentals-Chapter 1: Introduction, Fuel Cell Fundamentals, Wiley Online Library, 2016, pp. 1–24.
7. A. L. Dicks and D. A. J. Rand., Fuel Cell Systems Explained-Introducing Fuel Cells, Fuel Cell Systems Explained, 2018, pp. 1–26.
8. M. A. Laguna-Bercero, Recent advances in high temperature electrolysis using solid oxide fuel cells: A review, J. Power Sources, 2012, 203, 4–16.
9. C. Sun and U. Stimming, Recent anode advances in solid oxide fuel cells, J. Power Sources, 2007, 171, 247–260.
10. A. J. Jacobson, Materials for solid oxide fuel cells, Chem. Mater., 2009, 22, 660–674.
11. M. A. Hassan, O. B. Mamat and M. Mehdi, Review: Influence of alloy addition and spinel coatings on Cr-based metallic interconnects of solid oxide fuel cells, Int. J. Hydrogen Energy, 2020, 45, 25191–25209.
12. M. Dewa, W. Yu, N. Dale, A. M. Hussain, M. G. Norton and S. Ha, Recent progress in integration of reforming catalyst on metal-supported SOFC for hydrocarbon and logistic fuels, Int. J. Hydrogen Energy, 2021, 46, 33523–33540.
13. H. A. Shabri, M. H. D. Othman, M. A. Mohamed, T. A. Kurniawan and S. M. Jamil, Recent progress in metal-ceramic anode of solid oxide fuel cell for direct hydrocarbon fuel utilization: A review, Fuel Process. Technol., 2021, 212, 106626.
14. S. Senthil Kumar and S. T. Aruna, Hydrocarbon compatible SOFC anode catalysts and their syntheses: A review, Sustainable Chem., 2021, 2, 707–763.
15. V. Marcantonio, L. Del Zotto, J. P. Ouweltjes and E. Bocci, Main issues of the impact of tar, H₂S, HCl and alkali metal from biomass-gasification derived syngas on the SOFC anode and the related gas cleaning technologies for feeding a SOFC system: A review, Int. J. Hydrogen Energy, 2022, 47, 517–539.
16. E. P. Murray, T. Tsai and S. A. Barnett, A direct-methane fuel cell with a ceria-based anode, Nature, 1999, 400, 649–651.
17. F. Yu, T. Han, Z. Wang, Y. Xie, Y. Wu, Y. Jin, N. Yang, J. Xiao and S. Kawi, Recent progress in direct carbon solid oxide fuel cell: Advanced anode catalysts, diversified carbon fuels, and heat management, Int. J. Hydrogen Energy, 2021, 46, 4283–4300.
18. P. Zhang, Z. Yang, Y. Jin, C. Liu, Z. Lei, F. Chen and S. Peng, Progress report on the catalyst layers for hydrocarbon-fueled SOFCs, Int. J. Hydrogen Energy, 2021, 46, 39369–39386.
19. L. Fan, C. e Li, L. van Biert, S.-H. Zhou, A. N. Tabish, A. Mokhov, P. V. Aravind and W. Cai, Advances on methane reforming in solid oxide fuel cells, Renewable Sustainable Energy Rev., 2022, 166, 112646.
20. W. Wang, C. Su, Y. Wu, R. Ran and Z. Shao, Progress in solid oxide fuel cells with nickel-based anodes operating on methane and related fuels, Chem. Rev., 2013, 113, 8104–8151.
21. S. S. Rathore, S. Biswas, D. Fini, A. P. Kulkarni and S. Giddey, Direct ammonia solid-oxide fuel cells: A review of progress and prospects, Int. J. Hydrogen Energy, 2021, 46, 35365–35384.
22. W. Winkler and H. Lorenz, The design of stationary and mobile solid oxide fuel cell-gas turbine systems, J. Power Sources, 2002, 105, 222–227.
23. B. C. H. Steele, Running on natural gas, Nature, 1999, 400, 619–621.
24. T. Takeguchi, Y. Kani, T. Yano, R. Kikuchi, K. Eguchi, K. Tsujimoto, Y. Uchida, A. Ueno, K. Omoshiki and M. Aizawa, Study on steam reforming of CH₄ and C₂ hydrocarbons and carbon deposition on Ni–YSZ cermets, J. Power Sources, 2002, 112, 588–595.
25. A. Atkinson, S. Barnett, R. J. Gorte, J. T. S. Irvine, A. J. McEvoy, M. Mogensen, S. C. Singhal and J. Vohs, Advanced anodes for high-temperature fuel cells, Nat. Mater., 2004, 3, 17–27.
26. M. Flytzani-Stephanopoulos, M. Sakbodin and Z. Wang, Regenerative adsorption and removal of H₂S from hot fuel gas streams by rare earth oxides, Science, 2006, 312, 1508–1510.
27. Z. Tao, M. Fu and Y. Liu, A mini-review of carbon-resistant anode materials for solid oxide fuel cells, Sustainable Energy Fuels, 2021, 5, 5420–5430.
28. S. Zarabi Golkhatmi, M. I. Asghar and P. D. Lund, A review on solid oxide fuel cell durability: Latest progress, mechanisms, and study tools, Renewable Sustainable Energy Rev., 2022, 161, 112339.
29. C. Sun and R. Hui, and J. Roller, Cathode materials for solid oxide fuel cells: a review, J. Solid State Electrochem., 2009, 14, 1125–1144.
30. W. S. Wang, Y. Y. Huang, S. W. Jung, J. M. Vohs and R. J. Gorte, A comparison of LSM, LSF, and LSCo for solid oxide electrolyzer anodes, J. Electrochem. Soc., 2006, 153, A2066–A2070.
31. M. A. Laguna-Bercero, J. A. Kilner and S. J. Skinner, Development of oxygen electrodes for reversible solid oxide fuel cells with scandia stabilized zirconia electrolytes, Solid State Ionics, 2011, 192, 501–504.
32. J. Suntivich, K. J. May, H. A. Gasteiger, J. B. Goodenough and Y. Shao-Horn, A perovskite oxide optimized for oxygen evolution catalysis from molecular orbital principles, Science, 2011, 334, 1383–1385.
33. W. Zhang, B. Yu and J. Xu, Investigation of single SOEC with BSCF anode and SDC barrier layer, Int. J. Hydrogen Energy, 2012, 37, 837–842.
34. L. Jiang, T. Wei, R. Zeng, W.-X. Zhang and Y.-H. Huang, Thermal and electrochemical properties of PrBa_0.5Sr_0.5Co_2−xFe_xO_5+δ (x = 0.5, 1.0, 1.5) cathode materials for solid-oxide fuel cells, J Power Sources, 2013, 232, 279–285.
35. R. Hui, C. Sun, S. Yick, C. Decès-Petit, X. Zhang, R. Maric and D. Ghosh, Ba_1−xPr_xCo_1−yFe_yO_3−δ as cathode materials for low temperature solid oxide fuel cells, Electrochim. Acta, 2010, 55, 4772–4775.
36. W. Yang, T. Hong, S. Li, Z. Ma, C. Sun, C. Xia and L. Chen, Perovskite Sr_1−xCe_xCoO_3−δ (0.05 ≤ x ≤ 0.15) as superior cathodes for intermediate temperature solid oxide fuel cells, ACS Appl. Mater. Interfaces, 2013, 5, 1143–1148.
37. Y. Gong, C. Sun, Q. A. Huang, J. A. Alonso, M. T. Fernandez-Diaz and L. Chen, Dynamic octahedral breathing in oxygen-deficient Ba_0.9Co_0.7Fe_0.2Nb_0.1O_3−δ Perovskite Performing as a Cathode in Intermediate-Temperature SOFC, Inorg. Chem., 2016, 55, 3091–3097.
38. D. J. Brett, A. Atkinson, N. P. Brandon and S. J. Skinner, Intermediate temperature solid oxide fuel cells, Chem. Soc. Rev., 2008, 37, 1568–1578.
39. H. Moon, S. Kim, S. Hyun and H. Kim, Development of IT-SOFC unit cells with anode-supported thin electrolytes via tape casting and co-firing, Int. J. Hydrogen Energy, 2008, 33, 1758–1768.
40. L. Zhang, S. P. Jiang, W. Wang and Y. Zhang, NiO/YSZ, anode-supported, thin-electrolyte, solid oxide fuel cells fabricated by gel casting, J. Power Sources, 2007, 170, 55–60.
41. X. Xu, C. Xia, S. Huang and D. Peng, YSZ thin films deposited by spin-coating for IT-SOFCs, Ceram. Int., 2005, 31, 1061–1064.
42. W. Zhou, H. Shi, R. Ran, R. Cai, Z. Shao and W. Jin, Fabrication of an anode-supported yttria-stabilized zirconia thin film for solid-oxide fuel cells via wet powder spraying, J. Power Sources, 2008, 184, 229–237.
43. J.-H. Kim, R.-H. Song, K.-S. Song, S.-H. Hyun, D.-R. Shin and H. Yokokawa, Fabrication and characteristics of anode-supported flat-tube solid oxide fuel cell, J. Power Sources, 2003, 122, 138–143.
44. Y. Du and N. M. Sammes, Fabrication and properties of anode-supported tubular solid oxide fuel cells, J. Power Sources, 2004, 136, 66–71.
45. S. D. Kim, S. H. Hyun, J. Moon, J.-H. Kim and R. H. Song, Fabrication and characterization of anode-supported electrolyte thin films for intermediate temperature solid oxide fuel cells, J. Power Sources, 2005, 139, 67–72.
46. Y. Yang, Y. Zhang and M. Yan, A review on the preparation of thin-film YSZ electrolyte of SOFCs by magnetron sputtering technology, Sep. Purif. Technol., 2022, 298, 121627.
47. X.-F. Ding, J.-Y. Shen, X.-J. Gao and J. Wang, Enhanced electrochemical properties of Sm_0.2Ce_0.8O_1.9 film for SOFC electrolyte fabricated by pulsed laser deposition, Rare Metals, 2014, 40, 1294–1299.
48. S. McIntosh and R. J. Gorte, Direct hydrocarbon solid oxide fuel cells, Chem. Rev., 2004, 104, 4845–4865.
49. I. Sreedhar, B. Agarwal, P. Goyal and S. A. Singh, Recent advances in material and performance aspects of solid oxide fuel cells, J. Electroanal. Chem., 2019, 848, 113315.
50. D. Simwonis, F. Tietz and D. Stover, Nickel coarsening in annealed Ni/8YSZ anode substrates for solid oxide fuel cells - In memoriam to Professor H. Tagawa, Solid State Ionics, 2000, 132, 241–251.
51. S. Tao and J. T. Irvine, A redox-stable efficient anode for solid-oxide fuel cells, Nat. Mater., 2003, 2, 320–323.
52. J. B. Young, Thermofluid modeling of fuel cells, Annu. Rev. Fluid Mech., 2007, 39, 193–215.
53. L. S. Skutina, A. I. Vylkov, I. N. Bainov, K. A. Chistyakov, D. K. Kuznetsov, O. B. Pavlenko and D. A. Medvedev, Catalytic properties of Sr₂Ni_0.75Mg_0.25MoO_6–δ based composites for application in hydrocarbon-fuelled solid oxide fuel cells, Int. J. Hydrogen Energy, 2021, 46, 16899–16906.
54. S. A. Saadabadi, B. Illathukandy and P. V. Aravind, Direct internal methane reforming in biogas fuelled solid oxide fuel cell; the influence of operating parameters, Energy Sci. Eng., 2021, 9, 1232–1248.
55. A. Abdulrasheed, A. A. Jalil, Y. Gambo, M. Ibrahim, H. U. Hambali and M. Y. Shahul, Hamid, A review on catalyst development for dry reforming of methane to syngas: Recent advances, Renewable Sustainable Energy Rev., 2019, 108, 175–193.
56. L. Cai, T. He, Y. Xiang and Y. Guan, Study on the reaction pathways of steam methane reforming for H₂ production, Energy, 2020, 207, 118296.
57. L. Bobrova, D. Andreev, E. Ivanov, N. Mezentseva, M. Simonov, L. Makarshin, A. Gribovskii and V. Sadykov, Water–gas shift reaction over Ni/CeO₂ Catalysts, Catalysts, 2017, 7, 310.
58. H. Sumi, Y.-H. Lee, H. Muroyama, T. Matsui, M. Kamijo, S. Mimuro, M. Yamanaka, Y. Nakajima and K. Eguchi, Effect of carbon deposition by carbon monoxide disproportionation on electrochemical characteristics at low temperature operation for solid oxide fuel cells, J. Power Sources, 2011, 196, 4451–4457.
59. A. L. Dicks, Advances in catalysts for internal reforming in high temperature fuel cells, J. Power Sources, 1998, 71, 111–122.
60. S. McIntosh, J. M. Vohs and R. J. Gorte, Role of hydrocarbon deposits in the enhanced performance of direct-oxidation SOFCs, J. Electrochem. Soc., 2003, 150, A470–A476.
61. M. D. Gross, J. M. Vohs and R. J. Gorte, Recent progress in SOFC anodes for direct utilization of hydrocarbons, J. Mater. Chem., 2007, 17, 3071–3077.
62. C. M. Chun and T. A. Ramanarayanan, Mechanism and control of carbon deposition on high temperature alloys, J. Electrochem. Soc., 2007, 154, C465–C471.
63. D. Lozano-Castello, D. Cazorla-Amoros, A. Linares-Solano, K. Oshida, T. Miyazaki, Y. J. Kim, T. Hayashi and M. Endo, Comparative characterization study of microporous carbons by HRTEM image analysis and gas adsorption, J. Phys. Chem. B, 2005, 109, 15032–15036.
64. T. Nakamura, N. Ohmura, K. Yahiro and K. Amezawa, Structural changes of Ni/YSZ cermet for solid oxide fuel cells in hydrocarbon gas containing atmospheres, ECS Trans., 2013, 57, 1539–1544.
65. H. Sumi, T. Yamaguchi, K. Hamamoto, T. Suzuki and Y. Fujishiro, Impact of direct butane microtubular solid oxide fuel cells, J. Power Sources, 2012, 220, 74–78.
66. D. Lee, J. Myung, J. Tan, S.-H. Hyun, J. T. S. Irvine, J. Kim and J. Moon, Direct methane solid oxide fuel cells based on catalytic partial oxidation enabling complete coking tolerance of Ni-based anodes, J. Power Sources, 2017, 345, 30–40.
67. T. Chen, W. G. Wang, H. Miao, T. Li and C. Xu, Evaluation of carbon deposition behavior on the nickel/yttrium-stabilized zirconia anode-supported fuel cell fueled with simulated syngas, J. Power Sources, 2011, 196, 2461–2468.
68. C. M. Chun and T. A. Ramanarayanan, Mechanism and control of carbon deposition on high temperature alloys, J. Electrochem. Soc., 2007, 154, 9.
69. X. Li, K. Blinn, Y. Fang, M. Liu, M. A. Mahmoud, S. Cheng, L. A. Bottomley, M. El-Sayed and M. Liu, Application of surface enhanced Raman spectroscopy to the study of SOFC electrode surfaces, Phys. Chem. Chem. Phys., 2012, 14, 5919–5923.
70. B. C. Bayer, D. A. Bosworth, F. B. Michaelis, R. Blume, G. Habler, R. Abart, R. S. Weatherup, P. R. Kidambi, J. J. Baumberg, A. Knop-Gericke, R. Schloegl, C. Baehtz, Z. H. Barber, J. C. Meyer and S. Hofmann, In situ observations of phase transitions in metastable nickel (carbide)/carbon nanocomposites, J. Phys. Chem. C, 2016, 120, 22571–22584.
71. B. Novosel, M. Marinsek and J. Macek, Deactivation of Ni–YSZ material in dry methane and oxidation of various forms of deposited carbon, J. Fuel Cell Sci. Technol., 2012, 9, 061003.
72. M. Rames, Y. Yu and G. Ren, Optimized negative staining: a high-throughput protocol for examining small and asymmetric protein structure by electron microscopy, J. Vis. Exp., 2014, 90, e51087.
73. H. Nishino, R. Nishida, T. Matsui, N. Kawase and I. Mochida, Growth of amorphous carbon nanotube from poly(tetrafluoroethylene) and ferrous chloride, Carbon, 2003, 41, 2819–2823.
74. Y. Chen, Y. Zhang, Y. Lin, Z. Yang, D. Su, M. Han and F. Chen, Direct-methane solid oxide fuel cells with hierarchically porous Ni-based anode deposited with nanocatalyst layer, Nano Energy, 2014, 10, 1–9.
75. A. Cuesta, P. Dhamelincourt, J. Laureyns, A. Martínez-Alonso and J. M. D. Tascón, Raman microprobe studies on carbon materials, Carbon, 1994, 32, 1523–1532.
76. H. Sumi, H. Shimada, T. Yamaguchi, K. Hamamoto, T. Suzuki and Y. Fujishiro, Development of microtubular solid oxide fuel cells using hydrocarbon fuels, Advances in Solid Oxide Fuel Cells and Electronic Ceramics, 2015, pp. 93–104.
77. M. Chlipała, P. Błaszczak, S. F. Wang, P. Jasiński and B. Bochentyn, In situ study of a composition of outlet gases from biogas fuelled solid oxide fuel cell performed by the Fourier Transform Infrared Spectroscopy, Int. J. Hydrogen Energy, 2019, 44, 13864–13874.
78. Y. Xie, N. Shi, X. Hu, M. Liu, Y. Yang, D. Huan, Y. Pan, R. Peng and C. Xia, Novel in-situ MgO nano-layer decorated carbon-tolerant anode for solid oxide fuel cells, Int. J. Hydrogen Energy, 2020, 45, 11791–11801.
79. C. Sun, R. Su, J. Chen, L. Lu and P. Guan, Carbon formation mechanism of C₂H₂ in Ni-based catalysts revealed by in situ electron microscopy and molecular dynamics simulations, ACS Omega, 2019, 4, 8413–8420.
80. B. Steele, Appraisal of Ce_1−yGd_yO_2−y/2 electrolytes for IT-SOFC operation at 500 °C, Solid State Ionics, 2000, 129, 95–110.
81. K. Sasaki and Y. Teraoka, Equilibria in fuel cell gases - II. The C-H-O ternary diagrams, J. Electrochem. Soc., 2003, 150, 885–888.
82. N. Galea, D. Knapp and T. Ziegler, Density functional theory studies of methane dissociation on anode catalysts in solid-oxide fuel cells: Suggestions for coke reduction, J. Catal., 2007, 247, 20–33.
83. E. Fabbri, D. Pergolesi and E. Traversa, Materials challenges toward proton-conducting oxide fuel cells: a critical review, Chem. Soc. Rev., 2010, 39, 4355–4369.
84. H. Zhu, R. J. Kee, V. M. Janardhanan, O. Deutschmann and D. G. Goodwin, Modeling elementary heterogeneous chemistry and electrochemistry in Solid-Oxide Fuel Cells, J. Electrochem. Soc., 2005, 152, 12.
85. S. Ahmed, A. Aitani, F. Rahman, A. Al-Dawood and F. Al-Muhaish, Decomposition of hydrocarbons to hydrogen and carbon, Appl. Catal., A, 2009, 359, 1–24.
86. B. C. H. Steele, I. Kelly, H. Middleton and R. Rudkin, Oxidation of methane in solid state electrochemical reactors, Solid State Ionics, 1988, 28–30, 1547–1552.
87. H. He, Characterization of YSZ-YST composites for SOFC anodes, Solid State Ionics, 2004, 175, 171–176.
88. O. A. Marina, N. L. Canfield and J. W. Stevenson, Thermal, electrical, and electrocatalytical properties of lanthanum-doped strontium titanate, Solid State Ionics, 2002, 149, 21–28.
89. J. C. Ruiz-Morales, J. Canales-Vazquez, C. Savaniu, D. Marrero-Lopez, W. Zhou and J. T. Irvine, Disruption of extended defects in solid oxide fuel cell anodes for methane oxidation, Nature, 2006, 439, 568–571.
90. Y. H. Huang, R. I. Dass, Z. L. Xing and J. B. Goodenough, Double perovskites as anode materials for solid-oxide fuel cells, Science, 2006, 312, 254–257.
91. M. L. Toebes, J. H. Bitter, A. J. van Dillen and K. P. de Jong, Impact of the structure and reactivity of nickel particles on the catalytic growth of carbon nanofibers, Catal. Today, 2002, 76, 33–42.
92. R. T. K. Baker and J. J. Chludzinski, Jr., In-situ electron microscopy studies of the behavior of supported ruthenium particles. 2. Carbon deposition from catalyzed decomposition of acetylene, J. Phys. Chem., 1986, 90, 4734–4738.
93. R. Shiozaki, A. G. Andersen, T. Hayakawa, S. Hamakawa, K. Suzuki, M. Shimizu and K. Takehira, Partial oxidation of methane over a Ni/BaTiO₃ catalyst prepared by solid phase crystallization, J. Chem. Soc., Faraday Trans., 1997, 93, 3235–3242.
94. H. Kim, C. Lu, W. L. Worrell, J. M. Vohs and R. J. Gorte, Cu-ni cermet anodes for direct oxidation of methane in Solid-Oxide Fuel Cells, J. Electrochem. Soc., 2002, 149, 3.
95. Z. Lü, L. Pei, T.-M. He, X.-Q. Huang, Z.-G. Liu, Y. Ji, X.-H. Zhao and W.-H. Su, Study on new copper-containing SOFC anode materials, J. Alloys Compd., 2002, 334, 299–303.
96. S. Song, M. Han, J. Zhang and H. Fan, NiCu–Zr_0.1Ce_0.9O_2−δ anode materials for intermediate temperature solid oxide fuel cells using hydrocarbon fuels, J. Power Sources, 2013, 233, 62–68.
97. Y. Akdeniz, B. Timurkutluk and C. Timurkutluk, Development of anodes for direct oxidation of methane fuel in solid oxide fuel cells, Int. J. Hydrogen Energy, 2016, 41, 10021–10029.
98. M. Li, B. Hua, J. Pu, B. Chi and L. Jian, Electrochemical performance and carbon deposition resistance of M-BaZr_0.1Ce_0.7Y_0.1Yb_0.1O_3−δ (M=Pd, Cu, Ni or NiCu) anodes for solid oxide fuel cells, Sci. Rep., 2015, 5, 7667.
99. S. Kim, C. Kim, J. H. Lee, J. Shin, T.-H. Lim and G. Kim, Tailoring Ni-based catalyst by alloying with transition metals (M = Ni, Co, Cu, and Fe) for direct hydrocarbon utilization of energy conversion devices, Electrochim. Acta, 2017, 225, 399–406.
100. S. Islam and J. M. Hill, Preparation of Cu–Ni/YSZ solid oxide fuel cell anodes using microwave irradiation, J. Power Sources, 2011, 196, 5091–5094.
101. E. W. Park, H. Moon, M.-S. Park and S. H. Hyun, Fabrication and characterization of Cu–Ni–YSZ SOFC anodes for direct use of methane via Cu-electroplating, Int. J. Hydrogen Energy, 2009, 34, 5537–5545.
102. S. Senthil Kumar, V. Jayaram and S. T. Aruna, Co-fired anode-supported solid oxide fuel cell for internal reforming of hydrocarbon fuel, Energy, Ecol. Environ., 2020, 6, 55–68.
103. A. Ringuedé, D. Bronine and J. R. Frade, Ni_1−xCo_x/YSZ cermet anodes for solid oxide fuel cells, Electrochim. Acta, 2002, 48, 437–442.
104. C.-K. Cho, B.-H. Choi and K.-T. Lee, Effect of Co alloying on the electrochemical performance of Ni–Ce_0.8Gd_0.2O_1.9 anodes for hydrocarbon-fueled solid oxide fuel cells, J. Alloys Compd., 2012, 541, 433–439.
105. W. Nicharee, S. Chaianansutcharit and K. Sato, Electrochemical performance and stability of Ni_1−xCo_x-based cermet anode for direct methane-fuelled solid oxide fuel cells, MATEC Web Conf., 2017, 130, 03005.
106. M. R. Cesario, G. S. Souza, F. J. A. Loureiro, A. J. M. Araújo, J. P. F. Grilo, S. Aouad, H. L. Tidahy, D. A. Macedo, D. P. Fagg, C. Gennequin and E. Abi-Aad, Synthesis of Co–Ni and Cu–Ni based-catalysts for dry reforming of methane as potential components for SOFC anodes, Ceram. Int., 2021, 47, 33191–33201.
107. H. Kan and H. Lee, Enhanced stability of Ni–Fe/GDC solid oxide fuel cell anodes for dry methane fuel, Catal. Commun., 2010, 12, 36–39.
108. S. H. Lee and H. Kim, Dual layered anode support for the internal reforming of methane for solid oxide fuel cells, Ceram. Int., 2014, 40, 5959–5966.
109. Y. M. Park and H. Kim, An additional layer in an anode support for internal reforming of methane for solid oxide fuel cells, Int. J. Hydrogen Energy, 2014, 39, 16513–16523.
110. S. Lee, J. H. Park, K. T. Lee and Y.-W. Ju, Anodic properties of Ni-Fe bimetallic nanofiber for solid oxide fuel cell using LaGaO₃ electrolyte, J. Alloys Compd., 2021, 875, 159911.
111. B. Hua, M. Li, Y.-Q. Zhang, J. Chen, Y.-F. Sun, N. Yan, J. Li and J.-L. Luo, Facile synthesis of highly active and robust Ni–Mo bimetallic electrocatalyst for hydrocarbon oxidation in Solid Oxide Fuel Cells, ACS Energy Lett., 2016, 1, 225–230.
112. Q. Bkour, F. Che, K.-M. Lee, C. Zhou, N. Akter, J. A. Boscoboinik, K. Zhao, J. T. Gray, S. R. Saunders, M. Grant Norton, J.-S. McEwen, T. Kim and S. Ha, Enhancing the partial oxidation of gasoline with Mo-doped Ni catalysts for SOFC applications: An integrated experimental and DFT study, Appl. Catal., B, 2020, 266, 118626.
113. Q. Bkour, K. Zhao, L. Scudiero, D. J. Han, C. W. Yoon, O. G. Marin-Flores, M. G. Norton and S. Ha, Synthesis and performance of ceria-zirconia supported Ni–Mo nanoparticles for partial oxidation of isooctane, Appl. Catal., B, 2017, 212, 97–105.
114. X. Hou, K. Zhao, O. A. Marina, M. Grant Norton and S. Ha, NiMo-ceria-zirconia-based anode for solid oxide fuel cells operating on gasoline surrogate, Appl. Catal., B, 2019, 242, 31–39.
115. A. J. Majewski, S. K. Singh, N. K. Labhasetwar and R. Steinberger-Wilckens, Nickel–molybdenum catalysts for combined solid oxide fuel cell internal steam and dry reforming, Chem. Eng. Sci., 2021, 232, 116341.
116. E. Nikolla, J. Schwank and S. Linic, Promotion of the long-term stability of reforming Ni catalysts by surface alloying, J. Catal., 2007, 250, 85–93.
117. E. Nikolla, J. Schwank and S. Linic, Comparative study of the kinetics of methane steam reforming on supported Ni and Sn/Ni alloy catalysts: The impact of the formation of Ni alloy on chemistry, J. Catal., 2009, 263, 220–227.
118. E. Nikolla, J. Schwank and S. Linic, Direct electrochemical oxidation of hydrocarbon fuels on SOFCs: improved carbon tolerance of Ni alloy anodes, J. Electrochem. Soc., 2009, 156, 11.
119. E. Nikolla, J. Schwank and S. Linic, Hydrocarbon steam reforming on Ni alloys at solid oxide fuel cell operating conditions, Catal. Today, 2008, 136, 243–248.
120. H. Kan, S. H. Hyun, Y.-G. Shul and H. Lee, Improved solid oxide fuel cell anodes for the direct utilization of methane using Sn-doped Ni/YSZ catalysts, Catal. Commun., 2009, 11, 180–183.
121. H. Kan and H. Lee, Sn-doped Ni/YSZ anode catalysts with enhanced carbon deposition resistance for an intermediate temperature SOFC, Appl. Catal., B, 2010, 97, 108–114.
122. Z. Jiang, N. A. Arifin, P. Mardle and R. Steinberger-Wilckens, Electrochemical performance and carbon resistance comparison between tin, copper and silver-doped nickel/yttria-stabilized zirconia anodes SOFCs operated with biogas, J. Electrochem. Soc., 2019, 166, F393–F398.
123. A. Singh and J. M. Hill, Evaluation of Sn-modified Ni/YSZ SOFC anodes for the direct utilization of methane, ECS Trans., 2011, 35, 1397–1406.
124. B. Farrell and S. Linic, Direct electrochemical oxidation of ethanol on SOFCs: Improved carbon tolerance of Ni anode by alloying, Appl. Catal., B, 2016, 183, 386–393.
125. D. Yoon and A. Manthiram, Hydrocarbon-fueled solid oxide fuel cells with surface-modified, hydroxylated Sn/Ni–Ce_0.8Gd_0.2O_1.9 heterogeneous catalyst anode, J. Mater. Chem. A, 2014, 2, 17041–17046.
126. J.-h Myung, S.-D. Kim, T. H. Shin, D. Lee, J. T. S. Irvine, J. Moon and S.-H. Hyun, Nano-composite structural Ni–Sn alloy anodes for high performance and durability of direct methane-fueled SOFCs, J. Mater. Chem. A, 2015, 3, 13801–13806.
127. Q. Yang, J. Chen, C. Sun and L. Chen, Direct operation of methane fueled solid oxide fuel cells with Ni cermet anode via Sn modification, Int. J. Hydrogen Energy, 2016, 41, 11391–11398.
128. P. Li, Z. Wang, X. Yao, N. Hou, L. Fan, T. Gan, Y. Zhao, Y. Li and J. W. Schwank, Effect of Sn addition on improving the stability of Ni–Ce_0.8Sm_0.2O_1.9 anode material for solid oxide fuel cells fed with dry CH4, Catal. Today, 2019, 330, 209–216.
129. N. A. Arifin, A. H. Shamsuddin and R. Steinberger-Wilckens, Evaluating the drop of electrochemical performance of Ni/YSZ and Ni/ScSZ solid oxide fuel cells operated with dry biogas, Int. J. Energy Res., 2020, 45, 6405–6417.
130. A. Singh and J. M. Hill, Carbon tolerance, electrochemical performance and stability of solid oxide fuel cells with Ni/yttria stabilized zirconia anodes impregnated with Sn and operated with methane, J. Power Sources, 2012, 214, 185–194.
131. H. Jang, J. Eom, H. Ju and J. Lee, Ameliorated performance in a direct carbon fuel cell using Sn mediator on Ni–YSZ anode surface, Catal. Today, 2016, 260, 158–164.
132. A. Cantos-Gómez, R. Ruiz-Bustos and J. van Duijn, Ag as an alternative for Ni in direct hydrocarbon SOFC anodes, Fuel Cells, 2011, 11, 140–143.
133. X. Wu, X. Zhou, Y. Tian, X. Kong, J. Zhang, W. Zuo, X. Ye and K. Sun, Preparation and electrochemical performance of silver impregnated Ni–YSZ anode for solid oxide fuel cell in dry methane, Int. J. Hydrogen Energy, 2015, 40, 16484–16493.
134. X. Wu, Y. Tian, J. Zhang, W. Zuo, X. Kong, J. Wang, K. Sun and X. Zhou, Enhanced electrochemical performance and carbon anti-coking ability of solid oxide fuel cells with silver modified nickel-yttrium stabilized zirconia anode by electroless plating, J. Power Sources, 2016, 301, 143–150.
135. T. Hibino, A. Hashimoto, M. Yano, M. Suzuki, S.-I. Yoshida and M. Sano, High performance anodes for SOFCs operating in methane-air mixture at reduced temperatures, J. Electrochem. Soc., 2001, 149, A133–A136.
136. A. Babaei, L. Zhang, E. Liu and S. P. Jiang, Performance and carbon deposition over Pd nanoparticle catalyst promoted Ni/GDC anode of SOFCs in methane, methanol and ethanol fuels, Int. J. Hydrogen Energy, 2012, 37, 15301–15310.
137. V. Modafferi, G. Panzera, V. Baglio, F. Frusteri and P. L. Antonucci, Propane reforming on Ni–Ru/GDC catalyst: H₂ production for IT-SOFCs under SR and ATR conditions, Appl. Catal., A, 2008, 334, 1–9.
138. M. Boaro, V. Modafferi, A. Pappacena, J. Llorca, V. Baglio, F. Frusteri, P. Frontera, A. Trovarelli and P. L. Antonucci, Comparison between Ni–Rh/gadolinia doped ceria catalysts in reforming of propane for anode implementations in intermediate solid oxide fuel cells, J. Power Sources, 2010, 195, 649–661.
139. T. Hibino, A. Hashimoto, M. Yano, M. Suzuki and M. Sano, Ru-catalyzed anode materials for direct hydrocarbon SOFCs, Electrochim. Acta, 2003, 48, 2531–2537.
140. A. M. Hussain, J. V. T. Høgh, W. Zhang and N. Bonanos, Efficient ceramic anodes infiltrated with binary and ternary electrocatalysts for SOFCs operating at low temperatures, J. Power Sources, 2012, 216, 308–313.
141. L. Liu, Q. Yang, W. Yang, X. Qi, C. Sun and L. Chen, Li/Na modified Ni-SDC anode for methane-fueled Solid Oxide Fuel Cells, ECS Trans., 2015, 68, 1403–1409.
142. M. Asamoto, S. Miyake, K. Sugihara and H. Yahiro, Improvement of Ni/SDC anode by alkaline earth metal oxide addition for direct methane–solid oxide fuel cells, Electrochem. Commun., 2009, 11, 1508–1511.
143. D. La Rosa, A. Sin, M. L. Faro, G. Monforte, V. Antonucci and A. S. Aricò, Mitigation of carbon deposits formation in intermediate temperature solid oxide fuel cells fed with dry methane by anode doping with barium, J. Power Sources, 2009, 193, 160–164.
144. S. Islam and J. M. Hill, Barium oxide promoted Ni/YSZ solid-oxide fuel cells for direct utilization of methane, J. Mater. Chem. A, 2014, 2, 1922–1929.
145. M. D. McIntyre, J. D. Kirtley, A. Singh, S. Islam, J. M. Hill and R. A. Walker, Comparing in situ carbon tolerances of Sn-infiltrated and BaO-infiltrated Ni–YSZ cermet anodes in Solid Oxide Fuel Cells exposed to methane, J. Phys. Chem. C, 2015, 119, 7637–7647.
146. Y. Itagaki, S. Yamaguchi and H. Yahiro, Anodic performance of BaO-added Ni/SDC for solid oxide fuel cell fed with dry CH₄, Front. Energy Res., 2021, 9, 652239.
147. Y. H. Hu, Solid-solution catalysts for CO₂ reforming of methane, Catal. Today, 2009, 148, 206–211.
148. R. Zanganeh, M. Rezaei and A. Zamaniyan, Dry reforming of methane to synthesis gas on NiO–MgO nanocrystalline solid solution catalysts, Int. J. Hydrogen Energy, 2013, 38, 3012–3018.
149. Q. Yang, F. Chai, C. Ma, C. Sun, S. Shi and L. Chen, Enhanced coking tolerance of a MgO-modified Ni cermet anode for hydrocarbon fueled solid oxide fuel cells, J. Mater. Chem. A, 2016, 4, 18031–18036.
150. S.-I. Lee, K. Ahn, J. M. Vohs and R. J. Gorte, Cu-Co bimetallic anodes for direct utilization of methane in SOFCs, Electrochem. Solid-State Lett., 2005, 8, 1.
151. M. D. Gross, J. M. Vohs and R. J. Gorte, A study of thermal stability and methane tolerance of Cu-based SOFC anodes with electrodeposited Co, Electrochim. Acta, 2007, 52, 1951–1957.
152. A. Fuerte, R. X. Valenzuela, M. J. Escudero and L. Daza, Effect of cobalt incorporation in copper-ceria based anodes for hydrocarbon utilisation in Intermediate Temperature Solid Oxide Fuel Cells, J. Power Sources, 2011, 196, 4324–4331.
153. N. Yan, J. Pandey, Y. Zeng, B. S. Amirkhiz, B. Hua, N. J. Geels, J.-L. Luo and G. Rothenberg, Developing a thermal- and coking-resistant cobalt–tungsten bimetallic anode catalyst for Solid Oxide Fuel Cells, ACS Catal., 2016, 6, 4630–4634.
154. G. Kaur and S. Basu, Performance studies of copper–iron/ceria–yttria stabilized zirconia anode for electro-oxidation of butane in solid oxide fuel cells, J. Power Sources, 2013, 241, 783–790.
155. G. Kaur and S. Basu, Physical characterization and electrochemical performance of copper-iron-ceria-YSZ anode-based SOFCs in H₂ and methane fuels, Int. J. Energy Res., 2015, 39, 1345–1354.
156. M. Hasegawa, Ellingham Diagram, Treatise Process Metallurgy, 2014, pp. 507–516.
157. C. Sun, H. Li and L. Chen, Nanostructured ceria-based materials: synthesis, properties, and applications, Energy Environ. Sci., 2012, 5, 8475–8505.
158. N. V. Skorodumova, S. I. Simak, B. I. Lundqvist, I. A. Abrikosov and B. Johansson, Quantum origin of the oxygen storage capability of ceria, Phys. Rev. Lett., 2002, 89, 166601.
159. C. Sun, J. Sun, G. Xiao, H. Zhang, X. Qiu, H. Li and L. Chen, Mesoscale organization of nearly monodisperse flowerlike ceria microspheres, J. Phys. Chem. B, 2006, 110, 13445–13452.
160. J. W. Yun, S. P. Yoon, H. S. Kim, J. Han and S. W. Nam, Effect of Sm_0.2Ce_0.8O_1.9 on the carbon coking in Ni-based anodes for solid oxide fuel cells running on methane fuel, Int. J. Hydrogen Energy, 2012, 37, 4356–4366.
161. E. S. Putna, J. Stubenrauch, J. M. Vohs and R. J. Gorte, Ceria-based anodes for the direct oxidation of methane in Solid Oxide Fuel Cells, Langmuir, 1995, 11, 4832–4837.
162. S. Park, R. Craciun, J. M. Vohs and R. J. Gorte, Direct oxidation of hydrocarbons in a Solid Oxide Fuel Cell: I. methane oxidation, J. Electrochem. Soc., 1999, 146, 3603–3605.
163. C. Lu, W. L. Worrell, J. M. Vohs and R. J. Gorte, A comparison of Cu-ceria-SDC and Au-ceria-SDC composites for SOFC anodes, J. Electrochem. Soc., 2003, 150, A1357–A1359.
164. O. Marina, A solid oxide fuel cell with a gadolinia-doped ceria anode: preparation and performance, Solid State Ionics, 1999, 123, 199–208.
165. M. Wisniewski, A. Boréave and P. Gélin, Catalytic CO₂ reforming of methane over Ir/Ce_0.9Gd_0.1O_2−x, Catal. Commun., 2005, 6, 596–600.
166. M.-F. Han, J. Xiong, T. Zhu and S. C. Singhal, Performance of Solid Oxide Fuel Cells on H₂, NH₃ and hydrocarbon fuels, ECS Trans., 2013, 50, 163.
167. K. Ahn, H. He, J. M. Vohs and R. J. Gorte, Enhanced thermal stability of SOFC anodes made with CeO₂-ZrO₂ Solutions, Electrochem. Solid-State Lett., 2005, 8, A414–A417.
168. X. Zhou, J. Zhen, L. Liu, X. Li, N. Zhang and K. Sun, Enhanced sulfur and carbon coking tolerance of novel co-doped ceria based anode for solid oxide fuel cells, J. Power Sources, 2012, 201, 128–135.
169. D. Neagu and J. T. S. Irvine, Enhancing electronic conductivity in strontium titanates through correlated A and B-Site doping, Chem. Mater., 2011, 23, 1607–1617.
170. C. Zhao, Y. Li, W. Zhang, Y. Zheng, X. Lou, B. Yu, J. Chen, Y. Chen, M. Liu and J. Wang, Heterointerface engineering for enhancing the electrochemical performance of solid oxide cells, Energy Environ. Sci., 2020, 13, 53–85.
171. Y. Li, J. Zhang, Q. Chen, X. Xia and M. Chen, Emerging of heterostructure materials in energy storage: A review, Adv. Mater., 2021, 33, e2100855.
172. D. Neagu, G. Tsekouras, D. N. Miller, H. Menard and J. T. Irvine, In situ growth of nanoparticles through control of non-stoichiometry, Nat. Chem., 2013, 5, 916–923.
173. S. He, M. Li, J. Hui and X. Yue, In-situ construction of ceria-metal/titanate heterostructure with controllable architectures for efficient fuel electrochemical conversion, Appl. Catal., B, 2021, 298, 120588.
174. S. Vecino-Mantilla, P. Simon, M. Huvé, G. Gauthier and P. Gauthier-Maradei, Methane steam reforming in water-deficient conditions on a new Ni-exsolved Ruddlesden-Popper manganite: Coke formation and H₂S poisoning, Int. J. Hydrogen Energy, 2020, 45, 27145–27159.
175. Y. Li, W. Zhang, Y. Zheng, J. Chen, B. Yu, Y. Chen and M. Liu, Controlling cation segregation in perovskite-based electrodes for high electro-catalytic activity and durability, Chem. Soc. Rev., 2017, 46, 6345–6378.
176. Y. Sun, J. Li, Y. Zeng, B. S. Amirkhiz, M. Wang, Y. Behnamian and J. Luo, A-site deficient perovskite: the parent for in situ exsolution of highly active, regenerable nano-particles as SOFC anodes, J. Mater. Chem. A, 2015, 3, 11048–11056.
177. A. Marcucci, F. Zurlo, I. N. Sora, E. Placidi, S. Casciardi, S. Licoccia and E. Di Bartolomeo, A redox stable Pd-doped perovskite for SOFC applications, J. Mater. Chem. A, 2019, 7, 5344–5352.
178. N. K. Monteiro, F. B. Noronha, L. O. O. da Costa, M. Linardi and F. C. Fonseca, A direct ethanol anode for solid oxide fuel cell based on a chromite-manganite with catalytic ruthenium nanoparticles, Int. J. Hydrogen Energy, 2012, 37, 9816–9829.
179. H. Tanaka, M. Uenishi, M. Taniguchi, I. Tan, K. Narita, M. Kimura, K. Kaneko, Y. Nishihata and J. I. Mizuki, The intelligent catalyst having the self-regenerative function of Pd, Rh and Pt for automotive emissions control, Catal. Today, 2006, 117, 321–328.
180. M. van den Bossche and S. McIntosh, Pulse reactor studies to assess the potential of La_0.75Sr_0.25Cr_0.5Mn_0.4X_0.1O_3−δ(X = Co, Fe, Mn, Ni, V) as direct hydrocarbon Solid Oxide Fuel Cell anodes, Chem. Mater., 2010, 22, 5856–5865.
181. B. H. Park and G. M. Choi, Ex-solution of Ni nanoparticles in a La_0.2Sr_0.8Ti_1−xNi_xO_3−δ alternative anode for solid oxide fuel cell, Solid State Ionics, 2014, 262, 345–348.
182. M. Niania, R. Podor, T. B. Britton, C. Li, S. J. Cooper, N. Svetkov, S. Skinner and J. Kilner, In situ study of strontium segregation in La_0.6Sr_0.4Co_0.2Fe_0.8O_3−δ in ambient atmospheres using high-temperature environmental scanning electron microscopy, J. Mater. Chem. A, 2018, 6, 14120–14135.
183. W. Lee, J. W. Han, Y. Chen, Z. Cai and B. Yildiz, Cation size mismatch and charge interactions drive dopant segregation at the surfaces of manganite perovskites, J. Am. Chem. Soc., 2013, 135, 7909–7925.
184. B. Koo, K. Kim, J. K. Kim, H. Kwon, J. W. Han and W. Jung, Sr segregation in perovskite oxides: why it happens and how it exists, Joule, 2018, 2, 1476–1499.
185. Y. Chen, W. Jung, Z. Cai, J. J. Kim, H. L. Tuller and B. Yildiz, Impact of Sr segregation on the electronic structure and oxygen reduction activity of SrTi_1−xFe_xO₃ surfaces, Energy Environ. Sci., 2012, 5, 7979–7988.
186. E. N. Sgourou, Y. Panayiotatos, K. Davazoglou, A. L. Solovjov, R. V. Vovk and A. Chroneos, Self-diffusion in perovskite and perovskite related oxides: insights from modelling, Appl. Sci., 2020, 10, 2286.
187. D. Neagu, V. Kyriakou, I. L. Roiban, M. Aouine, C. Tang, A. Caravaca, K. Kousi, I. Schreur-Piet, I. S. Metcalfe, P. Vernoux, M. C. M. van de Sanden and M. N. Tsampas, In situ observation of nanoparticle exsolution from perovskite oxides: from atomic scale mechanistic insight to nanostructure tailoring, ACS Nano, 2019, 13, 12996–13005.
188. Y. Gao, D. Chen, M. Saccoccio, Z. Lu and F. Ciucci, From material design to mechanism study: Nanoscale Ni exsolution on a highly active A-site deficient anode material for solid oxide fuel cells, Nano Energy, 2016, 27, 499–508.
189. T. S. Oh, E. K. Rahani, D. Neagu, J. T. Irvine, V. B. Shenoy, R. J. Gorte and J. M. Vohs, Evidence and model for strain-driven release of metal nanocatalysts from perovskites during exsolution, J. Phys. Chem. Lett., 2015, 6, 5106–5110.
190. D. Neagu, T. S. Oh, D. N. Miller, H. Menard, S. M. Bukhari, S. R. Gamble, R. J. Gorte, J. M. Vohs and J. T. S. Irvine, Nano-socketed nickel particles with enhanced coking resistance grown in situ by redox exsolution, Nat. Commun., 2015, 6, 8120.
191. K. J. Kim, H. Han, T. Defferriere, D. Yoon, S. Na, S. J. Kim, A. M. Dayaghi, J. Son, T. S. Oh, H. M. Jang and G. M. Choi, Facet-dependent in situ growth of nanoparticles in epitaxial thin films: the role of interfacial energy, J. Am. Chem. Soc., 2019, 141, 7509–7517.
192. N. Hou, T. Yao, P. Li, X. Yao, T. Gan, L. Fan, J. Wang, X. Zhi, Y. Zhao and Y. Li, A-site ordered double perovskite with in situ exsolved core-shell nanoparticles as anode for Solid Oxide Fuel Cells, ACS Appl. Mater. Interfaces, 2019, 11, 6995–7005.
193. Y. Nishihata, J. Mizuki, T. Akao, H. Tanaka, M. Uenishi, M. Kimura, T. Okamoto and N. Hamada, Self-regeneration of a Pd-perovskite catalyst for automotive emissions control, Nature, 2002, 418, 164–167.
194. T. Cao, O. Kwon, R. J. Gorte and J. M. Vohs, Metal exsolution to enhance the catalytic activity of electrodes in Solid Oxide Fuel Cells, Nanomaterials, 2020, 10, 2445.
195. Z. Li, M. Li and Z. Zhu, Perovskite cathode materials for low-temperature Solid Oxide Fuel Cells: fundamentals to optimization, Electrochem. Energy Rev., 2021, 5, 263–311.
196. Z. Du, H. Zhao, S. Yi, Q. Xia, Y. Gong, Y. Zhang, X. Cheng, Y. Li, L. Gu and K. Świerczek, High-performance anode material Sr₂FeMo_0.65Ni_0.35O_6−δ with in situ exsolved nanoparticle catalyst, ACS Nano, 2016, 10, 8660–8669.
197. D.-M. Amaya-Dueñas, G. Chen, A. Weidenkaff, N. Sata, F. Han, I. Biswas, R. Costa and K. A. Friedrich, A-site deficient chromite with in situ Ni exsolution as a fuel electrode for solid oxide cells (SOCs), J. Mater. Chem. A, 2021, 9, 5685–5701.
198. T. Zhu, H. E. Troiani, L. V. Mogni, M. Han and S. A. Barnett, Ni-substituted Sr(Ti,Fe)O₃ SOFC anodes: achieving high performance via metal alloy nanoparticle exsolution, Joule, 2018, 2, 478–496.
199. W. Kobsiriphat, B. D. Madsen, Y. Wang, L. D. Marks and S. A. Barnett, La_0.8Sr_0.2Cr_1−xRu_xO_3−δ–Gd_0.1Ce_0.9O_1.95 solid oxide fuel cell anodes: Ru precipitation and electrochemical performance, Solid State Ionics, 2009, 180, 257–264.
200. Y. F. Sun, Y. Q. Zhang, J. Chen, J. H. Li, Y. T. Zhu, Y. M. Zeng, B. S. Amirkhiz, J. Li, B. Hua and J. L. Luo, New opportunity for in situ exsolution of metallic nanoparticles on perovskite parent, Nano Lett., 2016, 16, 5303–5309.
201. B. D. Madsen, W. Kobsiriphat, Y. Wang, L. D. Marks and S. A. Barnett, Nucleation of nanometer-scale electrocatalyst particles in solid oxide fuel cell anodes, J. Power Sources, 2007, 166, 64–67.
202. C. Sun, J. A. Alonso and J. Bian, Recent advances in perovskite-type oxides for energy conversion and storage applications, Adv. Energy Mater., 2020, 11, 2000459.
203. C. Arrivé, T. Delahaye, O. Joubert and G. Gauthier, Exsolution of nickel nanoparticles at the surface of a conducting titanate as potential hydrogen electrode material for solid oxide electrochemical cells, J. Power Sources, 2013, 223, 341–348.
204. P. B. Managutti, S. Tymen, X. Liu, O. Hernandez, C. Prestipino, A. Le Gal La Salle, S. Paul, L. Jalowiecki-Duhamel, V. Dorcet, A. Billard, P. Briois and M. Bahout, Exsolution of Ni nanoparticles from A-Site-deficient layered double perovskites for dry reforming of methane and as an anode material for a Solid Oxide Fuel Cell, ACS Appl. Mater. Interfaces, 2021, 13, 35719–35728.
205. Y. F. Sun, J. H. Li, L. Cui, B. Hua, S. H. Cui, J. Li and J. L. Luo, A-site-deficiency facilitated in situ growth of bimetallic Ni-Fe nano-alloys: a novel coking-tolerant fuel cell anode catalyst, Nanoscale, 2015, 7, 11173–11181.
206. Z. Ma, C. Sun, C. Ma, H. Wu, Z. Zhan and L. Chen, Ni doped La_0.6Sr_0.4 FeO_3−δ symmetrical electrode for solid oxide fuel cells, Chin. J. Catal., 2016, 37, 1347–1353.
207. Y. Tang, H. Wang, R. Wang, Q. Liu, Z. Yan, L. Xu and X. Liu, Synergistically promoting coking resistance of a La_0.4Sr_0.4Ti_0.85Ni_0.15O_3−δ anode by Ru-doping-induced active twin defects and highly dispersed Ni nanoparticles, ACS Appl. Mater. Interfaces, 2022, 14, 44002–44014.
208. W. Wang, R. Ran and Z. Shao, Combustion-synthesized ru–Al₂O₃ composites as anode catalyst layer of a solid oxide fuel cell operating on methane, Int. J. Hydrogen Energy, 2011, 36, 755–764.
209. Z. Zhan and S. Barnett, Use of a catalyst layer for propane partial oxidation in solid oxide fuel cells, Solid State Ionics, 2005, 176, 871–879.
210. C. Sun, Z. Xie, C. Xia, H. Li and L. Chen, Investigations of mesoporous CeO₂–Ru as a reforming catalyst layer for solid oxide fuel cells, Electrochem. Commun., 2006, 8, 833–838.
211. L.-l Sun, L.-l Liu, L.-h Luo, Y.-f Wu, J.-j Shi, L. Cheng, X. Xu and Y.-m Guo, Facile synthesis of flower-like Pd catalyst for direct ethanol solid oxide fuel cell, J. Fuel Chem. Technol., 2016, 44, 607–612.
212. K. Wang, R. Ran and Z. Shao, Methane-fueled IT-SOFCs with facile in situ inorganic templating synthesized mesoporous Sm_0.2Ce_0.8O_1.9 as catalytic layer, J. Power Sources, 2007, 170, 251–258.
213. J. Zhao, X. Xu, W. Zhou, I. Blakey, S. Liu and Z. Zhu, Proton-conducting La-doped ceria-based internal reforming layer for direct methane Solid Oxide Fuel Cells, ACS Appl. Mater. Interfaces, 2017, 9, 33758–33765.
214. P. Qiu, S. Sun, X. Yang, F. Chen, C. Xiong, L. Jia and J. Li, A review on anode on-cell catalyst reforming layer for direct methane solid oxide fuel cells, Int. J. Hydrogen Energy, 2021, 46, 25208–25224.
215. C. Jin, C. Yang, F. Zhao, A. Coffin and F. Chen, Direct-methane solid oxide fuel cells with Cu_1.3Mn_1.7O₄ spinel internal reforming layer, Electrochem. Commun., 2010, 12, 1450–1452.
216. Y. Chen, B. deGlee, Y. Tang, Z. Wang, B. Zhao, Y. Wei, L. Zhang, S. Yoo, K. Pei, J. H. Kim, Y. Ding, P. Hu, F. F. Tao and M. Liu, A robust fuel cell operated on nearly dry methane at 500 °C enabled by synergistic thermal catalysis and electrocatalysis, Nat. Energy, 2018, 3, 1042–1050.
217. Z. Zhan and S. A. Barnett, An octane-fueled solid oxide fuel cell, Science, 2005, 308, 844–847.
218. S. Futamura, A. Muramoto, Y. Tachikawa, J. Matsuda, S. M. Lyth, Y. Shiratori, S. Taniguchi and K. Sasaki, SOFC anodes impregnated with noble metal catalyst nanoparticles for high fuel utilization, Int. J. Hydrogen Energy, 2019, 44, 8502–8518.
219. J. Kim, Y. Choi, D.-K. Lim, J. Yoo, H. G. Seo, S. Kim, S. Kim and W. Jung, Electrochemically plated nickel-decorated ceria nanostructures for direct hydrocarbon solid oxide fuel cell electrodes, J. Mater. Chem. A, 2022, 10, 20886–20895.
220. B. Hua, M. Li, Y.-F. Sun, Y.-Q. Zhang, N. Yan, J. Chen, J. Li, T. Etsell, P. Sarkar and J.-L. Luo, Biogas to syngas: flexible on-cell micro-reformer and NiSn bimetallic nanoparticle implanted solid oxide fuel cells for efficient energy conversion, J. Mater. Chem. A, 2016, 4, 4603–4609.
221. K. Zhao, X. Hou, M. G. Norton and S. Ha, Application of a NiMo–Ce_0.5Zr_0.5O_2−δ catalyst for solid oxide fuel cells running on gasoline, J. Power Sources, 2019, 435, 226732.
222. L. Yang, Y. Choi, W. Qin, H. Chen, K. Blinn, M. Liu, P. Liu, J. Bai, T. A. Tyson and M. Liu, Promotion of water-mediated carbon removal by nanostructured barium oxide/nickel interfaces in solid oxide fuel cells, Nat. Commun., 2011, 2, 357.
223. J. Xu and G. F. Froment, Methane steam reforming, methanation and water-gas shift: I. Intrinsic kinetics, AIChE J., 1989, 35, 88–96.
224. K. Ahmed and K. Föger, Analysis of equilibrium and kinetic models of internal reforming on solid oxide fuel cell anodes: Effect on voltage, current and temperature distribution, J. Power Sources, 2017, 343, 83–93.
225. A. C. van Duin, B. V. Merinov, S. S. Han, C. O. Dorso and W. A. Goddard, 3rd, ReaxFF reactive force field for the Y-doped BaZrO₃ proton conductor with applications to diffusion rates for multigranular systems, J. Phys. Chem. A, 2008, 112, 11414–11422.
226. A. C. T. van Duin, B. V. Merinov, S. S. Jang and W. A. Goddard, ReaxFF reactive force field for solid oxide fuel cell systems with application to oxygen ion transport in yttria-stabilized zirconia, J. Phys. Chem. A, 2008, 112(14), 3133–3140.
227. H. Lu, D. Hua, T. Iqabl, X. Zhang, G. Li and D. Zhang, Molecular dynamics simulations of the coke formation progress on the nickel-based anode of solid oxide fuel cells, Int. Commun. Heat Mass Transfer, 2018, 91, 40–47.
228. J. W. Kim, K. Bae, H. J. Kim, J.-W. Son, N. Kim, S. Stenfelt, F. B. Prinz and J. H. Shim, Three-dimensional thermal stress analysis of the re-oxidized Ni–YSZ anode functional layer in solid oxide fuel cells, J. Alloys Compd., 2018, 752, 148–154.
229. Y. Guo, M. Bessaa, S. Aguado, M. C. Steil, D. Rembelski, M. Rieu, J.-P. Viricelle, N. Benameur, C. Guizard, C. Tardivat, P. Vernoux and D. Farrusseng, An all porous solid oxide fuel cell (SOFC): a bridging technology between dual and single chamber SOFCs, Energy Environ. Sci., 2013, 6, 2119–2123.
230. Y. M. Guo, G. Largiller, C. Guizard, C. Tardivat and D. Farrusseng, Coke-free operation of an all porous solid oxide fuel cell (AP-SOFC) used as an O₂ supply device, J. Mater. Chem. A, 2015, 3, 2684–2689.
231. H. Xu, B. Chen, P. Tan, W. Cai, W. He, D. Farrusseng and M. Ni, Modeling of all porous solid oxide fuel cells, Appl. Energy, 2018, 219, 105–113.
232. E. Audasso, F. R. Bianchi and B. Bosio, 2D simulation for CH₄ internal reforming-SOFCs: an approach to study performance degradation and optimization, Energies, 2020, 13, 4116.
233. Y. Su, T. Wei, Y. Li, B. Yin, Y. Huan, D. Dong, X. Hu and B. Huang, A highly active CH₄ catalyst correlated with solid oxide fuel cell anode performance, J. Mater. Chem. A, 2021, 9, 5067–5074.
234. X.-F. Ye, B. Huang, S. R. Wang, Z. R. Wang, L. Xiong and T. L. Wen, Preparation and performance of a Cu–CeO₂–ScSZ composite anode for SOFCs running on ethanol fuel, J. Power Sources, 2007, 164, 203–209.
235. X.-F. Ye, S. R. Wang, Q. Hu, J. Y. Chen, T. L. Wen and Z. Y. Wen, Improvement of Cu–CeO₂ anodes for SOFCs running on ethanol fuels, Solid State Ionics, 2009, 180, 276–281.
236. X. Hao, D. Han, J. Wang, Y. Liu, D. Rooney, W. Sun, J. Qiao, Z. Wang and K. Sun, Co-tape casting fabrication, field assistant sintering and evaluation of a coke resistant La_0.2Sr_0.7TiO₃–Ni/YSZ functional gradient anode supported solid oxide fuel cell, Int. J. Hydrogen Energy, 2015, 40, 12790–12797.
237. M. Wang, N. Li, Z. Wang, C. Chen and Z. Zhan, Electrochemical performance and redox stability of solid oxide fuel cells supported on dual-layered anodes of Ni–YSZ cermet and Ni–Fe alloy, Int. J. Hydrogen Energy, 2022, 47, 5453–5461.
238. S. Sengodan, S. Choi, A. Jun, T. H. Shin, Y. W. Ju, H. Y. Jeong, J. Shin, J. T. Irvine and G. Kim, Layered oxygen-deficient double perovskite as an efficient and stable anode for direct hydrocarbon solid oxide fuel cells, Nat. Mater., 2015, 14, 205–209.
239. M. Li, B. Hua, Y. Zeng, B. Shalchi Amirkhiz and J.-L. Luo, Thermally stable and coking resistant CoMo alloy-based catalysts as fuel electrodes for solid oxide electrochemical cells, J. Mater. Chem. A, 2018, 6, 15377–15385.
240. R. Lang, T. Li, D. Matsumura, S. Miao, Y. Ren, Y. T. Cui, Y. Tan, B. Qiao, L. Li, A. Wang, X. Wang and T. Zhang, Hydroformylation of olefins by a rhodium single-atom catalyst with activity comparable to RhCl(PPh₃)₃, Angew. Chem., Int. Ed., 2016, 55, 16054–16058.
241. Y. Pan, S. Liu, K. Sun, X. Chen, B. Wang, K. Wu, X. Cao, W. C. Cheong, R. Shen, A. Han, Z. Chen, L. Zheng, J. Luo, Y. Lin, Y. Liu, D. Wang, Q. Peng, Q. Zhang, C. Chen and Y. Li, A bimetallic Zn/Fe polyphthalocyanine-derived single-atom Fe-N₄ catalytic site: a superior trifunctional catalyst for overall water splitting and Zn-air batteries, Angew. Chem., Int. Ed., 2018, 57, 8614–8618.
242. J. Han, J. Bian and C. Sun, Recent advances in single-atom electrocatalysts for oxygen reduction reaction, Research, 2020, 2020, 9512763.
243. Y. Tang, Y. Wei, Z. Wang, S. Zhang, Y. Li, L. Nguyen, Y. Li, Y. Zhou, W. Shen, F. F. Tao and P. Hu, Synergy of single-atom Ni₁ and Ru₁ sites on CeO₂ for dry reforming of CH₄, J. Am. Chem. Soc., 2019, 141, 7283–7293.
244. J. Wu, J. Gao, S. Lian, J. Li, K. Sun, S. Zhao, Y. D. Kim, Y. Ren, M. Zhang, Q. Liu, Z. Liu and Z. Peng, Engineering the oxygen vacancies enables Ni single-atom catalyst for stable and efficient C-H activation, Appl. Catal., B, 2022, 314, 121516.
245. C. T. Campbell, The energetics of supported metal nanoparticles: relationships to sintering rates and catalytic activity, Acc. Chem. Res., 2013, 46, 1712–1719.
246. K. Liu, X. Zhao, G. Ren, T. Yang, Y. Ren, A. F. Lee, Y. Su, X. Pan, J. Zhang, Z. Chen, J. Yang, X. Liu, T. Zhou, W. Xi, J. Luo, C. Zeng, H. Matsumoto, W. Liu, Q. Jiang, K. Wilson, A. Wang, B. Qiao, W. Li and T. Zhang, Strong metal-support interaction promoted scalable production of thermally stable single-atom catalysts, Nat. Commun., 2020, 11, 1263.
247. X. Li, Z. Chen, Y. Yang, D. Huan, H. Su, K. Zhu, N. Shi, Z. Qi, X. Zheng, H. Pan, Z. Zhan, C. Xia, R. Peng, S. Wei and Y. Lu, Highly stable and efficient Pt single-atom catalyst for reversible proton-conducting solid oxide cells, Appl. Catal., B, 2022, 316, 121627.

---