Progress in Solid Oxide Fuel Cells with Nickel-Based Anodes

Review pubs.acs.org/CR

Progress in Solid Oxide Fuel Cells with Nickel-Based Anodes Operating on Methane and Related Fuels Wei Wang,† Chao Su,‡ Yuzhou Wu,§ Ran Ran,† and Zongping Shao*,† †

State Key Laboratory of Materials-Oriented Chemical Engineering, College of Chemistry & Chemical Engineering, Nanjing University of Technology, No. 5 Xin Mofan Road, Nanjing 210009, People’s Republic of China ‡ Department of Chemical Engineering, Curtin University, Perth, WA 6845, Australia § Department of Chemical Engineering, Monash University, Clayton, VIC 3800, Australia 4.1. Types and Critical Amount of Sulfur Contaminants 4.2. Possible Poisoning Mechanisms 4.3. Strategies for Sulfur Mitigation 4.3.1. Desulfurization 4.3.2. Optimization of the Operating Conditions 4.3.3. Sulfur-Tolerant Anode Materials 5. Redox Stability 5.1. Redox Behavior of Ni-Based Cermet Anodes 5.2. Possible Solutions 6. Cogeneration of Electric Power and Syngas 7. SOFCs Operating on Oxygenated Methane Fuels 7.1. Dimethyl Ether 7.2. Methanol 8. Conclusions and Prospects Author Information Corresponding Author Notes Biographies Acknowledgments Abbreviations References

CONTENTS 1. Introduction 2. Methane Electro-oxidation Mechanism and Kinetics 2.1. Direct Oxidation and Indirect Oxidation 2.2. Reaction Kinetics 2.3. Mechanistic Studies Based on Model Electrodes 2.3.1. Model Anodes and Anode Models 2.3.2. Detailed Investigations 3. Coke Formation over Nickel Catalysts 3.1. Mechanism and Kinetics 3.2. Factors Controlling the Coke Formation 3.2.1. Thermodynamics 3.2.2. Nickel Particle Size 3.2.3. Alloying Effect 3.2.4. Second Phase(s) 3.2.5. Gas Composition 3.3. Coke Formation over Sintered SOFC Nickel Cermet Anodes 3.4. Strategies To Suppress Coke Formation over Nickel Cermet Anodes 3.4.1. Addition of Other Gases 3.4.2. Polarization Current 3.4.3. Decoration of the Electrode Surface or Alloying of the Nickel with Other Metal(s) 3.4.4. Anode Catalyst Layer 3.4.5. Tailoring the Ceramic Phase 3.4.6. Modification of Electrode Surface with Other Oxide(s) 4. Sulfur Poisoning

© XXXX American Chemical Society

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1. INTRODUCTION During the past century, fossil fuels, in particular coal and oils, have contributed significantly to the development of our industry, economy, and modern conveniences. However, the excessive use of such nonrenewable natural resources based on low-efficiency combustion technology has also introduced serious problems, such as increasing CO2 concentrations in the atmosphere, which have caused a detrimental global warming effect, and worsening environmental conditions due to the emissions of NOx and SOx impurities. In addition, although there is still no consensus on how long fossil fuels may remain the main energy materials of our society because new resources are discovered from time to time, the harvesting of those new energy resources is increasingly difficult. With increasing demand for energy and ever-diminishing resources of fossil fuels, the exploration for new energy sources and improvements in the efficiency of energy utilization for the currently available resources has become increasingly important to achieve a sustainable future.

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Received: December 11, 2012

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dx.doi.org/10.1021/cr300491e | Chem. Rev. XXXX, XXX, XXX−XXX

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charged species they transport, and the operating temperatures.26−31 In the case of polymer electrolyte membrane fuel cells (PEMFCs), which operate at low temperatures, methane cannot be directly electro-oxidized because of the inactivity of platinum electrodes to methane at room temperature. Instead, it should be reformed into a gas mixture that is composed mainly of H2. Furthermore, because the CO would poison the platinum catalyst in the anode of PEMFCs, the selective oxidation of CO to CO2 under hydrogen-rich conditions is also required to reduce the CO concentration to less than 10 ppm.32 When the operating temperature is increased, the fuel cell anodes become increasingly active toward the electro-oxidation of methane and other fuels. The solid oxide fuel cell (SOFC) is a type of fuel cell that typically operates at approximately 600− 1000 °C.33 At such high temperatures, either the direct electrochemical oxidation or the internal reforming/partial oxidation of methane becomes kinetically favorable. Recently, the operation of SOFCs with methane or methane-based fuels has attracted tremendous attention.34−41 Some significant advances have been made in both the both fundamental and the technological aspects. For example, a 100 kW SOFC system fed by biogas mixtures has recently been demonstrated, and the electrical efficiency of this system was approximately 48.7%,42 whereas the efficiency of conventional electricity-generation systems based on gas turbines is typically 41.5%.43 Conventional SOFCs employ nickel-based cermet anodes, which exhibit good compatibility with electrolytes composed of stabilized zirconia or doped ceria, high activity toward the electrocatalytic oxidation of hydrogen, and high electronic conductivity.44 By direct feeding of a nickel cermet anode-based SOFC with methane, the anode may play dual roles: catalyzing the methane reforming reactions and promoting the electrochemical oxidation of hydrogen, carbon monoxide, and methane. However, because the nickel-based anodes tend to carbon deposition in methane fuel, quick carbon accumulation over the anode is often observed and leads to rapid degradation of the fuel cell’s performance. In addition, sulfur in the fuel gas could cause serious poisoning effects on the nickel cermet anodes. However, the nickel cermet anodes are also prone to reoxidation by the oxidant, which may significantly affect the cell integrity. Recently, tremendous efforts have been devoted to improve the coking resistance, sulfur resistance, and redox stability of methane and other hydrocarbons fueled SOFCs, that is, by adopting perovskite-based or non-nickel-based anodes.45−48 Several excellent review papers are available in the literature on the development of SOFC anode materials and the relative challenges facing their continued development.49−54 For example, McIntosh and Gorte have studied the development of anode materials for SOFCs operating on hydrocarbon fuels directly,49 whereas Tao and Irvine have reviewed the development of alternative oxide anodes with different structures for hydrocarbon-fueled SOFCs.50 Sun and Stimming have also summarized the anode materials for SOFCs operating on hydrocarbons, the development of anode kinetics and reaction mechanisms, as well as the anode models and the cost-effective processing technologies for anode preparation.51 Huang and Goodenough have noted several constraints on the design of anode materials for SOFCs, and also discussed the development of anodes for SOFCs operating on H2 and CO.52 One important drawback of perovskite anodes is their poor electrical conductivity under anodic atmospheres, which introduces the difficulty of current collection. Furthermore, the catalytic activity of perovskite anodes for hydrocarbon

Natural gas is a fossil fuel that is much cleaner than coal and oil because of the greater hydrogen-to-carbon ratio in its molecular composition and because it contains smaller quantities of N and S impurities. As the simplest hydrocarbon, methane is the principal component of natural gas and coal-bed gas. Methane is also the main component of biogas (the remainder is CO2), which can be generated by the fermentation of organic materials, including manure, wastewater sludge, municipal solid waste (including landfill waste), or any other biodegradable feedstock, under anaerobic conditions;1−4 thus methane could be considered a renewable energy resource. Methane can also be produced by the hydrogenation of CO2 through the Sabatier process.5,6 However, methane is a greenhouse gas with a global warming potential of approximately 25, which means that every kilogram of methane emitted to the atmosphere has the equivalent effect on the Earth’s climate of 25 times that amount of CO2 over a 100 year period.7 Therefore, proper utilization of methane is critically important to avert an energy crisis as well as to reduce the impact of global warming. Natural gas (methane) has already been widely used as an important energy source in numerous applications, including the heating of buildings and cooking by direct combustion, the generation of electricity by gas turbines or steam turbines,8 and as a fuel for combustion engines in vehicles.9 It can also be used to produce hydrogen via steam reforming.10−14 To achieve the sustainable development of our society in future decades, it is important to improve the energy efficiency and reduce CO2 emissions during the utilization of methane because the world’s energy demand for natural gas is predicted to increase greatly in the future. One of the promising solutions is the development of fuel cell technology. Fuel cells are electrochemical devices that convert the chemical energy of fuels directly into electric power without the limitation of the Carnot cycle.15,16 Therefore, fuel cell technology is considered to be a more efficient and cleaner alternative to the traditional electricity generation technique that involves heat engines, steam and gas turbines, and combined cycles. An external reforming or partial oxidation process is normally required for the conversion of methane to hydrogen before it is fed as fuel into fuel cell systems.17−19 To increase the hydrogen content, after the steam reforming reaction, a subsequent water gas shift (WGS) reaction is conducted to obtain a gas mixture that contains primarily H2 and CO2.20 A large volume of research focused on the production of hydrogen via the reforming of methane has recently been reported,21−24 and an excellent review paper has been published by Armor.24 However, the complicated reforming process will consume additional energy and lead to reduced overall fuel efficiency. For example, it was reported that the efficiency of internal reforming was 8 percentage points higher than that of external reforming in the same SOFC system.25 On the other hand, Sangtongkitcharoe et al. have found that, due to the presence of extra H2O from the electrochemical reaction at the anode chamber, internalreforming SOFC can be operated at lower values of the H2O/CH4 ratio as compared to the external reforming mode. The low requirement of H2O could reduce the energy of water gasification.17 Hence, the feeding of methane directly into the fuel cell system without an external reforming process is the preferred approach. Several types of fuel cells exist, and they are distinguished from each other according to their electrolyte materials, the B


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pathways because the demands on the anode material differ significantly between the two cases. The anode for the direct electrochemical oxidation of hydrocarbons should have a high turnover rate for the hydrocarbon bond breaking, whereas cracking followed by an electrochemical conversion of the cracking products in the indirect oxidation process requires an anode that is highly active for both the methane cracking reaction and the further oxidation of the cracking products. To distinguish between the possible electrochemical and chemical reactions in methane-fueled SOFCs, the possible reactions in the fuel cells are first discussed. In a fuel cell system operating at high temperatures, methane may decompose easily not only at the three-phase boundary (TPB) but also in the whole anode region. The anode region contains a maximum of six components that form predominantly due to their thermodynamic stabilities: the deposited carbon, CO, H2, CO2, H2O, and some residual CH4. Certain species, such as CH4, CO, and H2, may react with the O2− electrochemically. The possible electrochemical reactions of these species in the anode side are listed as follows:

conversion was low, consequently, a relative low power output was obtained. On the other hand, the CeO2−Cu-based electrodes usually show poor electro-catalytic activity toward the oxidation of methane or related fuels, which results in lower power output and diminished economical attractiveness. The nickel cermet anodes have shown superior performance in terms of activity toward hydrogen electrocatalytic oxidation, electronic conductivity, and thermo-mechanical compatibility with state-of-the-art electrolytes. Therefore, nickel cermets are still one of the most attractive anodes for SOFCs. During the past decades, the modification of nickel cermet anodes to improve their performance has been extensively exploited. In this Review, the recent progress in direct-methane (or related fuels) SOFCs with nickel-based cermet anodes is comprehensively reviewed. Emphasis is placed on the coke formation over nickel-based anodes, and related strategies for improving the coking resistance of the anodes and the power output of the related fuel cells. The methodologies used to reduce the coke formation thermodynamically or kinetically, such as the introduction of additional gas (H2, O2, CO2, and H2O) into the fuel gas, the formation of alloy anodes, the modification of the nickel cermet with other metal(s) or oxide(s), the application of anode catalyst layer, and so on, are extensively reviewed. The sulfur poisoning and redox stability issues are also addressed. The electricity and syngas cogeneration using SOFC technology is discussed. Methanol and dimethyl ether (DME) are also promising alternative fuels of the future and are easier to store and transport than methane because they are liquids at room temperature under ambient pressure (methanol) or medium pressure (DME). From a structural point of view, both methanol and DME can be treated as oxygenated methane. Consequently, the recent progress in methanol- and DME-fueled SOFCs with nickelbased anodes is also briefly summarized. This Review aims to provide some useful guidelines for the further development of practical nickel-based cermet anodes for SOFCs operating directly on methane or related fuels. Because methane is the simplest hydrocarbon, the information presented in this Review about methane-fueled SOFCs is also instructive for the development of materials and mechanistic understanding of SOFCs operating on other types of hydrocarbons or oxygenated hydrocarbons.

CH4 + O2 − → CO + 2H 2 + 2e−

(1)

CH4 + 2O2 − → CO + H 2 + H 2O + 4e−

(2)

CH4 + 2O2 − → CO2 + 2H 2 + 4e−

(3)

CH4 + 3O2 − → CO + 2H 2O + 6e−

(4)

CH4 + 3O2 − → CO2 + H 2 + H 2O + 6e−

(5)

CH4 + 4O2 − → CO2 + 2H 2O + 8e−

(6)

H 2 + O2 − → H 2O + 2e−

(7)

CO + O2 − → CO2 + 2e−

(8)

whereas the possible chemical reactions are:

2. METHANE ELECTRO-OXIDATION MECHANISM AND KINETICS 2.1. Direct Oxidation and Indirect Oxidation

Over the past few years, many researchers have been developing SOFC anodes for operating on hydrocarbon fuels. Mogensen and Kammer defined the direct conversion of hydrocarbons in SOFCs as either the direct electrochemical oxidation of the fuel or the electrochemical oxidation of the cracking products, and direct oxidation means that all steps in the reaction of methane conversion must be electrochemical in nature.55 The primary reason for using this specific definition for direct oxidation is that the open circuit voltage (OCV) of the fuel cell should be equal to the theoretical Nernst potential if no other electrochemical losses are occurring and all steps in the oxidation mechanism are electrochemical. According to this definition, both electrochemical reactions and chemical reactions occur in the indirect oxidation process, whereas only electrochemical reactions occur in the direct oxidation process. It is important to distinguish between the two reaction

CH4 + H 2O → CO + 3H 2

(9)

CH4 + CO2 → 2CO + 2H 2

(10)

H 2 + CO2 → CO + H 2O

(11)

CH4 → C + 2H 2

(12)

2CO ↔ CO2 + C

(13)

C + H 2O → CO + H 2

(14)

The direct electrochemical oxidation of methane in a SOFC was first reported in detail by Steele et al., who observed that methane could be effectively converted into CO2 and H2O on oxide-based anodes. 56 Ceria was also mentioned as a particularly good candidate for an anode material for the direct electrochemical oxidation of methane. Later, detailed investigations proved that reduced ceria is almost inactive with respect to the C−H bond breaking reaction.57 Thus, the main processes that occurred on the widely studied copper−ceria anodes and perovskite anodes were also the direct oxidation process. However, the processes on the state-of-art Ni-based anodes that operated on methane represent indirect oxidation because Ni is a good catalyst for C−H bond breaking as well as for the conversion of methane to hydrogen and carbon monoxide. C


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mechanism was dominant.76 In the MVK mechanism, the reaction occurs via the reduction and oxidation of lattice oxygen sites. The reduction step in the SOFC anode occurs via fuel oxidation at the active surface sites with a two-electron charge transfer, and the reoxidation step occurs by bulk ion transport. On the basis of the observed strong influence of the B-site cations, McIntosh et al. have used a point-defect approach to assume that the active site was likely a Bn+−O−Bn+ complex, which was reduced to form an oxygen vacancy, B(n−1)+−□− B(n−1)+.77 The main oxidation process on Ni-based anodes differs from that on perovskite anodes in that it is likely an indirect pathway because of the high catalytic activity of Ni toward the breakage of C−H bond in hydrocarbons. The reaction kinetics with methane and the products of chemical reactions such as H2, CO, and CO2 over nickel anodes have been studied recently. The mechanism of H2O dissociation was also studied because H2O is the product of the hydrogen electrochemical oxidation process. Some researchers have reported that the anodic oxidation of hydrogen could be limited by several processes, such as the electrochemical and chemical reaction steps at the TPB, physical transport restrictions for electrons and oxygen ions in the solid structure, surface adsorption/diffusion of species such as Had, Oad, and OH−ad on the anode surface or the diffusion of H from the anode metal to the active reaction sites, and gas diffusion in the anode pores/channels.66,78,79 Some electrochemical models are focused mainly on the adsorption and desorption behavior of hydrogen as well as the formation of hydroxyl ions. For example, Holtappels et al. considered the hydrogen oxidation reaction to be a two-step electrochemical reaction mechanism that involves the dissociative adsorption of H2 onto Ni and a charge-transfer process, where the water formation reaction was believed to have occurred via intermediates such as hydroxides on the Ni surface.78,80 Similarly, Jiang and Badwal demonstrated that hydrogen oxidation in Ni anodes was determined by two processes on the surface of Ni particles: hydrogen dissociative adsorption or diffusion and a charge transfer process on the zirconia electrolyte surface.65 In fact, the two previously discussed reaction mechanisms for H2 oxidation are similar in nature because the authors all considered that the H2 oxidation process was determined by both an adsorption process and a charge transfer process on the surface of the Ni particles. Their major differences are the role of the YSZ electrolyte plays in the hydrogen oxidation and the adsorption/desorption behavior of water. Jiang and Badwal found that, when H2O was mixed with H2, the H2O preferentially adsorbed onto the Ni surface, which formed OH−Ni as a reaction species. They also found that the dissociative adsorption/diffusion of hydrogen on the metal surface was enhanced, most likely through a spillover effect. The rate of this process could be enhanced via an increase in the oxygen partial pressure.81 Although CO has seldom been directly used as a fuel for SOFCs because of its toxicity for humans, it is an important intermediate product during the indirect oxidation of methane. Thus, the electrochemical oxidation of CO over SOFC anodes has also attracted attention recently. For example, Matsuzaki et al. have studied the electrochemical characteristics of a H2+H2O+CO+CO2 system within a Ni−YSZ cermet anode. They found that the electrochemical oxidation rate of hydrogen was approximately 2−3 times greater than that of CO depending on the operating temperature.82 They attributed the slower oxidation rate of CO relative to that of hydrogen to

As previously mentioned, the OCV of a fuel cell can be used to predict the possible oxidation mechanism on the anodes. A fuel cell with a BaTiO3 anode was observed to deliver an OCV greater than 1.20 V with 0.5% H2S−CH4 as fuel at 900 °C,58 and the theoretical value at the same temperature was 1.27 V. However, an OCV of approximately 1.18 V was obtained with a Ni−ScSZ anode using 3% H2O−CH4 at 1000 °C,59 and the theoretical value at the same temperature was 1.33 V. For comparison, the OCV was only approximately 1.05 V for a Ni− YSZ anode using 3% H2O−97% H2 fuel,59 and the theoretical value at the same temperature was 1.067 V. Some researchers also studied the OCV of carbon monoxide-rich fuels for SOFCs. They found that the OCV was approximately 1.04 V for a Ni−YSZ anode with 5% CO2−95% CO fuel, which was approximately 0.02 V lower than the theoretical value at the same temperature.60 The previously discussed results further suggest that the degree of direct oxidation is greater over the perovskite anodes than over the Ni-based anodes at operating temperatures higher than 700 °C. However, the OCV of perovskite anodes operated on methane decreased significantly as the operating temperature was decreased,58 which implies that the reaction kinetics likely also played an important role at the reduced temperatures with respect to the cell potential. The theoretical OCV is well-known to be a thermodynamic parameter that is, in principle, dependent on the chemical oxygen partial pressure of the cathode and anode while being independent of kinetics. However, in practice, significant oxygen exchange must occur between the electrode and the fuel gas, which could affect the actual OCV values. Low OCV values in CH4 and enhanced OCV values upon an increase in the catalytic rate have been reported numerous times. For example, a Cu−YSZ anode yielded an OCV of only 0.60 V at 700 °C in dry CH4; however, the addition of a ceria catalyst increased this OCV to ∼0.90 V.61 Van den Bossche and McIntosh have demonstrated that the rate and selectivity of bare La0.75Sr0.25Cr0.5Mn0.5O3−δ (LSCM) toward CH4 oxidation are low and significantly lower than the rate and selectivity toward H2.62 Indeed, the reported oxygen exchange rate k of LSCM at approximately 700 °C is extremely low, with a value of only 5.87 × 10−8 cm s−1.63 The combination of low oxidation selectivity and activity leads to a low OCV in pure CH4 fuel on a LSCM anode. The study of the kinetics of the reactions that occur over the anode is important to understand the OCV of a cell operated on hydrocarbon fuels and also for the future development and optimization of the anode materials. The anode performance may be considerably enhanced if the limiting processes for anode reactions are improved. 2.2. Reaction Kinetics

Numerous mechanistic theories have been proposed to identify the rate-limiting mechanism under various operating conditions, including reactivity and charge transfer,64−66 adsorption,65,67 hydrogen desorption rates,64,67 catalytic effects of water,67 the role of the YSZ support,66 and so on. The methane oxidation activity of perovskite SOFC anodes and those of related materials have been extensively studied.62,68−73 The activity is primarily determined by the nature and oxidation state of the B-site cation,68,71 with strong dependence on the oxygen stoichiometry and lattice oxygen ion mobility.69 At lower temperatures, the reaction was considered to proceed via an Eley−Rideal mechanism with mobile adsorbed surface oxygen as the active species.69,74,75 However, at higher temperatures, a Mars−van Krevelen (MVK) D


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the greater surface diffusion resistance of CO at low temperatures and to the diffusion and charge transfer at high temperatures. However, the SOFC, when operated on CO+H2 gas mixtures, still exhibited a performance comparable to that of hydrogen because the WGS reaction was found to be much faster than the CO electrochemical oxidation reaction over the temperature range studied.82 Because the electrochemical oxidation rate of CO is generally considered to be significantly slower than that of H2, the electrochemical oxidation of carbon monoxide has often been completely ignored in the modeling of fuel cell performance, except for the WGS reaction. For example, Aguiar et al. only considered H2 oxidation and completely neglected any carbon monoxide oxidation during their investigation of the effects of methane steam reforming and the WGS reaction in a one-dimensional model of an anodesupported SOFC operated on a prereformed methane−steam mixture.83 However, some recent studies also demonstrated that even pure CO could still produce a power output comparable to that of hydrogen at 850 °C.60 The electrode microstructure likely significantly influences the CO oxidation rate. More systematic research is still needed to establish reliable kinetic models for the operation of SOFCs on methane fuel through indirect oxidation. Figure 1. Schematic illustration of (a) a technical porous cermet anode and (b) geometrically well-defined model anodes for the Ni−YSZ materials system. Reproduced with permission from ref 84 (http://dx. doi.org/10.1039/C0CP005415). Copyright 2010 The Royal Society of Chemistry.

2.3. Mechanistic Studies Based on Model Electrodes

Notably, most of the previously mentioned kinetic models were established on the basis of electrodes with uncontrollable TPB length due to complicated electrode morphology, which may introduce large errors. A SOFC anode is a multiphase system that includes a solid-electrolyte phase, a gas phase, and a solidelectrode phase, which leads to the concept of the TPB. Electrochemical reactions can only occur in locations where all three components are spatially close to each other. Technical SOFC anodes are designed to have a high TPB length connected to percolating transport pathways for electrons, ions, and gaseous species. This design is schematically shown for a technical Ni−YSZ cermet anode in Figure 1a.84 Although the porous structure results in good electrochemical performance, it is not well suited for studying the processes at the TPB itself because the active regions are deep inside the porous structure and are not easily accessible using microscopic or spectroscopic techniques; furthermore, the exact microstructural properties and TPB length are usually unknown. To allow more precise investigations of the electrode reaction processes both experimentally and theoretically, model anodes with a wellcontrolled TPB length are required. 2.3.1. Model Anodes and Anode Models. Model electrodes have experimentally accessible active surfaces, porous transport effects are absent, and they are expected to have geometrically well-defined two-phase and three-phase boundaries. The best known model anodes include micropatterned thin-film electrodes and point electrodes, as shown schematically in Figure 1b. For the patterned anodes,67,85−87 metal patterns are photolithographically applied onto a solidelectrolyte substrate. This technique allows well-defined contact areas and long TPB lengths. In the fabrication of point electrodes,88−91 a metal wire with a diameter in the range of 0.5 mm was bent and pressed against a solid-electrolyte pellet or crystal, which led to elliptical contact areas in the range of 0.1 mm2 and absolute TPB lengths in the range of 1 mm. Researchers have made significant progress in the microfabrication of dense, patterned electrodes in which geometric dimensions, such as the thickness and the TPB length, can be

specifically controlled. In this way, various asymptotes of behavior (such as the surface and bulk paths) have been examined and various electrode mechanisms have been probed using a variety of techniques, including impedance92 and isotope tracer experiments.93 Recently, the mechanism of the oxygen reduction reaction (ORR) on the cathodes of SOFCs has been widely studied with the use of model electrodes.94 Mathematical modeling opens the possibility of a physically based interpretation of experimental data and the possibility of a model-based optimization of the electrode design and performance in a further step. To understand the complex and nonlinear behavior observed in the modeled anode performance, a number of different modeling approaches have been developed and applied, and these approaches range from a nanoscopic (quantum chemistry, molecular dynamics) to a mesoscopic (elementary kinetics) to a macroscopic (empirical global kinetics) description. The most investigated models include atomistic modeling, such as density functional theory (DFT) studies and elementary kinetic modeling, which refers to the description of chemical reactions via their individual reaction steps.84 2.3.2. Detailed Investigations. DFT has been used for ab initio thermodynamics calculations of a number of metal surfaces in the context of the electrochemical oxidation of H2 and CH4.95,96 For example, Linic and co-workers have presented a theoretical approach, combining DFT calculations and statistical thermodynamics, aimed at studying electrochemical surface reactions. They used this approach to investigate a mechanism for the electrochemical oxidation of hydrogen and methane, steam and dry reforming of methane, electrochemical oxidation of H2 and CO, on various metal surfaces under SOFC operating conditions. All calculations were performed with the Dacapo19 code by plane wave basis set and Perdew−Wang 91 (PW91) functional on a slab E


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model.97 DFT has also been used to directly study hydrogen spillover.98,99 Anderson and Vayner have explored the hydrogen anode reaction chemistry at the Ni−zirconia TPB in SOFCs by using hybrid DFT (B3LYP) calculations and cluster models. The activation energy for hydrogen spillover is calculated to be the same order of magnitude as experimental estimates at the reversible potential. Shishkin and Ziegler have studied the activation of fuel molecules (H2, CH4, and CO) at the anode TPB modeled by a Ni/YSZ/fuel interface using DFT.99 All calculations presented in that work were performed employing a periodic DFT method implemented in the Vienna ab initio Simulation Package (VASP) code. The exchangecorrelation interactions were treated with the help of the Perdew−Burke−Ernzerhof (PBE) functional.100 They demonstrated that, by employing ab initio calculations, it is possible to elucidate the mechanisms of electronic charge transfer and current generation as a result of electrochemical oxidation of fuel in the anodic TPB. They showed that the transfer of hydrogen atoms from Ni to YSZ (i.e., hydrogen spillover reactions) with subsequent water formation was another possible scenario of hydrogen oxidation on the YSZ+O surface. These results for oxygen and hydrogen spillover reactions could be used for the kinetic modeling of electrochemical methane oxidation in the anode TPB of the SOFC.98 The considerable accuracy and the capacity to unveil a reaction on atomic scale have made DFT a standard technique in most branches of chemistry and materials science. However, DFT has many limitations in its present form: the high computational cost makes it barely available to predict conditions rather than ideal model, and there are some other limitations like failures to describe systems with strongly correlated d and f electrons (which is very important for metal oxide modeling), incapacity to describe the long-range London dispersion, etc. Despite the enormous advances in density functional implementations and the sufficiently documented accuracy of results for many applications, there is no systematic way of improving DFT or converging its results to the “correct” answer, in contrast to some of the traditional methods like second-order approximation of Møller−Plesset perturbation theory (MP2) or Coupled Cluster Singles and Doubles (CCSD). Moreover, the success of a particular functional in one setting does not guarantee its performance in a different one. Therefore, to further enhance the credibility of DFT, DFT applications must include some form of validation or estimation of the error range on the basis of careful comparison between calculated and measured observables.101 Atomistic models could give a detailed insight into the structure and energetic on the subnanometer scale. However, they cannot predict macroscopic electrochemical behavior (e.g., polarization curves or impedance spectra). In elementary kinetics modeling, the reaction rates are described by mass actions laws. Elementary kinetic models are based on a continuum spatial description (i.e., no atomic-scale resolution) that can be coupled to meso- and macro-scale transport processes. This coupling allows direct modeling of the geometries of pattern anodes and the diffusion processes that occur on the submicrometer scale in the bulk or on the surface. Elementary kinetic models of SOFC model anodes have been published in 0D,102 1D (including surface diffusion),103 and 2D (including both surface diffusion and bulk diffusion).104 For example, Yurkiv et al. have presented a combined modeling and experimental analysis of the electrochemical oxidation of CO on well-defined Ni−YSZ pattern model anodes.105 They

developed a computational model that represents the coupled behavior of heterogeneous chemistry and electrochemistry in terms of elementary reactions, which allowed a quantitative description of the electrochemical impedance spectra and current−voltage response. Excellent agreement between the model and experiments was achieved with a wide range of CO +CO2+N2 gas compositions and operating temperatures. The authors found that the CO and CO2 concentrations had a strong and nonlinear influence on the electrode kinetics and that an increase in the partial pressure of either CO or CO2 could enhance the electrode kinetics. The authors attributed this counterintuitive effect to a simultaneous change in the electric potential and in the coverage of species that participate in the charge-transfer reaction. The gas composition also influenced the macroscopically observed activation energy. Charge transfer was found to proceed by two different mechanisms. At high CO-to-CO2 ratios, oxygen spillover from the YSZ to the Ni surface occurred, and CO was oxidized in a Langmuir−Hinshelwood-type heterogeneous reaction on the Ni surface. At low CO-to-CO2 ratios, the largest portion of oxygen ions was fully reduced on the YSZ surface without undergoing spillover and CO was oxidized in an Eley−Ridealtype heterogeneous reaction on the YSZ surface.102 The authors believed that the level of understanding achieved in the present study could be the basis for further investigations of the behavior of more complex SOFC fuels, especially reformate gases or hydrocarbons. In addition to the various detailed investigations on anode models, model anodes have also been widely used in kinetics studies on hydrogen and/or carbon monoxide oxidation processes.85−87,106,107 Sukeshini et al. have used thin-film, sputter-deposited Ni pattern anodes microfabricated on singlecrystal YSZ electrolyte disks to examine the electrochemical oxidation of H2, CO, and CO+H2 mixtures.86 They found that the anode potentials and polarization resistances were correlated with the original Ni pattern area for the various pattern geometries. At temperatures greater than 750 °C, they found the Ni−YSZ anode showed a significantly increased rate of H2 oxidation in comparison to that of CO oxidation; furthermore, H2 oxidation was the dominant process during the oxidation of CO/H2 gas mixtures. The authors considered that adsorption/desorption equilibration on the anode surface played a decisive role in the oxidation of both H2 and CO at higher temperatures. Thus, the further exploration of this hypothesis through examination of the model’s predicted impedance spectra and polarization curves using multistep chemistry for H2 and CO electrochemical oxidation on Ni− YSZ anodes is important. Greater activation overpotentials and polarization resistances were observed for CO than H2, and they were not observed for CO+H 2 mixtures with H 2 concentration as low as 25%. The results indicated that the detrimental effects of H2O on the oxidation of CO were decreased OCVs, whereas reduced polarization resistances and enhancements due to WGS reactions were not observed. These results could provide the basis for insights into H2 and CO electro-oxidation on Ni−YSZ anodes. Notably, the sputter-deposited thin Ni pattern films often transformed into an array of interconnected Ni agglomerates during testing at high temperatures, which contributed to large and not well-quantified TPB lengths.86 Thus, control of the Ni agglomeration behavior in patterned thin film anodes is critical, and more work is needed on this aspect of anode development. F


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3. COKE FORMATION OVER NICKEL CATALYSTS 3.1. Mechanism and Kinetics

For fuel cells with nickel cermet anodes that operate on methane fuel, the high operation temperatures and the presence of methane combined with steam or CO2 create a harsh environment for the nickel-based anodes. In particular, four main challenges impede the further development of anode catalysts: low catalytic activity, sulfur poisoning, sintering, and coke formation. Among these challenges, coke formation is the most important one and is closely related to the other three challenges. Therefore, understanding the mechanisms and kinetics of coke formation that occur over Ni-based cermet anodes is of primary importance and will be helpful with respect to the development and optimization of anode materials for SOFCs that operate on methane. The formation of coke over nickel-based catalysts during the partial oxidation or the steam and CO2 reforming of methane or other hydrocarbons has been extensively investigated during the past decades.108−112 Coke accumulation has been suggested to be a more accurate description than coke formation because a balance between the formation and the removal of coke determines the degree of coking.113 For convenience, we use the term “coke formation” throughout this Review. Although coke formation on nickel surfaces has been fairly well investigated, not all aspects of the process are completely clear due to its complexity. In general, hydrocarbons (including methane) are believed to first dissociate to produce highly reactive monatomic carbon (Cα),114,115 which is easily gasified through reaction with the oxidant to form CO. However, the excessive formation of Cα, which is nongasified, could polymerize to Cβ, which is much less reactive than Cα.113,116 Consequently, Cβ is easy to accumulate on the nickel surface or dissolve into the bulk of nickel particles. The carbon growth on the Ni catalysts is also generally accepted to occur via the dissolution−precipitation mechanism.117−119 For example, methane adsorbs onto the nickel surface and then decomposes into C and H; later, the carbon atoms dissolve into the bulk of the Ni particles, diffuse through the Ni particles, and finally precipitate as carbon at their rear side. The growth of carbon on the nickel surface during a hydrocarbon reforming process has been explored experimentally in detail by Sehested.120 Three types of carbon have been observed in a reformer: pyrolytic, encapsulating, and whisker carbons. The whisker carbon is one of the most destructive forms of carbon produced in the steam reforming of methane on Ni-based catalysts. The whisker carbon is formed by two steps: first, the hydrocarbons or CO decompose at one side of the nickel particle; second, the graphitized carbon nucleates with the formation of a carbon whisker on the other side of the nickel particle.120 The process is believed to begin by the formation of nickel carbide.117 The possible pathways for syngas production and for coke formation are shown in Figure 2,121 and illustrative TEM images of carbon nanofibers (whisker carbon) and graphite carbon, as observed experimentally, are also presented. On the basis of the dissolution−precipitation mechanism and the Langmuir−Hinshelwood mechanism, as reported by Asai et al.,122 the kinetic expression for the coke formation rate, v, was derived by Asai et al.:122

Figure 2. The pathways for syngas production and for coke formation. Reprinted with permission from ref 121. Copyright 2011 John Wiley and Sons.

v=

kMStotal(PM − C NiPH22 /K ) 1 + C Ni /KD + (KHPH2)1/2 + KMC NiPH22 /K

(15)

where kM and kM′ are rate constants for the adsorption and desorption of CH4; SV is the number of vacant sites; SM, SC, and SH are the numbers of surface sites occupied by CH4 molecules, C, and H atoms, respectively; CNi is the concentration of carbon in the nickel particles; KM (=kM/kM′), KH, KD, and KS are equilibrium constants for methane and hydrogen adsorptions, for carbon dissolution, and for the surface reaction; K = KMKDKS/KH2; and Stotal = SV + SM + SC + SH. The key assumption of the Langmuir−Hinshelwood mechanism is that the CH4 molecules, and C and H atoms were adsorbed onto the same sites. Therefore, the decomposition of one methane molecule requires five adsorption sites.119 Although the denominator in eq 15 is rather complicated, its second, third, and fourth terms are equal to SC/SV, SH/SV, and SM/SV, respectively. By assuming one of the four sites, vacant (SV), carbon (SC), hydrogen (SH), or methane (SM) sites, is predominant, the equations for expressing the coke formation rate (eq 15) reduces to simpler expressions, which are summarized below:123 methane: v =

kMStotalKDKSPM C NiKH 2PH22

vacant: v = kMStotalPM

hydrogen: v =

carbon: v =

(16) (17)

kMStotalPM (KHPH2)1/2

kMStotalKDPM C Ni

(18)

(19)

For simplicity, the second term of the numerator in eq 15 can be neglected. A comparison of the observed reaction orders with the kinetic expressions revealed that the rapid preequilibrium adsorption of methane was the rate-determining step at lower temperatures, whereas the dissolution of surface carbon atoms into nickel particles was the rate-determining step at high temperatures.123 3.2. Factors Controlling the Coke Formation

3.2.1. Thermodynamics. In principle, the coke formation can be avoided thermodynamically through the introduction of G


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atures. As shown in Figure 3b, the temperature region associated with the carbon deposition became narrower when the steam was added into the system with steam-to-carbon (S/ C) molar ratio of 1.0, whereas almost no solid carbon formed at an S/C ratio of 1.5 (Figure 3c). Similarly, the thermodynamic formation of coke as a function of temperature for the mixture of methane and steam can also be predicted. As shown in Figure 4, the amount of coke

a sufficient amount of steam or other oxygen-containing oxidant(s) into the fuel system. Sasaki and Teraoka have studied the equilibrium diagrams of methane within the temperature range of 100−1000 °C.124,125 As shown in Figure 3a, CH4(g) was stable at lower temperatures, whereas it decomposed to form H2(g) and solid carbon (C) at higher temperatures. The amount of solid carbon derived the thermal decomposition of methane increased with increasing temper-

Figure 4. Thermodynamically predicted coke formation toward methane as a function of temperature.

increased as the operating temperature increased at CH4-toH2O molar ratios less than 2:1. However, the coke formation is obviously reduced when the molar ratio of CH4 to H2O is reduced to 1:1. The coking is negligible at temperatures greater than 750 °C. If the CH4-to-H2O ratio is further reduced to 1:1.5, the coking is totally inhibited at all operation temperatures. 3.2.2. Nickel Particle Size. Numerous experimental and theoretical studies have confirmed that the nickel particle size significantly affects the coke formation properties of Ni-based catalysts.126−131 For example, Ermakova et al. have found that the coke formation by the decomposition of methane is related to the average particle size of nickel in the fresh catalysts, and they found that the most serious coke formation occurs at a nickel particle size between 20 and 60 nm.127,128 The strong effect of the Ni particle size on coke formation can be directly related to the thermodynamic equilibrium constant or the coking threshold.126,132 It was proposed that this can be related to the changes in thermodynamic properties of carbon nanofibers due to the external tension energy of carbon nanofibers as a function of the crystal size.126 The coking resistance of catalysts can thus be improved through tailoring the particle size of nickel, which could be implemented through the use of a promoter, changing the composition, or using different preparation methods for the catalysts. 3.2.2.1. Preparation Method. Most of the work reported in the literature involved the use of physical mixing, the impregnation method, or the sol−gel method for the preparation of nickel-based catalysts. Recently, some new preparation methods have been developed. These methods involve various techniques, such as combustion,133 supercritical treatment,134 microwave irradiation,135 and plasma treatment.136 The preparation method can strongly affect the nickel particle size and therefore result in catalysts with different catalytic activities and coking resistances. For example, Kim et al. have studied the coke formation behavior of Ni/Al2O3 catalysts prepared via impregnation and aerogel methods for

Figure 3. Equilibrium products from methane (CH4) mixed with H2O in different steam-to-carbon ratios (S/C): (a) S/C = 0, (b) S/C = 1.0, and (c) S/C = 1.5. Reprinted with permission from ref 124. Copyright 2003 The Electrochemical Society. H


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the CO2 reforming of methane.137 As shown in Figures 5 and 6, the preparation methods affected both the morphology and the

the CO2 reforming of methane.139 Enger et al. have also reported that the inactive NiAl2O4 spinel could be a catalyst precursor for the steam reforming of methane, and they prepared nanosized nickel particles (15−25 nm) with excellent initial stability (for a period of 20−40 h) by exposing the NiAl2O4 spinel to hydrogen at 800 °C for 2 h.141 Good catalytic activity toward the partial oxidation of methane was also achieved with NiAl2O4 after it was fully reduced.142 3.2.2.2. Nickel Content. Considerable research has been conducted to understand the effect of nickel content on the catalytic activity and the coke formation behavior of Ni-based catalysts in the partial oxidation, steam, and CO2 reforming of methane.138,143−146 Because nickel is the active component for the electrochemical and chemical oxidation of methane, a proper amount of nickel is required to achieve sufficient catalytic activity. However, excessive nickel may lead to the aggregation of nickel particles and thus affect the activity and coking resistance of the catalyst. The best catalytic activity is normally achieved when the active metal forms a monolayer or monatomic dispersion over a support.147 Interestingly, under such conditions, the active metal content is often approximately 15 wt % in Ni-based catalysts.143,145 If the nickel content exceeds this limit, the nickel particles are expected to aggregate. For example, Dong et al. have studied the effect of Ni content on the performance of the Ni/Ce−ZrO2 catalyst for the steam reforming of methane,143 in which a Ni/Ce−ZrO2 catalyst with 15 wt % Ni exhibited not only the highest catalytic activity and selectivity but also remarkable stability; by contrast, catalysts with Ni contents greater than 15 wt % suffered from serious coke formation. Similar results were also demonstrated by Hao et al.138 and Wang et al.145 for Ni/Al2O3 catalysts in the CO2 reforming of methane. 3.2.2.3. Nickel Sintering. Sintering describes a process in the fabrication of heterogeneous catalysts where small particles grow into larger particles. Because large nickel particles are prone to coke formation, the sintering of nickel catalysts becomes an important factor.120,148 Prevention of the sintering of nickel particles is therefore another essential approach to suppress coke formation. Two mechanisms are known for the growth of metal particles:120 (i) particle migration, where entire crystallites migrate over the support, followed by amalgamation; and (ii) Ostwald ripening, where metal transport species are emitted from one crystallite, then migrate over the support or via the gas phase and are finally captured by another crystallite. One way to suppress the sintering of nickel is to introduce another metal, such as Cr, to the Ni catalyst such that the metal does not form an alloy with Ni. The foreign metal acts as an inhibitor for the Ni sintering and enables good dispersion of the nickel phase.149 3.2.3. Alloying Effect. The electronic structure of carbon is very similar to those of tetra- and pentavalent p metals, such as Ge, Sn, and Pb, which contain “spare” p electrons in their outer shell, close to a stable s-orbital. With the presence of carbon, nickel carbide may be formed due to the interaction between the 2p electrons of the carbon and the 3d electrons of the nickel. A potential way to reduce coke formation is to add another metal, which can interact with the 3d electrons of Ni to form an alloy with nickel, thereby reducing the nickel carbide formation kinetically.113,150 The formation of an alloy to reduce the carbide formation is a potential solution to the coke formation issue; however, the loss of the overall catalytic properties of nickel is undesirable because the steam reforming of methane must be catalyzed. Thus, the selection of metals

Figure 5. TEM micrographs of (a) impregnated Ni/Al2O3 catalyst and (b) aerogel Ni/Al2O3 catalyst after reduction at 973 K. Reprinted with permission from ref 137. Copyright 2000 Elsevier.

Figure 6. Ni particle size distributions derived from TEM for (a) impregnated Ni/Al2O3 catalyst and (b) aerogel Ni/Al2O3 catalyst after reduction at 973 K. Reprinted with permission from ref 137. Copyright 2000 Elsevier.

size of the nickel particles distributed over the support. The impregnated catalyst contained numerous large Ni particles with a broad size distribution, as well as irregular-shaped particles deposited over the Al2O3 support. By contrast, small Ni particles with an average diameter of approximately 3.3 nm and with a narrow size distribution were observed in the aerogel catalyst. Large amounts of whisker carbon were observed on the impregnated catalyst, whereas small amounts of filamentous carbon were detected on the aerogel catalyst. Hao et al. have also reported similar results for Ni/Al2O3 catalysts for the CO2 reforming of methane,138 and they demonstrated that the nickel particle size was closely related to the preparation method, the nickel content, and the sintering behavior of the active nickel metal. The preparation method can also affect the catalytic activity, stability, and coking resistance of the catalysts by altering the interaction between nickel and the support. The NiAl2O4 spinel, which is not readily reduced, is well-known to form in catalysts prepared using special preparation techniques. For example, NiAl2O4 was formed during the preparation of a Ni/ Al2O3 catalyst via a combustion technique with urea as fuel.139 Although the formation of NiAl2O4 may reduce the amount of active nickel species and consequently caused a low activity for the reforming reactions,140 it improved the coking resistance of the catalysts for the CO2 reforming of methane.138 For example, homogeneous nickel nanocrystallites with an average particle size of 4 nm were produced in the Ni/Al2O3 catalyst prepared via the combustion method with urea as fuel, and the catalyst showed an excellent catalytic activity and stability for I


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formation of surface carbonate and bicarbonate species. These species provided the possibility of improving the coke resistance of Ni/SiO2.170 3.2.5. Gas Composition. Watwe et al.171 and Bengaard et al.172 have reported that the methylidyne dehydrogenation (i.e., CH* → C* + H*) step is the rate-limiting step for methane dehydrogenation because this step exhibits the highest activation energy. The presence of hydrogen in the feed gas will increase the concentration of adsorbed hydrogen species and prevent this dissociation, thereby suppressing coke formation. Numerous researchers have focused on the addition of co-fed hydrogen to suppress coke formation and/or methane selectivity in the steam reforming of (oxygenated) hydrocarbon reactions.173−176 For example, Laosiripojana et al.173 have studied the effect of hydrogen addition on the coke formation behavior of Ni/CeO2 and Ni/Al2O3 catalysts for the steam reforming of ethanol. They observed that the yields of CH4, C2H4, and C2H6 decreased as the molar ratio of H2/C2H5OH increased, until it reached 3.0; the effect of the hydrogen feed became less significant at higher H2/C2H5OH ratios, which resulted in a slightly enhanced CH4 concentration. With the presence of hydrogen, less carbon was observed to form over both high-surface-area Ni/CeO2 and Ni/Al2O3 catalysts. However, excessive amounts of hydrogen slightly decreased the catalytic activity for both catalysts because of the competition between active sites on nickel particles. Furthermore, in the case of the Ni/CeO2 catalyst, the decrease in the catalytic activity as the amount of hydrogen was increased was also due to the inhibition of the reactions between hydrocarbons and the lattice oxygen (OOx) on the surface of CeO2. Figure 7 presents the critical hydrogen content necessary to avoid coke formation during the methane cracking reaction

that will form an alloy without decreasing the catalytic activity of the nickel catalyst is limited. For example, Sn−Ni alloy catalysts are widely used in the steam reforming of methane and other hydrocarbons.151,152 It is demonstrated that the coke resistance of Ni can be improved by Sn−Ni surface alloys, by preferentially oxidizing the carbon atoms rather than form C− C bonds, and have a lower thermodynamic driving force, associated with the nucleation of carbon atoms on lowcoordinated Ni sites.151 Using the molecular insights obtained in the DFT calculations, the authors have identified Sn−Ni surface alloy as a potential coking-resistant reforming catalyst. The Dacapo pseudopotentials plane wave code was employed for all of the calculations. The Generalized Gradient Approximation (GGA)-PW91 functional was employed for self-consistent spin-polarized electronic structure calculations. The predictions of the DFT calculations were supported by the experimental studies, which showed that Sn−Ni is much more resistant to coke formation than Ni in the steam reforming of methane, propane, and isooctane at moderate steam-to-carbon ratios.151,152 3.2.4. Second Phase(s). The catalytic activity and coking resistance of Ni-based catalysts are determined not only by the active nickel phase but also by the support(s) and promoter(s). The supports and/or promoters play a major role in providing thermal stability and participating in the coke removal reaction. Carbon deposition on a catalyst is closely related to its surface structure and acidity.153 Basic additives or promoters that favor water adsorption and OH surface mobility can decrease the coke formation rate on catalyst surfaces. Alkaline earth oxides (i.e., MgO and CaO) have been widely used as additives in catalysts for the reforming reactions to neutralize the acidity of the Al2O3 support.154−157 Oxycarbonates were found to be the main intermediate on La2O3-supported Ni catalysts, and these intermediates suppressed the coke formation.158 In addition, basic oxides can further enhance coking resistance by promoting the reaction of steam/CO2 and solid carbon.113,121 With respect to the promoters, emphasis has been placed on the role of alkaline or alkaline-earth oxides in modifying the coke formation behavior of Ni-based catalysts.159−163 Horiuchi et al. have studied the effect of basic metal oxides on the catalytic activity and carbon deposition properties of Ni/Al2O3 catalysts for the CO2 reforming of methane.159 They revealed that the basic metal oxides affected both the catalytic activity and the coke formation rate of the catalysts. Their results indicated that the surface of the Ni/Al2O3 catalysts with basic metal oxides contained an abundance of adsorbed CO2, whereas the surface of catalysts without basic metal oxides contained an abundance of adsorbed CH4; thus, basic metal oxides limited the occurrence of the CH4 decomposition reaction. Recently, researchers have shown increased interest in the promotion effect of rare-earth oxides, which have been revealed to slightly enhance the catalytic activity toward steam reforming and to improve the coking resistance of catalysts by enhancing the carbon gasification rate.164−168 Yang et al. have reported La2O3 and CeO2 co-promoted Ni/Al2O3 catalysts for the CO2 reforming of methane. Not only was the amount of deposited carbon decreased, but the catalytic activity was also slightly improved through the use of La2O3−CeO2 as a promoter in the Ni/γ-Al2O3 catalyst.166 In addition to the previously mentioned oxides, other promoters have also been investigated. For example, Pan et al. used a Ga2O3-doped Ni/ SiO2 catalyst for the CO2 reforming of methane;169 CO2 was activated on the Ga2O3-doped catalyst, which resulted in the

Figure 7. Thermodynamically predicted critical hydrogen content in methane to avoid coke formation and the related hydrogen to methane ratios as a function of temperature.

(CH4 → C + 2H2); this hydrogen content was obtained from thermodynamic calculations in which the critical hydrogen amount and the related hydrogen-to-methane ratios were calculated according to the equations: lg K = −

K= J

Δr Gm 2.303RT

(20)

2 4XCH 4

1 − XCH4

(21) dx.doi.org/10.1021/cr300491e | Chem. Rev. XXXX, XXX, XXX−XXX

Chemical Reviews C H2 =

Review

2XCH4 1 + XCH4

R H2 /CH4 =

ionic conductivity, high chemical and thermal compatibility, a good thermal expansion match with other cell components, and high electrical conductivity; thus state-of-the-art anode materials are always composed of interpenetrated nickel and electrolyte materials (doped ceria or stabilized zirconia). The nickel content in the reduced composite anode is typically in the range of 44−65 wt %, which is significantly higher than the 1−25 wt % in the typical catalysts for methane steam reforming. In addition, the nickel in the cermet anode is easily sintered at elevated temperatures of 1300−1500 °C; thus, the nickel particles in the anode are typically large, which implies that, under methane-rich conditions, the coking should be even more serious in SOFCs with a nickel cermet anode than in those with normal nickel-based catalysts. The factors controlling the coke formation of Ni-based catalysts could also be applied to the development of coking-resistant Ni-based anodes for SOFC. Considerable research has been conducted on the coke formation behavior of conventional Ni-based anodes that operate directly on methane fuel.179−183 For example, He et al. have studied the carbon deposition behavior of Ni−yttriastabilized zirconia (YSZ) anodes after they were exposed to 3 vol % steam humidified methane at temperatures between 500 and 800 °C.182 Figure 10 shows the scanning electron microscope (SEM) images of the fresh as-prepared Ni−YSZ anodes and the anodes treated in the humidified methane atmosphere for 4 h at various temperatures. After the anodes were exposed to the humidified methane at temperatures greater than 700 °C, carbon nanotubes were observed on the surface of the nickel particles. However, the deposited carbon was dissolved into the bulk of the Ni particles and destroyed the anode structure in all of the as-prepared Ni−YSZ anodes (Figure 10f) when the temperature was increased to 800 °C. After the carbon deposition test, the pellet size increased, and higher reaction temperatures resulted in larger pellet sizes. After the oxygen-temperature programmed oxidation (O2-TPO) process was performed to remove the deposited carbon, the Ni structure was damaged to various degrees, depending on the exposure temperature. The Ni−YSZ anodes treated at different temperatures all showed irreversibility during their recovery, where higher temperatures led to more extensive damage. In fuel cells, the carbon could also fill the pores of the electrode and lead to increased mass-transport polarization resistance, consequently, diminished cell performance. In addition to nickel, the ceramic phase in the anode also affects the coke formation behavior. For example, Sumi et al. have studied the effects of the crystal structure of YSZ and scandia-stabilized zirconia (ScSZ) in nickel-based SOFC anodes on carbon deposition and the oxidation behavior for methane fuel.183 They observed that the lattice parameter of zirconia decreased after 1−2 mol % of nickel was dissolved into the YSZ and ScSZ lattices. The lattice parameter and crystal structure of Ni-doped YSZ did not change during the redox treatment, whereas the lattice parameter of the Ni-doped ScSZ increased after the reduction treatment; these results suggest that the nickel solubility in ScSZ was decreased by a partial change in the crystal structure. New, finer Ni particles were formed around the original Ni grains, accompanied by a decrease in Ni solubility in ScSZ after the reduction treatment. In addition, rod-shaped carbon was observed to grow from the new, finer Ni particles on the ScSZ substrate. The morphology and crystallinity of the deposited carbon were strongly affected by the nickel particle size, which was dependent on the nickel

(22)

2XCH4 1 − XCH4

(23)

where K is the equilibrium constant for the methane cracking reaction; ΔrGm is the Gibbs free energy; R is the universal gas constant (8.314 J mol−1 K−1); XCH4 is the methane conversion; XH2 is the hydrogen content; and XH2/CH4 is the ratio of hydrogen to methane when the reaction reached equilibrium at a certain temperature. According to Figure 7, the hydrogen-tomethane ratio should be greater than 12 to thermodynamically avoid carbon deposition at 800 °C; this ratio is significantly higher than the ratio determined experimentally and reported in the literature.177 This discrepancy suggests that kinetics plays a more important role than thermodynamics for the suppression of coke formation, as reported by Laosiripojana et al., under conditions where a small amount of hydrogen was added.173 Similar results were also reported by Eguchi et al.178 according to their equilibrium phase diagram of methane, hydrogen, and carbon, which is shown in Figure 8. Shown in Figure 9 are the schematics of the high-level summary of section 3.

Figure 8. The boundary of carbon deposition region in the C−H−O phase diagram at 1 atm. Reprinted with permission from ref 178. Copyright 2002 Elsevier.

Figure 9. The schematics of the high-level summaries of section 3.

3.3. Coke Formation over Sintered SOFC Nickel Cermet Anodes

As compared to typical methane reforming catalysts, the anodes of SOFCs require some additional properties, such as sufficient K


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Figure 10. SEM images of Ni−YSZ (a) as-prepared, after carbon deposition in humidified methane for 4 h at (b) 500 °C, (c) 600 °C, and on the surface at (d) 600 °C, (e) 700 °C, and (f) 800 °C. All images are from the center of the pellet except for (c). Reprinted with permission from ref 182. Copyright 2007 Elsevier.

could improve the steam reforming, WGS, and carbon-removal reactions. For the Ni−ScSZ cermet, the decreased carbon deposition with increased temperatures, which was different from thermodynamic predictions, could be explained by an actual local O/C ratio on Ni (referred as (Oad/C)Ni) that was significantly greater than the O/C in the fuel atmosphere due to the enrichment of H2Oad on the Ni surface. However, a highly enriched H2Oad content was not formed on Ni−YSZ, which resulted in a coke formation trend similar to the thermodynamic prediction. The authors believed that the difference in (Oad/C)Ni between Ni−YSZ and Ni−ScSZ was attributable to different effects of the oxide in the ceramic phase on the H2Oad accumulation on the nickel surface. Similar to the coke formation over conventional nickel-based catalysts, the coke formation over Ni-based anodes may also be affected by the preparation methods. Asamoto et al. fabricated a Ni−Sm2O3-doped CeO2 (SDC) anode via an electrophoretic deposition (EPD) technique for the direct oxidation of dry methane in SOFCs.186 The effects of the Ni content, the heattreatment temperature, the supporting materials, and the anode layer thickness on the electrochemical performance were investigated. The results revealed that the power density and anode overpotential were influenced by all of the previously mentioned factors. The optimal fabrication conditions were a Ni content of 20 wt %, a heat-treatment temperature of 900 °C, SDC as the support material, and an anode thickness of 20 μm. The electrochemical performance of a Ni(20)−SDC (via EPD)

solubility in zirconia. At high temperatures, amorphous carbon was easily deposited over the Ni−YSZ cermet, which caused the rapid deterioration of the anode performance. A comparative study of the performance and durability of Ni−YSZ and Ni−ScSZ anodes for SOFCs operated on internal steam/CO2-reforming of methane has been reported by Sumi et al.184 The coke formation rate on Ni−YSZ was faster than that on Ni−ScSZ at 1000 °C, whereas the opposite trend was observed at 850 °C. These results suggest that the coke formation rates were affected by both the dopants in the zirconia and the operation temperatures. These phenomena can be explained by the change in the crystal structure of ScSZ under different heat-treatment conditions, which was in good agreement with the results reported by Gunji et al.185 Ke et al. have studied the effect of the ceramic phase in the Ni-zirconia cermets on the carbon deposition behavior toward the internal reforming of methane fuel in SOFCs.181 Two types of Ni-based anodes, that is, Ni−ScSZ and Ni−YSZ, were selected. Within the temperature range of 800−1000 °C, carbon deposition on the Ni−ScSZ cermet anode decreased as the temperature was increased at a low oxygen/carbon ratio (O/C = 0.03) under a certain polarization current density. The opposite trend was observed with the Ni−YSZ anode under the same conditions. The authors considered that the H2O produced from hydrogen electrochemical oxidation in the adsorption form was accumulated on the surface of the TPB region and that the adsorbed H2O (H2Oad) on the Ni surface L


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Figure 11. TEM image of YSZ particle following ramped reduction in hydrogen at 700 °C for 1 h (a) and following steam reforming of methane at 700 °C for 65 h at S/C = 3 (b). Small particles are pure Ni. Reprinted with permission modified from ref 192. Copyright 2008 Elsevier.

methane−steam gas mixtures as a direct feed for a SOFC with a Ni−YSZ anode and found that no coking was observed at S/C = 3 at temperatures greater than 950 °C;189 in contrast, only 10 vol % H2O addition was reported to be sufficient to fully suppress carbon deposition over a CeO2-modified Ni−YSZ anode operated under OCV conditions at 850 °C on methane fuel.187 Such a large discrepancy can be explained by the fact that thermodynamic calculations only provide insight into the expected equilibrium conditions and do not account for the impact of reaction kinetics.190 In the thermodynamic calculations, a significant assumption is that the rates of the carbon deposition and carbon removal processes are sufficiently rapid to achieve equilibrium; however, under real reaction conditions, this situation may not be the case. In some cases, the microstructure that was assigned to reaction kinetics may significantly affect the coking resistance of the electrode. For example, Chen et al. attributed the attractive coking resistance of a CeO2-modified Ni−YSZ anode to its novel microstructure, where the Ni and YSZ particles were homogenously covered with a layer of 100 nm CeO2 particles.187 Ceria is well-known to have a sufficiently high oxygen storage capacity (OSC) to store oxygen under an oxidative atmosphere and release it under a reductive atmosphere;191 thus, the addition of ceria could increase the O/C ratio in the anode and decrease the critical amount of steam in the methane fuel required to avoid coke formation. Although the introduction of steam into fuel gas can reduce coke formation thermodynamically, it can also detrimentally affect the operational stability of the anode by facilitating the sintering of nickel particles, which, in turn, affects the catalytic activity and coking behavior of the anodes.192,193 For example, King et al. found that methane conversion over a Ni−YSZ cermet anode catalyst under steam reforming conditions was significantly reduced from an initial value of 78% to 9% after the fuel cell was operated for 120 h at an S/C ratio of 3.0 at 700 °C.192 Two contributions of nickel particles to the catalytic activity were detected: large bulk nickel particles derived from the source material and small Ni crystallites located at the surface of the YSZ particles. NiO was dissolved into the YSZ to form a solid solution during the high-temperature sintering process and then evolved to form small Ni particles during the

anode was also compared to that of a Ni(20)−SDC (via slurry coating, SC) anode. The peak power density (PPD) of the SOFC with the anode prepared using the EPD technique was significantly greater than that with the anode prepared using the SC technique. In addition, the SOFCs with the Ni−SDC (via EPD) anode also exhibited significantly better cell stability and coking-resistance than those with the Ni−SDC (via SC) anode. 3.4. Strategies To Suppress Coke Formation over Nickel Cermet Anodes

As previously demonstrated, although Ni-based cermets are effective anodes for the direct utilization of methane in SOFCs in principle, they suffer from numerous drawbacks, such as their tendency to carbon deposition and their poor redox stability. Several strategies have been extensively investigated to increase the coking resistance of nickel cermet anodes, including (1) the addition of steam, CO2, O2, or H2 into the fuel gas to increase the O/C ratio and thereby avoid the coke formation thermodynamically, (2) the promotion of the carbon elimination reaction by increasing the polarization current, (3) the introduction of another metal(s) into the nickel cermet anode to modify the surface of the nickel particles or form an alloy−ceramic anode, (4) the application of an anode catalyst layer to alter the gas distribution within the anode and to increase the catalytic activity for the reforming/partial oxidation of methane, (5) the tailoring of the ceramic phase in the anode, and (6) the modification of the anode surface with other active oxide(s). 3.4.1. Addition of Other Gases. 3.4.1.1. Steam/Oxidant. The simplest way to suppress coke formation over a fuel cell anode that operates on methane fuel is via a thermodynamic approach that involves the addition of oxygen-containing gases into the fuel gas to increase the O/C ratio. For example, steam has been added to the methane fuel to reduce coking over Nibased anodes of SOFCs via internal steam reforming.187−189 S/ C ratios greater than the thermodynamic minimum have been demonstrated to be necessary to avoid carbon deposition and structural damage to the anode.189 In some cases, however, the actual S/C ratio for effectively avoiding coke formation was drastically different from that predicted from the thermodynamic calculations. For example, Laosiripojana et al. used M


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highly effective way to suppress coke formation because the hydrogen-to-methane ratio must be large to totally avoid coke formation thermodynamically, as shown in Figures 7 and 8. In addition, this approach may decrease the efficiency of the fuel cell and increase reliance on the hydrogen economy. 3.4.2. Polarization Current. For fuel cells under polarization, oxygen in the form of O2− transports from the cathode to the anode, where it reacts with the fuel. As the polarization current increases, the flux of O2− across the electrolyte from the cathode to the anode increases. However, the conversion of carbonaceous species, such as methane, into CO, CO2, H2, and H2O has also been shown to increase with the polarization current,197 which implies that the partial oxidation or steam/ CO2 reforming of methane occurs over the anode during the current polarization. Several researchers have successfully demonstrated the effectiveness in suppressing coke formation through the application of a polarization current, both theoretically and experimentally.180,198−201 For example, Koh et al. have used thermodynamic chemical equilibrium to study the potential reactions of methane in a SOFC and found that a critical current density (Jc) greater than 85 mA cm−2 could inhibit coke formation.180 The main assumption made in that study was that the carbon reacted with O2− as soon as it formed. This method could be used to predict a threshold current density for any given set of conditions where coke formation was mitigated by direct oxidation. However, the Jc values reported by different researchers are quite different from each other: they vary from 0.05 to 1.8 A cm−2 for different fuel cell systems and operating temperatures. The discrepancy between these results reflects the complexity of the reactions over the fuel cell anodes. AlzateRestrepo and Hill reported that the amount of deposited carbon on Ni−YSZ exposed to methane was affected by the current density in a fuel cell operated at 800 °C.198 The amount of deposited carbon decreased as the current density was increased from 1 to 50 mA cm−2. The coke formed under polarization-current conditions remained on the Ni surface and slightly affected the anode microstructure, whereas the carbon formed under OCV conditions dissolved into the bulk of the Ni structure. Lin et al. have also the stability of a cell operated on methane fuel at different current densities and temperatures.199 At 650 and 700 °C, the stability under methane was excellent at a small current density of 0.1 A cm−2, which suggests that the Jc at 650 and 700 °C was 0 < Jc < 0.1 A cm−2. At 750 and 800 °C, much larger critical current values, such as 1.2 A cm−2 for 750 °C and 1.8 A cm−2 for 800 °C, were required to maintain stable operation. These results imply that the O2− flux through the SOFC was at least partly responsible for preventing coking and thereby for maintaining stable operation. Liu et al. have reported that SOFCs with conventional Ni−YSZ anodes can be successfully operated on 3 vol % steam humidified methane and natural gas with high OCVs and power densities.200 The SOFCs were stably operated on methane fuel for more than 90 h at 700 °C and 0.6 V, which yielded a PPD of 0.35 W cm−2. Very little carbon was detected on the anodes, which suggests that carbon deposition was limited under the high polarization current density. Horita et al. developed a model for the elimination of deposited carbon over Ni-based anodes under polarization current conditions.201 They compared the catalytic activities of Ni-mesh/YSZ samples for CH4 decomposition and the reactions with reformed gases under polarization current conditions through image analysis of labeled gases with

hydrogen reduction. The significantly decreased methane conversion was attributed to the sintering of both types of Ni particles, as shown in Figure 11. Thus, care should be taken in suppressing coke formation by simply increasing the steam content in the fuel gas. Some researchers also attempted to suppress the coke formation over nickel cermet anodes by adding O2 or CO2 into the methane fuel gas; however, progress in this area remains slow.194−196 The challenge is that the amount of oxidant added should be kept to a minimum because an increase in the amounts of such oxidants may increase the polarization resistances and reduce the power density; furthermore, an excessive amount of oxygen also increases the risk of reoxidation of the Ni-based anodes, which results in an irreversible effect on the operational stability of the cell. For example, Wang et al. observed oscillatory behavior of Ni−YSZ anodes in a single-chamber SOFC (SC-SOFC) under methane oxidation conditions. They attributed the oscillations in the anode resistance, the cell voltage, and the actual temperature of the fuel cell to the Ni/NiO redox cycles.194 The corresponding oscillation patterns were mainly dependent on the methane-tooxygen ratio. Greater current densities promoted Ni oxidation, and Ni/NiO redox cycles occurred primarily near the anode surface. The authors hypothesized that the gradual reoxidation of the Ni-based anodes that accompanied the various oscillation behaviors played an important role in the degradation of the SC-SOFCs. Wang et al. have studied the effect of the addition of H2O, CO2, or O2 to methane on the power output at 850 °C of fuel cells with Ni−ScSZ anodes.142 They found that the PPDs of the fuel cells operated on all three gas mixtures first increased as the methane-to-O2/H2O/CO2 ratios were increased and subsequently decreased. The maximum PPDs were achieved at methane-to-O 2 /H 2 O/CO 2 ratios of 8:1, 2:1, and 4:1, respectively. Excessively low methane-to-O2/H2O/CO2 ratios resulted in high concentrations of unconverted O2, H2O, and CO2 due to the poor catalytic activity of the anode, and these unconverted gases subsequently diluted the H2/CO fuels and resulted in a lower H2/CO concentration in the fuel gas. However, excessively high methane-to-O2/H2O/CO2 ratios also resulted in decreased H2 and CO concentrations in the fuel gas due to insufficient H2/CO production. These results demonstrate that the sintered Ni-based anodes exhibited extremely low catalytic activity for the partial oxidation, steam, and CO2 reforming of methane and consequently resulted in low power outputs. Thus, the sintered Ni-based cermet anodes are inappropriate for SOFCs that operate on methane through internal partial oxidation/reforming; instead, proper modification of the anode is required. 3.4.1.2. Hydrogen. The introduction of hydrogen into the fuel gas to suppress coke formation over the anodes was also investigated in methane-fueled SOFCs, although the literature contains few such reports. As an example, Nikooyeh et al. have studied the effect of hydrogen addition on coke formation over Ni−YSZ anodes exposed to methane fuel.177 They found that the coking was effectively suppressed at 800 °C under OCV conditions at H2-to-CH4 molar ratios of 0.5 to 1.5. The coke formation rate and its damage to the anode structure were significantly reduced as the hydrogen content in the feed gas was increased. Even if coke formation was not completely avoided, the addition of hydrogen could result in less damage to the anode microstructure after the carbon was removed. However, the addition of hydrogen to methane fuel gas is not a N


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decorate the surface of a SOFC anode cermet. For example, the modification of catalytic activity against coke formation can be achieved through the decoration of a small quantity of another metal on the surface of Ni particles in Ni-based anodes. Among the various metal additives, Au is considered to be an effective promoter for improving the coking resistance of Ni-based anodes toward methane.202−204 In a theoretical study, Besenbacher et al. confirmed that the presence of small amounts of Au in a Ni-based catalyst can significantly affect the coke formation behavior during the steam reforming of methane.205 Rostrup-Nielsen et al. have also predicted, using DFT calculations, that Au effectively inhibits the decomposition of methane.202 Triantafyllopoulos and Neophytides have studied the nature of carbon species formed on a 1 mol % Au modified Ni−YSZ catalyst during the dissociative adsorption of methane and confirmed the previous statements.206 Gavrielatos et al. have reported a carbon-tolerant Ni−YSZ anode modified with 1 mol % Au with respect to Ni for SOFCs under methane-rich steam reforming conditions.203 Figure 13a shows the surface morphology of the NiAu−YSZ anode, which was porous and consisted of nanoparticles with size of 30−40 nm, before the reduction process. The reduction caused the formation of larger nickel particles with sizes in the range of 100−200 nm on the catalyst surface (Figure 13b). As shown in Figure 13c, large nickel particles were observed on the catalyst surface, whereas the Ni particles in the bulk were still nanostructured. In addition, the Ni−YSZ anode underwent an immediate and rapid weight increase due to carbon accumulation on the Ni surface, whereas no weight increase of the NiAu−YSZ sample was observed at 650 °C. These results indicate that the addition of Au inhibited the dissociative adsorption of methane or at least the dehydrogenation reaction step that led to coke formation. At temperatures greater than 650 °C, carbon deposition still occurred on the NiAu−YSZ anode. In any event, the NiAu−YSZ anode showed good stability in a methane-rich atmosphere (CH4/H2O > 3) at 850 °C under a certain voltage for 60 h. The cell delivered the same

secondary ion mass spectrometry (SIMS). A significant amount of deposited carbon was observed on the Ni-mesh during a labeling experiment with a mixture of CH4, D2O, and 18O2 at 800 °C. However, the applied polarization voltage could effectively eliminate the carbon deposited on the Ni mesh. The authors further proposed a possible reaction mechanism, which is shown in Figure 12, for the optimal metal−oxide interface of

Figure 12. Models of CH4 decomposition and elimination of carbon deposition by the applied voltage over Ni−YSZ interface. Reprinted with permission from ref 201. Copyright 2005 Elsevier.

SOFCs. Under OCV conditions, carbon deposition occurred on the Ni surface, whereas with the application of a polarization current, the O2− conducted through the YSZ electrolyte could spill over from the Ni−YSZ interface to the Ni surface, and, consequently, the supplied O2− could effectively react with the deposited carbon on the Ni surface. The deposited carbon on the Ni surface was believed to be totally eliminated if the amounts of O2− supplied were appropriately controlled. 3.4.3. Decoration of the Electrode Surface or Alloying of the Nickel with Other Metal(s). 3.4.3.1. Surface Modification. Carbon deposition can also be reduced through the introduction of another metal as a separate phase to

Figure 13. Scanning electron micrographs of NiAu (1 mol %)−YSZ electrode. (a) Top side before its reduction with H2. (b) Top side after its reduction with H2 for 2 h. (c) Cross section and surface|gas interface, (d) cross section and NiAu−YSZ|YSZ interface. Reprinted with permission from ref 203. Copyright 2008 Elsevier. O


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inhibiting the deposition of C. The formation of a surface alloy, in which the alloy element is enriched at the topmost surface, was found to be critical to the activity of the Ni-alloy catalysts. Jia et al. have also investigated the adsorption of C on Ni (111), Cu (111), and alloyed Ni−Cu (111) surfaces using DFT calculations.216 All of the calculations were performed by using the Projector Augmented Wave (PAW) pseudopotentials as implemented in VASP. The exchange correlation functional was treated within the GGA and parametrized by PBE functional. They confirmed that C energetically favors the hollow sites of Ni (111) and Cu (111) surfaces. Because less overlap occurs between the C 2p and the metallic 3d orbits, the adsorption energy of C could be lowered through the formation of a Ni− Cu alloy with the addition of Cu into Ni. Wang et al. have designed a copper-modified Ni−SDC anode by impregnating Cu into a preprepared Ni−SDC porous matrix, which was referred to as CuNi−SDC. 207 By comparison, Ni−SDC and Ni0.95Cu0.05(alloy)−SDC were also prepared using a similar procedure without impregnation. The authors found that Cu particles were distributed uniformly on the porous Ni−SDC matrix and that the particle size of Cu ranged from 50 to 250 nm, as shown in Figure 14. The higher-

power output when operated on hydrogen fuel before and after the durability test on methane−steam gas mixtures, which further demonstrated the high tolerance of the NiAu−YSZ anode to carbon deposition under internal methane-rich steam reforming conditions. Niakolas et al. have also reported an Aupromoted Ni−GDC anode for SOFCs operated on methane through internal steam reforming.204 They used a deposition− precipitation method to add fine, dispersed Au nanoparticles to Ni−GDC anodes and also studied the effect of the Au content and calcination temperature on the carbon deposition behavior of the anodes. Au dispersed homogeneously on the Ni−GDC anodes, and the Au crystallite size varied in the range of 10−50 nm in samples calcined at 850 °C. The Au crystallite size increased to 150 nm in the sample with 4 mol % Au calcined at 1100 °C. Obviously, the addition of Au nanoparticles had a strong positive effect in suppressing carbon deposition under severe conditions (10 vol % CH4/Ar at 600 °C). As the Au content was increased, the amount of deposited carbon decreased dramatically, and almost no carbon was deposited on the 4 mol % AuNi−GDC. Furthermore, an increase in the calcination temperature was found to have an additional positive effect on the coking resistance. The authors attributed such phenomena to the increase in nickel particle size that occurred during the thermal treatment and to the agglomeration of nickel particles, which caused a reduction in the number of surface active sites for the methane cracking reaction. These results were incongruent with the fact that sintered nickel typically favors coke formation. This discrepancy can be attributed to the addition of Au. The authors explained the stable and enhanced performance of the Au-modified anode by the kinetic and electro-kinetic mechanisms reported by Triantafyllopoulos and Neophytides,206 where the syngas production proceeded through the decomposition of oxyhydrogenated species, which were formed by the oxidation of CHx species. However, the stability of the Au nanoparticles was not evaluated after the high-temperature thermal treatment or after the cell operation. Because nanosize particles may coarsen during long time operation, this could be a possible limitation of the approach. In addition to Au, other metals, such as Cu, Ru, and Rh, have also been used as additives for the modification of Ni-based anodes.207−213 Little research into the modifications with Ru or Rh has been reported because these metals are prohibitively expensive; however, in the published literature, they did not facilitate coke formation due to poor carbon solubility.214 In addition to Au, another extensively investigated metal additive is copper, which has been added to Ni-based anodes to reduce or prevent coke formation in SOFCs operated on methane fuel because copper is inert toward coke formation and exhibits high electronic conductivity. An et al. have presented DFT calculations of the chemisorption of the atomic species O, S, C, and H and the reaction intermediates OH, SH, and CHn (n = 1, 2, and 3) on bimetal M/Ni alloy model catalysts (M = Bi, Mo, Fe, Co, and Cu).215 All calculations were performed using DFT, as implemented in VASP with spin-polarized PBE functional. They found that the binding of undesirable intermediates such as C and S could be inhibited and that the catalytic activity of planar Ni-based anodes could be tuned toward oxidation through the selective formation of a bimetallic surface alloy. Cu/Ni, Fe/Ni, and Co/Ni anodes have been found to be the most active anodes toward oxidation, and the Mo/Ni alloy surface has been predicted to be the most effective in terms of

Figure 14. (a) SEM image of the porous Ni−SDC matrix after impregnation with Cu, (b) EDS mapping of the Cu distribution in (a), and (c) higher resolution of the SEM image in (a). Reprinted with permission from ref 207. Copyright 2008 Elsevier.

magnification micrograph (Figure 14c) showed that the impregnated Cu particles adhered well to the surface of the Ni−SDC matrix. The electronic conductivities of the Ni−SDC and Ni0.95Cu0.05−SDC anodes at 600 °C were approximately the same, whereas the conductivity of the CuNi−SDC anode was ∼50% higher. The PPDs for the fuel cells with Ni−SDC, Ni0.95Cu0.05−SDC, and CuNi−SDC anodes were 240, 338, and 317 mW cm−2, respectively, when operated on dry methane at 600 °C. The authors observed that the power output of the Ni−SDC anode supported cell was obviously lower than that of the fuel cells with Cu-modified anodes. Similar results were reported by Ringuedé et al.217 The addition of copper by P


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impregnation also dramatically improved the cell stability and reduced carbon deposition when the cells were operated on methane fuel.207 Although wet impregnation has been proven to be a useful method for the modification of SOFC anodes, it provides limited control over the microstructure of the deposited components. Multiple impregnation and subsequent calcination steps are often required to obtain a sufficient Cu content, which makes impregnation a time-consuming process.218 In principle, electrochemical deposition is another method that can be used for the addition of active components to a conducting porous electrode structure. This approach has the potential to allow better control of the morphology. Park et al. have used the direct current electroplating method to add copper to Ni−YSZ anodes to improve the coking resistance and operating stability with methane fuel.208 Figure 15 shows the apparatus for

The electrodeposition technique typically requires a long time (in terms of hours) for the deposition of Cu in an anodesupported SOFC.208 In addition, this method requires that the substrate material be conductive.219 Recently, microwave irradiation processes have been used for the synthesis of metal nanoparticles and metal oxides.220−222 This process is faster than most other processing techniques because the time for microwave irradiation may be on the order of seconds, and a subsequent calcination step to remove the precursor is not required. Islam et al. have used a microwave irradiation process to deposit Cu nanoparticles onto the Ni−YSZ anode of an electrolyte-supported SOFC.210 The authors considered the obvious advantage of the microwave irradiation method to be the significant reduction in the time required for anode preparation because the irradiation time was only 15 s and no calcination step was used. The spherical Cu particles prepared using this method were less than 100 nm in size. The stability of the fuel cell with this CuNi−YSZ anode was improved when run on dry methane at 800 °C under a certain current density. The modification of Ni-based anodes with a highly active nickel catalyst was also proven to be an effective way to improve cell performance. Zhang et al. impregnated a highly active nickel catalyst into a sintered Ni−ScSZ anode to improve the cell performance of a traditional SC-SOFC operated on methane−air gas mixtures.223 The nickel catalyst exhibited a catalytic activity for the partial oxidation of methane that was greater than that of the sintered Ni−ScSZ anode between 700 and 900 °C; the improvement was especially pronounced at the lower temperatures because of the nickel catalyst’s finer particle size and larger specific surface area. The impregnated nickel catalyst was found to increase the roughness of the nickel surface in the anode, which increased the metal surface area for the partial oxidation of methane. When operated on methane− air gas mixtures with a methane-to-oxygen ratio of 1.3:1, the fuel cell with the modified anode showed an OCV and a PPD of 0.954 V and 119 mW cm−2, respectively, at a furnace temperature of 750 °C; these values are significantly better than those of the unmodified anode (0.893 V and 79 mW cm−2). Coking-resistant surface decoration of Ni-based cermet anodes would be a promising future research direction, although the choice of alloying metals is limited. Au was used to modify the Ni-based anodes, and considerable improvement in coking resistance and operational stability is achieved. However, the stability in the microstructure of the Ni/Au anodes should be a concern. There are several methods to add Cu into the Ni-based anodes such as wet impregnation, electrochemical deposition, and microwave irradiation methods; however, both the wet impregnation and the electrochemical deposition methods have some limitations. The microwave irradiation seems to be an effective method to prepare the Cu modified anodes for SOFC; however, the longterm stability of these anodes was still lacking up to now. Because of the high catalytic activity of impregnated nickel and the excellent compatibility with the Ni-based anodes, it should be a good approach to improve the power output of the SOFC with Ni-based anodes, although the present literature is still few. 3.4.3.2. Alloying. Another way to avoid or suppress coke formation over SOFC anodes is the replacement of Ni in the conventional anodes with electronic conductors that do not catalyze carbon formation, such as copper. However, the Cubased cermets are limited to lower operating temperatures because copper is easily sintered at higher temperatures. The

Figure 15. Apparatus for Cu-electroplating. Reprinted with permission from ref 208. Copyright 2009 International Association of Hydrogen Energy.

electroplating of Cu onto the Ni−YSZ anodes with an aqueous solution. The Cu-electroplating was performed by controlling both the plating time and the Cu2+ concentration of the plating solution. Two parallel Cu meshes were used as reference and counter electrodes, and a magnetic bar was used to thoroughly mix the plating solution during the electroplating process. The addition of copper greatly decreased the amount of deposited carbon on the anode, especially for electrodes with longer Cuelectroplating exposure times. The authors found that the addition of copper dramatically improved the operating stability when methane was used as fuel due to the substantial suppression of carbon deposition, although the PPDs obtained from the CuNi−YSZ anode-supported single cell operated on hydrogen and methane were slightly lower than those of a fuel cell with a Ni−YSZ anode operated at 700 °C. Jung et al. have also reported the successful introduction of copper onto Ni− YSZ anodes through electrodeposition, which consequently improved the coking resistance and operating durability for a cell operated on methane. However, the exact amount of electroplated copper using this method was difficult to control.209 Q


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Ni4Fe1−ZrO2 (Ni:Fe = 4:1 by weight) prepared via a glycine nitrate process (GNP) showed somewhat less carbon deposition than a Ni−ZrO2 catalyst in methane.235 The Fe− ZrO2 itself did not inhibit the coke formation effectively. Surprisingly, Ni1Fe2−ZrO2 showed even more carbon deposition than Ni−ZrO2. Such discrepancies may be related to the different fabrication methods adopted by the different authors. Wu et al. have systematically studied the influence of preparation methods on the properties and cell performance of a Ni0.75Fe0.25−SDC (60:40 by weight) alloy−ceramic anode for SOFCs.236 Three methods, that is, a physical mixing route, a simultaneous GNP/sol−gel route, and a combined GNP sol− gel route, were used. The authors reported that the phase structure of the anode components, the chemical interaction between nickel and iron, and the anode microstructure were all strongly dependent on the synthesis method. The coking resistance was more sensitive to the anode phase structure and the chemical interaction between the Ni and Fe phases, whereas the cell power output was mainly determined by the microstructure. The authors suggested that the iron content in the NiFe-alloy-based anodes should be carefully controlled in the chosen preparation method to achieve high cell performance. Other metals, such as Pd and Sn, have also been investigated. Nabae et al. have reported a Pd−Ni/composite anode used for SOFCs that operate on methane and studied the alloying effects of Pd and Ni on the catalytic behavior of direct methane oxidation.237−239 The Ni, Pd, and Pd−Ni alloy were supported on La0.8Sr0.2CrO3 (LSCr) and SDC. The authors found that the fuel cell with a bare LSCr−SDC anode delivered the lowest PPD when operated on hydrogen fuel at 800 °C and that the PPDs of the fuel cells with the other three anodes were comparable. When operated on methane, the fuel cell with the Pd−Ni−LSCr−SDC anode exhibited a much greater power output than those with Pd−LSCr−SDC and Ni−LSCr−SDC anodes, which suggests that a synergistic effect occurred between the Pd and Ni catalysts for the direct oxidation of dry methane. PPDs of 150 and 420 mW cm−2 were obtained for the fuel cell with the Pd−Ni/composite anode on dry methane fuel at 800 and 900 °C, respectively. The amount of carbon deposited on the Pd−Ni/composite anode was small under both open- and closed-circuit conditions. The authors also presented a model of the reaction scheme for the anodic oxidation of CH4 over the Pd−Ni/composite anodes, as shown in Figure 16.237 The Pd−Ni catalyst on the composite anodes catalyzed the CH4 decomposition to H2 and C. The CH4

requirement for lower operating temperatures and the low catalytic activity for methane conversion make Cu-based cermets less effective for the direct utilization of methane.224,225 A possible solution to these problems is the use of anodes based on an alloyed system. Copper has reportedly been added to a Ni-based anode to form an alloy that suppresses coke formation toward hydrocarbons.226−229 Kim et al. have examined Cu−Ni alloys as anodes for the direct oxidation of methane in SOFCs at 800 °C.226 The anodes with different compositions of 0%, 10%, 20%, 50%, and 100% Ni were exposed to dry methane at 800 °C, and they found that coke formation on the Cu−Ni alloys was greatly suppressed as compared to that on pure Ni and that the increased reduction temperatures also reduced the coke formation rate. The Cu−Ni alloys all showed a certain degree of coke formation; however, the amount of carbon was clearly not proportional to the amount of Ni. A fuel cell with a Cu (80%)−Ni (20%) cermet anode showed a significant increase in power density with time for 500 h, which was attributed to improved electronic conductivity of the anode, as evidenced by the impedance spectra. Sin et al. reported a similar NiCu−GDC anode for the direct electro-oxidation of methane in intermediate-temperature (IT) SOFCs.227 A PPD of 320 mW cm−2 was achieved at 800 °C for an electrolyte-supported fuel cell with a NiCu−GDC anode operated on dry methane fuel. Excellent performance was obtained for approximately 1300 h of operation with dry methane and in the presence of redox cycles at 750 °C. However, the Cu−Ni alloys were found to be unstable in the presence of hydrocarbons, depending strongly on the pretreatment conditions.230 Consequently, investigations of anodes with other alloyed metals, such as Fe and Pd, are necessary. Among the various transition metals, iron has been the most extensively studied because of its low cost, abundance, and high effectiveness in improving the performance of nickel-based anodes. To improve the electrocatalytic activity for fuel oxidation, a proper amount of iron added into a nickel-based anode may also enhance the coking resistance in SOFCs operated on hydrocarbon fuels.231,232 For example, Kan et al. have reported an iron modified Ni−GDC anode for an SOFC operated on dry methane with considerable power output and enhanced operating stability.231 A fuel cell with a Ni0.9Fe0.1− GDC alloy-based anode provided the highest power density, which was even higher than that of the fuel cell with Ni−GDC anode. The addition of 10% Fe was observed to decrease the electrode polarization resistance, whereas the addition of iron to concentrations of 30% or 50% increased the anode polarization resistance. The addition of a proper amount of Fe into Ni−GDC anodes also significantly improved the longterm stability of SOFCs. For example, a fuel cell with a Ni− GDC anode failed to operate after 12 h at a current density of 0.2 A cm−2 at 650 °C when operated on dry methane, whereas a similar cell with a Ni0.9Fe0.1−GDC anode showed no performance degradation during continuous operation over a period of 50 h. The authors also proposed that methane was oxidized more completely on a NiFe−GDC anode as compared to Ni−GDC. However, some inconsistent results about the iron-alloying effect have also been reported. For example, Lu et al. reported that the anodic overpotential increased after the introduction of 5 wt % Fe into a nickel anode,233 whereas Ishihara et al. gave the opposite conclusion.234 Zhu et al. have also studied the coke formation behavior of Ni−Fe alloy in Ni− Fe−ZrO2 catalysts for methane-fueled SOFCs and found that

Figure 16. Reaction scheme for the direct oxidation of CH4 over the Pd−Ni composite anode. Reprinted with permission from ref 237. Copyright 2006 The Electrochemical Society. R


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conversion was independent of the anodic potential because this reaction was not an electrochemical reaction. Under OCV conditions, the surface of the Pd−Ni catalyst was rapidly covered with a small amount of carbon. During the fuel cell reaction, H2 produced by the CH4 cracking was electrochemically oxidized to H2O, which could react with the deposited carbon to produce CO and H2. The produced CO was electrochemically oxidized to CO2 at lower cell voltages. Thus, the Pd−Ni catalyst could maintain a clean surface under the fuel cell conditions. Therefore, the continuous operation of the SOFC on dry methane fuel is possible with this alloyed anode. Through the use of a combination of DFT calculations and several experimental studies, the coke formation rate during hydrocarbon reforming reactions on a Ni catalyst was lowered significantly by the formation of Sn/Ni surface alloys.240−242 The Dacapo pseudopotentials plane wave code was employed for all of the calculations. The GGA-PW91 functional was employed for self-consistent spin-polarized electronic structure calculations.241 Kan et al. have demonstrated a Sn-doped Ni− YSZ anode used for SOFCs that operate on methane.243 The fuel cell with the Sn-doped Ni−YSZ anode delivered a power output comparable to that of the fuel cell with a Ni−YSZ anode when operated on either hydrogen or methane fuel. The authors found that the rapid coke formation rates on the Ni− YSZ anode led to the destruction of the fuel cell, whereas significantly less carbon was deposited on the Sn-doped Ni− YSZ anode, which resulted in improved stability of the fuel cell. A small amount of Sn added to the anode resulted in the best cell performance, whereas higher levels of Sn resulted in diminished performance due to reduced numbers of catalytically active sites. The authors also found that the properties of the deposited carbon on the anodes differed: graphitic carbon was deposited on the Ni−YSZ, whereas amorphous carbon was deposited on the Sn-doped Ni−YSZ. This effect is beneficial because amorphous carbon is typically easier to eliminate than graphite. The authors believed that the deposited carbon on the Sn-doped Ni−YSZ surface could be removed during operation under a certain polarization current density, and the long-term stability of the fuel cell was further improved and stable operation was achieved for 300 h without any degradation in performance. The Sn-doped Ni−YSZ anode also showed good performance with respect to being regenerated (i.e., carbon removal) with an oxidant. Furthermore, as shown in Figure 17, the quantity and distribution of Sn on the anode surface remained almost unchanged after the fuel cell was operated. The authors found that Sn was distributed uniformly on the Ni surface and that the change in Sn content was negligible after 300 h operation. The previous results confirm that Sn on the Ni surface was stable despite numerous hours of operation. Nikolla et al. reported the addition of Sn to Ni-based anodes to form a Sn/Ni-alloy anode electrocatalyst, which exhibited improved resistance to carbon-induced deactivation during the direct utilization of methane and iso-octane in SOFCs as compared to the performance of a conventional monometallic Ni electrocatalyst.244 They observed that the electrochemical activity and electronic conductivity of Ni were not suppressed significantly by the introduction of Sn. They considered that the improved coking resistance of Sn/Ni as compared to that of Ni was due to the formation of Sn/Ni surface alloys that, unlike monometallic Ni, preferentially oxidized carbon atoms and fragments. Recently, Singh and Hill studied the influence of Sn added to Ni−YSZ anodes on the carbon deposition behavior, electro-

Figure 17. EDX images of Sn-doped Ni−YSZ fuel cell (1Anode/1FL) after the cell operated for 300 h: bulk anode showing (a) Ni, (b) Sn and the catalyst layer showing (c) Ni, (d) Sn distribution. Reprinted with permission from ref 243. Copyright 2010 Elsevier.

chemical performance, and operational stability of anodes operated on CH4.245 The electrochemical performance under H2 and CH4 decreased when the Sn content was increased from 1% to 5%. The authors found that Sn was segregated into the Ni particle surface and that Sn then occupied the electrochemically active sites; thus, the polarization resistance was increased. However, the operational stability of the Ni−YSZ anodes under CH4 was not significantly improved by the addition of Sn, which could be attributed to the different preparation methods of the Sn/Ni alloy and/or the different cell configurations compared to those used in the previous study.243,244 The presence of Sn did not have a positive effect on the coke formation on the Ni−YSZ anodes. The greater amount of carbon accumulated on the SnNi−YSZ anodes as compared to that accumulated on the Ni−YSZ anodes, however, indicated that the presence of Sn impeded the carbon removal reactions. Performances of fuel cells with selected Nibased anodes with surface decoration or alloying of the nickel with other metals are listed in Table 1. As shown in Table 1, besides the anode composition, the fuel cell performance and operational stability are also strongly dependent on the fuel composition, operating temperature, cell configurations, and the materials of the electrolyte and the cathode. On the whole, the addition of steam or other oxidant is helpful to improve the operational stability. In summary, Cu is widely used to form the alloy with Ni to suppress the coke formation, and considerable cell performances are obtained. However, the technological barrier of NiCu alloy-based SOFC is its multiple impregnation-calcination routes and the selection of electrolyte-supported cell configuration due to the low melting points of copper oxides. The special anode fabrication processes appear tedious and costly. Moreover, the fuel cell performance of Cu-based cermets is relative poor because Cu is inert to the electrochemical oxidation reaction of the fuels. Furthermore, Cu−Ni alloy is not stable in the hydrocarbon atmospheres. To reduce carbon deposition to acceptable levels, high ratios of Cu to Ni are still required, and as a result Cu−Ni alloying anodes have only a little better thermal stability than those that with pure Cu-cermet anodes. Fe is superior to Cu in alloying with Ni for the anodes in SOFCs due to its relative high catalytic activity S


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Table 1. Performance of Fuel Cell with Selected Ni-Based Anode with Surface Decoration or Alloying of the Nickel with Other Metals fuel cell performance anode 4 wt % Au−Ni/GDC 5 mol % Cu−Ni/SDC Cu−Ni/YSZ Cu−Ni/YSZ 3 wt % Ru−Ni/GDC 2 wt % Ni−Ni/ScSZ Cu−Ni(4:1)−CeO2−YSZ NiCu−GDC Ni0.95Cu0.05−SDC Ni0.9Fe0.1/GDC Ni8−Fe2−LSGMC Pd−Ni(LSCr+SDC)/ SDC Sn−Ni/YSZ Sn−N/YSZ 1 wt % Sn−Ni/YSZ

electrolyte/cathode

temp (°C)

PPD (mW cm−2)

fuel

test period (h)

ref

YSZ/LSM−YSZ/ La0.8Sr0.2MnO3(LSM) SDC/SSC−SDC YSZ/LSM−YSZ YSZ/LSM−YSZ GDC/Sm0.5Sr0.5CoO3(SSC) ScSZ/La0.8Sr0.2Sc0.1Mn0.9O3−δ (LSSM) YSZ/LSM−YSZ GDC−LSCF GDC/SSC−SDC GDC/LSCF LSGMC/SSC−LSGMC LSGM/La0.8Sr0.2CoO3(LSC)

850

CH4:H2O (2.07:1)

0.410

200

204

600 700 800 600 750 800 800 600 650 800 900

dry CH4 3% H2O−CH4 dry CH4 dry CH4 CH4:O2 (1.3:1) dry CH4 dry CH4 BPG dry CH4 DME dry CH4

0.317 0.24 0.049 0.75 0.079 0.33 0.32 0.35 0.34 0.90 0.42

12 200 25 0.33

207 208 210 213 223 226 227 229 231 232 237

YSZ/LSM/YSZ YSZ/LSM−YSZ

650 740

YSZ/LSM−YSZ

800

3% H2O−CH4 5.3% iso-octane+7.8% air+balanced N2 3% H2O−CH4

and melting points of Fe oxides. However, there are some discrepancies about the effectiveness of Fe to improve the coke resistance, which could be attributed to the different preparation methods. Thus, it is critical to optimize the preparation methods for Ni−Fe alloy anodes. For the Snmodified Ni−YSZ anodes, there are also some discrepancies concerning the improvement of Sn addition in heightening the cell performance and the operational stability. 3.4.4. Anode Catalyst Layer. The deposition of a catalyst layer with high activity for the reforming/partial oxidation of hydrocarbons and good resistance toward coke formation on conventional Ni-based cermet anodes was also proposed to increase the operational stability and performance of hydrocarbon-fueled SOFCs.246−251 Zhan and Barnett first demonstrated that Ru−CeO2 could be used as a material for an anode catalyst layer.246 They applied a thin-layer of Ru−CeO2 catalyst to the outer surface of a Ni− YSZ anode. This layer was used to catalyze the reforming of hydrocarbons into syngas, which limited the direct exposure of the Ni-based anodes to the hydrocarbons. The lower sensitivity of Ru−CeO2 to coking as compared to that of Ni allowed the cells with this design to operate in an internal reforming mode using a wider range of fuels. It was also reported that the catalyst layer could serve as a diffusion barrier layer to suppress coke formation by reducing the methane concentration and increasing the concentration of steam and CO2 within the anode under current density conditions.192,252,253 For instance, Zhu et al. have extended and applied physical models to predict performance in relatively large tubular cells with barrier layers. The results show that barrier layers can be designed to develop SOFC systems that are capable of operating coke-free on hydrocarbons, with all reforming accomplished internally.253 Schematics of SOFCs with an anode catalyst layer that operate on CH4−O2 gas mixtures are shown in Figure 18. CH4 and O2 are first converted into CO and H2 over the catalyst layer, and then CO and H2 diffuse to the anode, where they are electrochemically oxidized to produce H2O, CO2, and electricity. Because CO and H2 have higher electrochemical activities than methane, improved cell performance is expected

500 1300 12 50

0.41 ∼0.40

120 ∼15

243 244

0.08

25

245

Figure 18. Schematic diagram of a SOFC with a catalyst layer during operation on CH4−O2 gas mixtures.

through the introduction of the anode catalyst layer. In an SOFC system, the anode catalyst layer may also perform as a diffusion barrier to reduce the respective diffusion rates of reactants and products into and out of the anode layer. These reduced diffusion rates increased the local concentration of electrochemically produced steam or CO2 within the porous anode structure and produced higher local O/C ratios that enhanced the selectivity of the internal reforming reactions toward H2 and CO and suppressed coke formation. However, the adoption of an anode barrier has also been found to decrease the power output due to increased gas-transport losses.254 Concentration polarization was observed in the SOFCs with a Ru−CeO2 catalyst layer, which was prepared via the physical mixing of RuO2 and CeO2 with low surface areas.255 T


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Figure 19. Representative SEM images (a,b) and TEM images (c,d) of the flowerlike mesoporous CeO2 microspheres; the inset in (a) is a fractured CeO2 microsphere. Reprinted with permission from ref 247. Copyright 2006 Elsevier.

Figure 20. Schematic diagram of possible formation processes of mesoporous SDC powder using in situ created NiO as template. Reprinted with permission from ref 256. Copyright 2007 Elsevier.

on iso-octane/CO2/H2O/air as fuel, whereas the performance of the fuel cell with a flowerlike mesoporous CeO2−Ru microsphere catalyst layer was obviously improved, with a PPD of 0.654 W cm−2 under the same operational conditions, due to its open 3D porous structure and large pore volume, as shown in Figure 19. The results indicated that the flowerlike CeO2− Ru microsphere catalyst layer was more beneficial for the mass transport processes. Wang et al. have also used a mesoporous SDC-supported Ru catalyst layer in methane-fueled SOFCs, which they prepared using GNP with NiO created in situ as a template, as shown in Figure 20.256 Significant differences in the crystallite sizes, the Brunauer−Emmett−Teller (BET) specific surface areas, and the pore volumes of SDC were observed between samples prepared with and without the presence of nickel nitrate in the precursor. The increase in the NiO-to-SDC weight ratio caused the decrease in the SDC crystallite size, which suggests that nickel oxide actually acted as an inhibitor for the crystallite size

The cell performance was shown to be further enhanced through improvements in the microstructure of the catalyst layer. Sun et al. used a flowerlike mesoporous CeO2microsphere-supported Ru catalyst as the catalyst layer for IT-SOFCs.247 As shown in Figure 19a and c, most of the CeO2 particles are monodisperse spherical particles. The average diameter of the microspheres is 1−3 μm. A fractured CeO2 microsphere clearly demonstrated a hollow microstructure. According to high-magnification SEM images, such as that shown in Figure 19b, the microsphere was composed of numerous nanosheets with an average thickness of approximately 20−30 nm, and the petals formed a flowerlike texture. These nanosheets interweave together to form an open porous structure. Figure 19d shows a TEM image of an individual microsphere, which displayed an obvious contrast between the dark edge and the light center. The fuel cells with a conventional Ru-loaded porous CeO2 catalyst layer yielded a PPD of approximately 0.39 W cm−2 at 600 °C when operated U


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based anodes. These results suggest that strategies should be developed to increase the coking resistance of the catalyst layer to ensure long-term operating stability of the fuel cells. Wang et al. used Li2O, La2O3, and CaO promoters to modify Ni/Al2O3 catalysts to improve its resistance to coking during methane conversion reactions.259 Among the various catalysts, LiLaNi/Al2O3 exhibited the best catalytic activity and stability. The amount of carbon deposited over the LiLaNi/Al2O3 catalyst was only 18.5% of that deposited over Ni/Al2O3 under the same conditions. Considerable power output comparable to that achieved with hydrogen fuel was obtained on methane-containing fuels with the fuel cells that contained a LiLaNi/Al2O3 catalyst layer. Wang et al. also studied the effect of lanthanide promoters on the catalytic activity and carbon deposition behavior of the Ni/Al2O3 catalyst for methane conversion reactions.167 The addition of the lanthanide promoter decreased the amount of carbon deposited on Ni/ Al2O3, whereas the Gd2O3-promoted Ni/Al2O3 catalyst exhibited the best coking resistance because of its high surface basicity. High power outputs comparable to those achieved with hydrogen fuel were obtained when the fuel cells with this selected catalyst layer were operated on methane-reforming gases at 650−850 °C. All of the catalysts investigated thus far have exhibited poor electrical conductivity, which has caused serious current collecting problems in practical applications. Copper exhibits one of the metals with highest electrical conductivities but also exhibits poor activity toward the methane cracking reaction.48,226 Wang et al. added different amounts of copper (50, 60, and 80 wt %) to the previously discussed LiLaNi/Al2O3 catalyst to improve its apparent electronic conductivity.260 The surface conductivity of LiLaNi/Al2O3−Cu (50:50) was comparable to that of conventional Ni−YSZ anodes, and the addition of 50 wt % copper did not impair the good coking resistance of the LiLaNi/Al2O3 catalyst. Considerable power output was obtained for a fuel cell with the LiLaNi/Al2O3−Cu (50:50) catalyst layer that was operated on methane-reforming gases. In addition to the Ru-based and Ni-based catalysts, some other catalysts were also investigated as the materials for the anode catalyst layer. Jin et al. have developed a Cu1.3Mn1.7O4 spinel-type oxide as an anode catalyst layer for direct-methane SOFCs.261 No carbon was detected in the fuel cell after it was operated on methane fuel. For the fuel cell with the Cu1.3Mn1.7 O4 catalyst layer, a reasonable stability was maintained for 60 h, whereas the fuel cell without a catalyst layer was deactivated quickly and failed after being operated for only 16 h. The authors believed that the highly stable MnO and its strong interaction with Cu particles led to the excellent durability of the catalyst. Recently, Suzuki et al. reported the use of pure ceria as the catalyst layer of low-temperature (LT) SOFCs operated on methane.262 At 554 °C, the performance of the fuel cell without the catalyst layer, when operated on methane−steam fuel, was comparable to that of the fuel cell operated on hydrogen; the good cell performance was attributed to the relatively high catalytic activity of the Ni− GDC anode at higher operating temperatures. However, the power output of the cell without the catalyst layer abruptly decreased if the operating temperature was decreased to 449 °C because of a lack of catalytic activity of the Ni−GDC anode, whereas the cell with a catalyst layer showed a PPD greater than 0.1 W cm−2 even at 449 °C. This PPD was more than 10 times greater than that of the cell without the catalyst layer, which

growth of the SDC powders. A SDC powder with a high surface area 77 m2 g−1 and a high pore volume 0.23 cm3 g−1 were obtained at a NiO-to-SDC weight ratio of 9; this surface area and pore volume are significantly greater than those of the corresponding powder prepared without NiO templating. The cells with and without the catalyst layer showed comparable performances when operated on hydrogen at intermediate temperatures, which suggests that the catalyst layer did not negatively affect the fuel-cell performance with hydrogen fuel. However, the PPD improved from 369 to 462 mW cm−2 at 650 °C when the fuel cell was operated on 3 vol % H2O-humidified methane fuel after the anode was modified with a Ru−SDC catalyst layer. Furthermore, concentration polarization was also reduced as compared to the results reported in the literature.251 No observable agglomerated Ru grain was observed by SEM, which suggests that Ru was homogeneously distributed over the pore walls of the mesoporous SDC. One shortcoming of Ru−CeO2 is its poor thermomechanical compatibility with Ni-based cermet anodes. Delamination of the Ru−CeO2 catalyst layer from the anode surface was observed after several thermal or redox cycles.142 In addition, the high Ru loading also makes it less economically attractive. Later studies consequently focused on the reduction of the Ru content and/or the replacement of Ru with some other inexpensive metal(s) with high catalytic activity for the conversion of methane into syngas. Wang et al. have studied the effect of the Ru content on the catalytic activities of Ru− Al2O3 catalysts for partial oxidation, steam, and CO2 reforming of methane under SOFC conditions.257 They found that 1 wt % Ru−Al2O3 exhibited the worst catalytic performance among all of the catalysts, whereas the other three catalysts exhibited high and comparable catalytic activities. To address the issue of economy, they then selected 3 wt % Ru−Al2O3 as the anode catalyst layer, and the related fuel cell delivered high power output when operated on methane-reforming gases (gas mixtures of CH4−O2 or CH4−H2O or CH4−CO2). They further demonstrated that this Ru−Al2O3 catalyst layer exhibited good thermo-mechanical compatibility with Ni− YSZ anode after repeated thermal and redox cycling tests. GNP was found to be a suitable method for the production of fairly fine, homogeneous, and complex compositional Nibased catalyst powders. In addition, GNP has numerous other advantages, including a relatively low cost, a high energy efficiency, a fast heating rate, a short reaction time, and high compositional homogeneity.133,258 Wang et al. have prepared a Ni/Al2O3 catalyst by GNP for use as an anode catalyst layer in SOFCs operating on methane.142 They found that the Ni/ Al2O3 catalyst exhibited high activity comparable to that of Ru− CeO2 and significantly greater than that of the Ni−ScSZ cermet anode for the partial oxidation, steam, and CO2 reforming of methane in the temperature range of 600−850 °C. They demonstrated that Ni/Al2O3 was superior to Ru−CeO2 in applications where repeated thermal cycling and redox cycling was required. The PPDs for the fuel cells increased significantly when operated on both methane-reforming gases and pure methane after the Ni/Al2O3 catalyst layer was added. The Ni/ Al2O3 catalyst layer also functioned as a gas diffusion barrier, which effectively increased the O/C ratio within the anode layer under a certain current density and suppressed coke formation to maintain cell integrity; consequently, the cell’s operating stability was effectively improved.249 However, O2− TPO analysis demonstrated that Ni/Al2O3 did not show intrinsically improved coking resistance as compared to NiV


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Table 2. Performance of the Fuel Cell with Selected Anode Catalyst Layer fuel cell performance catalyst layer/anode

electrolyte/cathode

temp (°C)

fuel

PPD (W cm−2)

test period (h)

ref

Ru−CeO2/Ni−YSZ Ru−CeO2/Ni−SDC Ru−SDC/Ni−SDC 3 wt % Ru−Al2O3/Ni−YSZ Cu−CeO2/Ni−YSZ Ni−Ce0.8Zr0.2O2/Ni−YSZ Cu1.3Mn1.7O4−SDC/Ni−SDC 7 wt % Ni−Al2O3/Ni−ScSZ LiLaNi−Al2O3/Ni−ScSZ LiLaNi−Al2O3+Cu/Ni−YSZ CeO2/Ni−GDC

YSZ/LSM−YSZ GDC/SSC SDC/BSCF YSZ/LSM−YSZ ScSZ/(Pr0.7Ca0.3)0.9MnO3 (PCM) YSZ/SDC/BSCF−SDC SDC/LSCF ScSZ/LSM ScSZ/LSM YSZ/LSM−YSZ GDC/LSCF−GDC

770 600 650 850 800 700 650 850 850 850 449

5% C8H18+9% air +86% CO2 C8H18/air/H2O/CO2 3%H2O−CH4 CH4:CO2 (2:1) C2H5OH−H2O C2H5OH−H2O dry CH4 CH4−O2 (4:1) CH4:H2O (2:1) CH4:O2 (4:1) CH4:H2O (1:4)

0.6 0.654 0.462 0.929 0.566 0.536 0.304 0.456 0.532 1.08 ∼0.1

50

246 247 256 257 248 249 250 252 259 260 262

80 80

∼7

the solid-state reactions between NiO and various oxides during the cell fabrication processes. The performance of the cells with Ni−Sm2O3 and Ni−Eu2O3 anodes was found to be limited by concentration polarization. He et al. have also investigated Ni− LnOx (Ln = Dy, Ho, Er, Yb, and Tb) cermets as the anodes of IT-SOFCs with ceria-based electrolytes.264 These electrodes exhibited similar electrochemical activities that were comparable to that of the common Ni−SDC cermets. The anode performance was found to strongly depend on the cermet composition and porosity. On the basis of the results of hydrogen-temperature-programmed reduction (H2-TPR) studies and electrochemical impedance spectroscopy (EIS) analyses under different hydrogen partial pressures, the catalytic activity of the Ni−LnOx cermets might originate from the hydrogen adsorption capability of the LnOx surface, which promotes the hydrogen spillover process and consequently enhances the electrochemical oxidation of the fuel. These results suggest that Ni−LnOx is an attractive alternative anode for SOFCs. He et al. have further investigated the performance of Ni− Dy2O3 cermet anodes in detail. The Dy2O3 exhibited high hydrogen adsorption capability, which might enhance the hydrogen spillover process and thus promote the anodic reaction.265 The catalytic behavior of Dy2O3, given its ability to promote the anodic reaction through hydrogen adsorption, is concluded to be comparable to that of SDC. However, further investigations, such as DFT calculations, are still needed to confirm the enhanced hydrogen adsorption on the Ni surface via spillover to the oxide support. Ni−Sm2O3 cermets have shown high performance as the anodes for SOFCs with doped ceria electrolytes; however, fuel cells composed of a Ni−Sm2O3 anode and a stabilized zirconia electrolyte exhibited relatively low performance when humidified hydrogen was used as the fuel.266 Significant solid-state reactions between Sm2O3 and Zr0.89Sc0.1Ce0.01O2−δ (ScCSZ) were found to occur during the cell fabrication process. To avoid such reactions, Ni−SDC, as a thin interlayer, was deposited between the Ni−Sm2O3 anode and the ScCSZ electrolyte. The cell performance was thus improved significantly from 202 to 410 mW cm−2 at 700 °C, which suggests that Ni−Sm2O3 cermets have electrochemical activities comparable to those of anodes for SOFCs with stabilized zirconia electrolytes. Notably, the Ni−LnOx anodes may also be suitable as anodes for direct-methane SOFCs because rare earth oxides exhibit high activity toward the steam reforming of methane and good coking resistance due to the high activity for carbon gasification and high basicity. Zhou et al. have reported a doubly doped CeO2-based anode with Y and Yb for direct-methane SOFCs.267 At 750 °C, a fuel

proved that the catalyst layer enhanced the catalytic activity toward the steam reforming of methane. This study showed promising performance of the CeO2 catalyst layer for a fuel cell operated on methane−steam fuel. However, the mechanism by which the pure ceria catalyst layer dramatically improved the cell performance is still not well understood because ceria should not exhibit such high catalytic activity toward the steam reforming of methane at a temperature as low as 449 °C. The difference between CeO2 and Ni−GDC anodes with respect to catalytic activity for the steam reforming of methane at lower temperatures thus requires further investigation. Performances of the fuel cells with selected anode catalyst layer are listed in Table 2. At higher operating temperatures, considerable power outputs are obtained operating on hydrocarbon fuels with the help of anode catalyst layer. The study about long-term stability of the fuel cells with anode catalyst layer is still lacking up to now. The progress in the improvement of performance with anode catalyst layer has been made, but there is still scope for further enhancement. For instance, Ru−CeO2 catalysts are expensive. In addition, Ru−CeO2 displays poor thermo-mechanical compatibility with nickel-cermet anodes during thermal and redox cycling, making it less attractive for the practical applications. Ni-based catalysts with low price are considered to be highly active and are good coking-resistant materials for anode catalyst layer. The optimization of Ni content and preparation methods, the modification of surface acidity, and the improvement of conductivity are very important for the large-scale applications. However, the stability of the fuel cell with Ni-based catalyst layer when using hydrocarbon fuels is still not ideal, and it demonstrated that the fuel cell performance is insufficient at reduced temperatures. The Cubased and the pure oxide catalysts could not be the ideal materials for anode catalyst layer due to the relatively low catalytic activity for hydrocarbon conversion, and proper modification is still required. 3.4.5. Tailoring the Ceramic Phase. Considerable research has also been conducted on the development of novel anodes through alteration of the ceramic phases in the anodes. He et al. have developed novel Ni−LnOx cermets (Ln = La, Ce, Pr, Nd, Sm, Eu, Gd) as anodes for IT-SOFCs based on GDC electrolytes, where LnOx are not O2− conductors.263 When 3 vol % H2O humidified H2 was used as the fuel and ambient air as the oxidant, Ni−CeO2 and Ni−Gd2O3 anodes exhibited good performances comparable to that of a fuel cell with a Ni−GDC anode, whereas Ni−La2O3, Ni−Pr6O11, and Ni−Nd6O11 showed poor performance, which was attributed to W


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hydrogen fuels, whereas the addition of excessive CeO2 led to a rapid decrease in power output, especially with methane fuel.270,271 Higher calcination temperatures were found to enhance the operating stability, and the anode calcined at 800 °C exhibited the best performance.271 The PPDs of a fuel cell with an anode that contained 10 wt % CeO2 and 25 wt % Ni, after being calcined at 800 °C, reached 480 mW cm−2 when operated on methane fuel at 800 °C.271 Some oxygen-ionconducting ceria materials, such as doped ceria, have been impregnated onto Ni−YSZ cermet anodes to improve their coking resistance and/or operational stability under hydrocarbon atmospheres.274−277 For example, Zhang et al. have developed a SDC-modified Ni−YSZ anode using an impregnation technique.275 They found that the performance of the fuel cell, when operated on hydrogen, was significantly improved after nanostructured SDC particles were incorporated into the anode. As compared to the unmodified cells, the PPD of the modified cells increased by 60% and the areaspecific resistance decreased by 47% at 700 °C. The authors attributed such improvement to the extended TPB in the anode and to the excellent electrocatalytic properties of nanosized SDC. Chen et al. have also reported the use of a similar Sm0.2(Ce1−xTix)0.8O1.9 (SCTx)-modified Ni−YSZ anode in SOFCs for the direct utilization of methane.277 They found that both the amount of Ti-doping and the SCTx loading in the anode substantially affected the electrochemical activity with respect to methane oxidation. Optimal anode performance was obtained for a SCT0.17-modified Ni−YSZ anode with a SCT0.17 loading of approximately 241 mg cm−2, which was accomplished through four repeated impregnation cycles. The fuel cell showed a cell polarization resistance of only 0.63 Ω cm−2 under OCV conditions and a PPD of 383 mW cm−2 at 700 °C. The anode modified with SCT0.17 exhibited better cell performance than that modified by SDC when operated on methane fuel, which indicates that the Ni−YSZ anode modified with SCT0.17 exhibited greater electrochemical activity toward methane oxidation; the authors attributed this improved electrochemical activity to the improved OSC and/or to the rapid redox equilibrium of SCT0.17. To improve the coking resistance and stability of SOFCs that operate on methane, some researchers have also attempted to modify the conventional nickel cermet anode with protonconducting oxides. Jin et al. have reported the fabrication of proton-conductor SrZr0.95Y0.05O3‑δ (SZY)-promoted Ni−YSZ anodes for the direct utilization of methane in SOFCs.278−280 They found that the performance of a cell with the SZYmodified Ni−YSZ anode was superior to that of a cell with the conventional Ni−YSZ anode and that the PPDs of the fuel cells increased as the amount of SZY phase in the anodes was increased when the cells were operated on both 1 vol % H2Ohumidified hydrogen and dry methane fuels; the optimal amount of infiltrated SZY was approximately 1/14 (SZY/NiO by weight). The degradation ratio of the cell with Ni−YSZ anode operated on methane fuel decreased from approximately 64% to 22% after the anode was infiltrated with SZY at a ratio of approximately 1/14 (SZY/NiO by weight).279 The authors further demonstrated that the addition of SZY dramatically reduced the amount of carbon deposited under OCV conditions, which they attributed to the increased amount of adsorbed oxygen near the TPB. Such increased adsorbed oxygen, which was derived from the interaction among Ni as an electronic conductor, YSZ as an O2− conductor, and SZY as a proton conductor near the TPB, consumed the deposited

cell with the Ni−Ce0.8Y0.1Yb0.1O1.9 anode displayed stable power output for 120 h at 200 mA cm−2 when operated on dry CH4; however, the fuel cell with Ni−YSZ anode dropped rapidly for only 1.25 h. This result suggested that carbon deposition was largely absent on this anode, which was confirmed by SEM observations and energy-dispersive X-ray spectroscopy (EDS) results. The remarkable performances suggested that co-doped CeO2 may be an attractive electrode component for direct-methane SOFCs. 3.4.6. Modification of Electrode Surface with Other Oxide(s). The conventional Ni cermet anodes of SOFCs are still the preferred choice for practical applications because of their low cost, high electronic and reasonable ionic conductivities, high stability under reducing atmospheres at high temperatures, and moderate thermal expansion coefficient (TEC) matching with the YSZ electrolyte. Various strategies have been applied to prevent the surface of Ni particles in Nibased anodes from being directly exposed to methane fuel, to prevent the sintering of Ni, and to reduce carbon deposition through surface modifications. In addition to metals, oxides have also been extensively used to modify the properties and performance of nickel cermet anodes, particularly with respect to coking resistance. Many research groups have devoted themselves to the development of anodes with high electrocatalytic activity and superior coking resistance toward methane fuel through the modification.268−271 The modification may improve the anode kinetics of the fuel oxidation through the addition of a catalytically active component, such as nanosized ceria-based materials, which exhibit superior ability to promote the WGS reaction, thermal stability, and OSC.268−271 For example, Takeguchi et al. have studied the effect of the addition of CeO2 to Ni−YSZ cermets with respect to their catalytic activity and tendency to form carbon deposits during the steam reforming of methane under SOFC operating conditions.265 The carbon deposition rate and catalytic activity were found to be intimately related to the concentration of CeO2, in good agreement with the results reported by Qiao et al.271 A small amount of CeO2 addition (1 wt %) suppressed the deposition of carbon; however, beyond a certain concentration limit, the amount of deposited carbon increased as the concentration of CeO2 was increased.269 The authors attributed the different catalytic activity and carbon deposition behavior of CeO2modified Ni−YSZ to the formation of CeO2−ZrO2 solid solutions. Because Ni particles on CeO2−ZrO2 solid solutions were stable,272,273 the CeO2-modified Ni−YSZ cermets with low CeO2 concentrations exhibited high performance for the steam reforming of methane. However, CeO2 concentrations in excess of the limit for the formation of CeO2−ZrO2 solid solutions deactivated the Ni catalysts;272 thus, the CeO2modified Ni−YSZ with high CeO2 concentrations exhibited diminished catalytic performance and also produced a large amount of carbon. In the same manner, Qiao et al. also prepared CeO2-promoted Ni−YSZ anodes for the direct oxidation of methane using a vacuum mix-impregnation method.270,271 With this method, NiO and CeO2 were obtained from nitrates, which eliminated the need for high-temperature sintering. The authors found that at least 25 wt % Ni was needed to ensure low ohmic resistance and that the impregnation of CeO2 into the anode could improve the cell performance, especially with methane fuel. They also found that the addition of a small amount of CeO2 improved the cell performance with respect to its operation on both methane and X


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Figure 21. Schematic of the effects of the SZY distribution on carbon deposition in the cross section of a fractured anode after an accelerated degradation test. Images of location of (a) nonactive area on a Ni−YSZ anode, (b) carbon deposition area on a Ni−YSZ anode, (c) protonconductor SZY in a Ni−YSZ−SZY anode, and (d) carbon deposition and SZY in a Ni−YSZ−SZY anode after an accelerated degradation test. Reprinted with permission from ref 279. Copyright 2010 The Electrochemical Society.

Figure 22. Microstructures of the cermet anodes A (a), AMg (b), and AMgAl (c) and Ni particles observed in A (d) and AMgAl (e). A, AMg, and AMgAl screen-printed on the ScSZ plate were sintered at 1250 °C for 3 h, then reduced in the hydrogen flow at 900 °C for 2 h. Reprinted with permission from ref 284. Copyright 2006 Elsevier.

carbon deposition area is localized near the anode’s surface side (Figure 21a). This result is based on the carbon distribution in the cross section of the Ni/YSZ anode after an accelerated degradation test (Figure 21b). After the nonactive area of the Ni−YSZ anode was selectively modified with SZY (Figure 21c), the performance and stability of the anode against carbon deposition significantly increased in dry methane fuel, although the concentration polarization resistance also increased due to the addition of SZY (Figure 21d). The authors considered that carbon deposition was suppressed under OCV conditions and that the carbon removal reaction was enhanced due to the increased amount of H2O electrochemically produced by SZY.

carbon. The amount of carbon deposited near the anode’s top surface was found to be approximately 2 times greater than that deposited near the electrolyte side, and SZY was distributed primarily near the anode’s top surface (i.e., the nonactive area). These results suggested that SZY in the nonactive area electrochemically promoted the production of water around the TPB and that SZY localized near the nonactive area significantly contributed to the generation of H2O and the adsorption of oxygen for the removal of deposited carbon. Figure 21 schematically shows the effect of SZY on the carbon deposition behavior of the Ni−YSZ anode.279 They considered that the nonactive area of the Ni/YSZ in terms of the favored Y


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A 3D homogeneous distribution of SZY in the Ni−YSZ cermet anode might yield the improved operational stability against carbon deposition. Recently, Shimada et al. have studied the dependence of the electrode performance on the distribution of proton-conducting yttrium-doped barium zirconate (BZY) in Ni−YSZ anodes to clarify the role of BZY in improving the anode performance with dry methane fuel;281 the anode microstructure and the BZY distribution in the anodes could be well controlled through the use of a spray pyrolysis technique for the electrode fabrication. The effective thickness of Ni−YSZ anode was found to be approximately 25 and 38 μm in 3 vol % humidified H2 and in dry CH4, respectively, at 900 °C, and BZY was distributed either inside or outside the electrochemically active zone. Consequently, the addition of BZY to the inside was effective, whereas its addition to the outside was not effective in improving the anode performance; these results indicate that BZY primarily promoted the electrochemical reactions. In addition to oxygen-ion-conducting and proton-conducting oxides, some nonconductive oxides can also improve its catalytic activity and/or coking resistance.282−291 For example, Yan et al. have reported that La2O3 successfully prevented the sintering of nickel in the anodes.282 Shiratori et al. have reported the co-modification of Ni−ScSZ anodes with MgO and Al2O3 effectively improved the degree of dispersion of Ni particles and also reduced the Ni particle size to suppress coke formation.284 Three types of cermets based on NiO−ScSZ (A), Ni0.9Mg0.1O−ScSZ (AMg), and Ni0.9Mg0.095Al0.005O−ScSZ (AMgAl) were used as the anodes. As shown in Figure 22, no microstructural differences were observed between AMg and AMgAl geometrically. Figure 22d and e shows the Ni particles formed in A and AMgAl, respectively. The nickel particles in AMgAl were highly dispersed with an average size of less than 50 nm; however, the particles grew to 1 μm in size in A. These results indicate that the addition of MgO and/or Al2O3 prevented the sintering of nickel. The Ni phases in AMg and AMgAl (dark gray regions in Figure 22b and c) exhibited microstructures analogous to the Ni phases that were prepared without sintering. In contrast, the clusters of small Ni particles with an average size of 30 nm were observed on the powder without sintering, and the clusters of Ni particles were observed as a single large particle with a smooth surface on A, as shown in Figure 22d. The authors considered that the microstructural changes induced by the addition of other oxides were closely related to the change in the reduction behaviors. When the fuel cells were operated on hydrogen, the addition of MgO and Al2O3 to the anodes barely influenced the anodic overvoltage; therefore, the TPB areas of AMg and AMgAl were expected to be similar to that of A under these conditions. The results are schematically summarized in Figure 23. Considerable research efforts have also been expended toward the addition of other basic oxides, such as BaO, SrO, and CaO, to Ni-based anodes to improve their coking resistance.269,289−291 Yang et al. have reported a new anode with nanostructured barium oxide/nickel (BaO/Ni) interfaces for low-cost SOFCs;291 these anodes exhibited high power density and stability in C3H8, CO, and gasified carbon fuels at 750 °C. Nanosized BaO islands grew on the Ni surface, which created numerous nanostructured BaO/Ni interfaces that could adsorb H2O and promote the water-mediated carbon removal reactions. DFT calculations predicted that the dissociated OH from H2O on BaO reacted with C on Ni near the BaO/Ni interface to produce CO and H species, which were then

Figure 23. Schematic diagrams of the cermet anodes: (a) conventional anode (A type) and (b) novel anode (AMg or AMgAl). Reprinted with permission from ref 284. Copyright 2006 Elsevier.

electrochemically oxidized at the TPBs of the anode. The constructed models for DFT calculations used the VASP and GGA with the PW91 exchange-correlation functional. The authors hypothesized that the dissociation of water occurred on the BaO, that the coke formation occurred on the Ni sites of BaO/Ni, and that all subsequent steps occurred at or near the BaO/Ni interfaces; this hypothesis implies that the high performance and coking tolerance of this new anode relied heavily on the direct participation of the BaO/Ni interfaces. They found that the important intermediate in water-mediated carbon removal reactions was hydroxylated BaO. Notably, however, the hydroxylated BaO is not sufficiently stable at high temperatures in the presence of CO2. Instead, barium carbonate may also be an important intermediate in the carbon-removal reaction, at least as important as the hydroxylated BaO. Other oxides have also been used as promoters in Ni-based anodes.179,292 Finnerty et al. demonstrated that the addition of small quantities of molybdenum to Ni−YSZ anodes substantially reduced the level of carbon deposition.179 The addition of as little as 1 wt % MoO3 decreased the amount of carbon deposited by 75%. Such low molybdenum content did not decrease the catalytic activity of the anode toward methane reforming, and cell performance was also only slightly diminished. Recently, however, Tavares et al. gave an opposite conclusion.293 They found that the deposition of MoOx could partially cover the electrocatalytic sites and result in poor anode performance and that the MoOx did not prevent carbon deposition. Additional research is required to clarify such discrepancies. Myung et al. have developed a Li2TiO3-modified Ni−YSZ anode for SOFCs that operate on methane.292 The Li2TiO3 catalyst showed excellent resistance to carbon deposition on Li2TiO3-catalyst-doped Ni−YSZ anodes. The amount of carbon deposited was greatly suppressed by the addition of a small amount of Li2TiO3, especially at a Li2TiO3 doping level of 4 wt %. The PPD of the 1 wt % Li2TiO3-doped Z


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Ni−YSZ anode-supported fuel cells was 0.4 W cm−2 at 800 °C when they were operated on methane fuel, which was slightly lower than that of a conventional fuel cell operated under the same conditions (0.47 W cm−2). In particular, the long-term stability of the fuel cell was dramatically improved with the addition of 1 wt % Li2TiO3 when operated on methane fuel, which was attributed to the high coking resistance of the Li2TiO3 catalyst. However, this new anode was tested only at 800 °C, and the carbon deposition behavior at lower temperatures still needs to be evaluated. In addition to the composition of the material, the morphology of foreign oxide particles also significant influences the suppression of coke formation over nickel cermet anodes. For example, obvious improvement in coking resistance was observed after a SDC thin film was coated over nickel particles.294,295 A related fuel cell with a SDC thin film modified Ni−SDC anode prepared using a dip-coating method was successfully operated on methane−air fuel for 500 h without a significant degradation in the performance.294 A similar cell with porous SDC thin film modified Ni−SDC anodes prepared using a sol−gel coating method also exhibited no significant performance degradation during 180 h of operation on methane fuel, whereas severe degradation of the cell performance was observed for cells with a conventional Ni−SDC or Ni−YSZ anode.295 According to electron-probe micro-analyzer (EPMA) mapping images shown in Figure 24, a

Figure 25. Schematic illustration of an anode microstructure coated with SDC. Reprinted with permission from ref 294. Copyright 2004 Elsevier.

method provided individual ionic pathways (through the continuous SDC layer) and electronic pathways (through the Ni phase); thus, the carbon deposited on the TPB sites was more readily oxidized by O2− over the whole anode surface. Furthermore, the fine SDC particles prevented the surface diffusion of Ni via a pinning effect, as a result, both the catalytic activity and the resistance to coke formation were improved. Fuel cell performances of selected oxide modified anode materials are listed in Table 3. Rare earth doped CeO2 is a good promoter for Ni-based cermet anodes of SOFCs due to its advantageously high electrocatalytic activity and redox properties to suppress the coke formation. Although considerable improvement in cell performance is reported, no excellent operational stability was obtained up to now. The fuel cells with SDC thin-film modification delivered a considerable power output and excellent operational stability toward methane, and the operating mechanisms of these anodes are available, suggesting this approach could be widely used in the research field of oxide-modification in the future. It is well-known that the basic oxides could improve the coking resistance of Nibased catalysts due to the decreased surface acidity. BaO was found to be an effective promoter to improve the operational stability without sacrificing the power output, while the other additives such as SrO, CaO would reduce the power output of the fuel cell with Ni-based anodes. Finally, with further optimization of the composition and microstructure, the performance of these anode materials may be further improved and hopefully replace the traditional Ni−YSZ anodes in the future.

4. SULFUR POISONING 4.1. Types and Critical Amount of Sulfur Contaminants

Sulfur compounds are known to poison the catalytic activity of numerous metals, even at very low concentrations. Most sources of hydrocarbon fuels, including natural gas, oil, and energy materials from biomass, are potential fuels for SOFCs contain sulfur-based impurities. Sulfur is present in several compounds, including carbonyl sulfide (COS), disulfides, mercaptans, and tetrahydrothiophene. These compounds can all be converted into H2S under SOFC operating conditions in a H2-rich atmosphere. Therefore, for simplicity, H2S is often used in experiments as the source of sulfur to study the impact of sulfur on SOFC performance. In addition, most investigations of the sulfur poisoning effect on fuel-cell anodes have been conducted using H2S-containing hydrogen to avoid the potential contribution of coke formation to the degradation of cell performance. Electrochemical studies conducted on Ni−YSZ cermet anodes have shown that, at a given temperature, the extent of degradation increased as the partial pressure of H2S was

Figure 24. Concentrations of carbon deposition (left bottom) from EPMA mapping on the surface of the Ni−YSZ anode (a), the Ni− SDC anode (b), and the SDC-coated Ni−SDC anode (c). Reprinted with permission from ref 295. Copyright 2012 International Association of Hydrogen Energy.

significant amount of bulk carbon and/or carbon particles was detected on the Ni−YSZ and Ni−SDC anode (Figure 24a and b), whereas almost no carbon was detected on the SDC-coated Ni−SDC anode (Figure 24c). The mechanisms for the improved coking resistance and the catalytic performance were proposed by Yoon et al. and are illustrated in Figure 25.290 The continuous SDC coating layer prepared by the dip-coating AA


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Table 3. Fuel Cell Performance of Selected Oxide Modified Anode Materials fuel cell performance anode

electrolyte/cathode

temp (°C)

fuel

PPD (W cm−2)

test period (h)

ref

25 wt % Ni−10 wt % CeO2/YSZ SDC−Ni/YSZ SDC−Ni/SDC SCT−Ni/YSZ SZY−Ni/YSZ Ni0.95La0.05O/SDC (Ni0.75Fe0.25−MgO)/YSZ Ni−CaO/SDC Ni0.53Cu0.47+BaO BaO−Ni/YSZ Li2TiO3−Ni/YSZ SDC−Ni/YSZ SDC−Ni/SDC

YSZ/LSM YSZ/LSM−SDC SDC/SSC YSZ/LSCF−SDC YSZ/LSM SDC/SSC−SDC YSZ/LSM−YSZ ScSZ/Pt GDC/LSCF YSZ/LSCF−SDC YSZ/LSM−YSZ YSZ/LSM YSZ/Pt

800 800 600 700 900 600 800 700 750 750 750 700 850

dry CH4 dry H2 3% H2O−H2 dry H2 dry CH4 3% H2O−CH4 3% H2O−CH4 dry CH4 dry CH4 dry C3H8 dry CH4 25% CH4+75% air dry CH4

0.48 0.525 0.75 0.383 0.28 0.55 0.648 0.03 ∼0.3 ∼0.9 0.4 0.3 0.248

∼5 70

271 275 276 277 279 282 286 289 290 291 292 294 295

80 ∼8 20 24 200 100 50 500 200

found to be caused primarily by the physical adsorption of sulfur that covered and/or blocked the hydrogen reaction sites, and the formation of nickel sulfide could lead to the complete catalytic loss of the Ni catalyst.303 Thus, sulfur adsorption was considered as the major cause for performance degradation at low H2S concentrations in the fuel (i.e., less than 20 ppm). For SOFCs that operate on fuel with higher H2S concentrations, the formation of nickel sulfide on the anode surface is primarily responsible for the deterioration in performance.304 Shown in Figure 27 are the schematics of the high-level summary of section 4.

increased. Furthermore, the critical H2S concentration for the initialization of cell degradation increased with increasing operating temperature, whereas the performance loss increased in the intermediate temperature range for H2S concentrations that ranged from 0.2 to 20 ppm.296,297 The poisoning effects become irreversible at low temperatures due to a more stable adsorption state of sulfur on the Ni surface. The adsorption and poisoning effect of sulfur on Ni was also studied by Bartholomew et al.298,299 Figure 26 shows the influence of

Figure 27. The schematics of the high-level summaries of section 4.

4.2. Possible Poisoning Mechanisms

On the basis of the available knowledge, the possible sulfur poisoning mechanisms have been summarized by Sasaki et al.; the mechanisms are shown in Figure 28.305 For sulfur poisoning at relatively low sulfur concentrations, reversible processes associated with the adsorption/desorption of sulfur are considered to be the predominant mechanism,296,306 as described in Figure 28b. The importance of the dissociative adsorption of sulfur on Ni has been verified both experimentally307 and theoretically.308,309 As schematically shown in Figure 28e and f, the larger voltage drop observed to occur with CO-rich fuels could be partially due to the preferable adsorption of sulfur on Ni in the equilibrium reaction (H2S(g) → Sad + H2(g)) at lower H2 concentrations. At a higher sulfur concentration and/or a lower operating temperature, an irreversible degradation was observed, which has been associated with the oxidation of Ni in cases where the

Figure 26. Equilibrium partial pressure of H2S versus reciprocal temperature (open symbols, θ = 0.5−0.6; closed symbols, θ = 0.8− 0.9) (θ, coverage rate). Reprinted with permission from ref 298. Copyright 1982 Elsevier.

the partial pressure of H2S in a hydrogen atmosphere on the coverage of a nickel surface with H2S (θ).298 The adsorbed sulfur was more stable than bulk sulfide at intermediate temperatures, and 90% surface coverage occurred at 1000 K for a H2S concentration of 10 ppm. The effects of H2S contamination on the Ni−YSZ anodes of SOFCs in H2 and methane fuels have been reported in numerous research papers.296,300−302 The degradation of Ni−YSZ anodes was AB


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Figure 28. Possible sulfur poisoning mechanisms of SOFCs. Reprinted with permission from ref 305. Copyright 2011 Elsevier.

anodic overpotential was very large, as shown in Figure 28g.296 Sulfur poisoning during internal reforming reactions is also a serious issue, as shown in Figure 28d as compared to Figure 28c; significantly larger fuel-cell voltage drops have been observed in cases where CH4-rich fuels were used.306 In addition, according to the C−H−O−Ni−S stability diagram,307 the formation of Ni3S2 (melting point: 787 °C) is possible when hydrogen-poor fuels are used, as shown in Figure 28h. Galea et al. have examined the adsorption and dissociation of consecutive molecules of hydrogen sulfide on a nickel (111) surface to gain insight into the details of poisoning on SOFC anodes by sulfur-based contaminants.310 The preferred adsorption sites, energies, transition states, and kinetic barriers were calculated for the resulting species, *SHx (x = 0−2) and *H. Systematically larger amounts of adsorbed sulfur (0, 25, 50, 75, 100 mol %) were calculated to determine the most energetically favorable coverage of sulfur on the nickel surface. The removal of the existing surface-sulfur atoms was studied to investigate the irreversibility of the H2S adsorption reaction. The authors found that, due to the extent of molecular hydrogen adsorption with different surface sulfur coverages, the presence of even 25 mol % surface sulfur could result in a 50% decrease in the amount of surface molecular hydrogen adsorbed. Their results, which were consistent with the experimental data, demonstrated that the equilibrium coverage of adsorbed sulfur on the surface was 50%. Because of the considerable exothermic nature of the H2S adsorption and dissociation reactions, the reaction exhibits partial irreversibility. This irreversibility poses a challenge for any attempt to remove the surface sulfur and regain the original electrochemical activity of the Ni catalyst.

the fuel before it enters the SOFC’s anode chamber. The two main processes for desulfurization are hydrodesulfurization and sulfur sorption. Hydrodesulfurization includes hydrogenation and hydrogenolysis of the fuels over a catalyst in a H2-rich environment. This approach is useful for fuels with high sulfur content; however, it is difficult to reduce the sulfur concentration to less than 10 ppm, which is required for stable operation of a SOFC with conventional nickel cermet anodes.311,312 Inexpensive sorbents that are capable of reducing sulfur levels to a greater extent than hydrodesulfurization have received increased interest recently. Various materials, including activated carbon313,314 (which may produce toxic COS gas), zeolites315,316 (which are difficult to regenerate), and ceria, have recently received significant interest as potential sulfur sorbents.317,318 4.3.2. Optimization of the Operating Conditions. In addition to the preremoval of sulfur from the fuel gas, which has the disadvantages of adding complexity and cost to the system, the sulfur-poisoning effects can also be mitigated to a certain extent through optimization of the operating conditions. For example, the poisoning effect can be effectively reduced through operation of a SOFC at higher temperatures because the adsorption process is exothermic. High current density was also found to enhance the mitigation of sulfur because of the increased flux of O2−, which could react with sulfur to form SO2. For example, Brightman et al. have studied the effect of current density on the H2S-poisoning of nickel-based anodes in SOFCs, and they observed that the degree of poisoning, as measured by the increase in anode polarization resistance (Rp), was decreased at higher current densities.319 Lakshminarayanan et al. have studied the effect of H2S on the catalytic activity for methane steam reforming and methane oxidation over Ni−YSZ and found that the oxidation sites on the Ni−YSZ catalysts were more resistant to attack, whereas the sites for the steam reforming reaction were significantly affected by sulfur;320 these

4.3. Strategies for Sulfur Mitigation

4.3.1. Desulfurization. The negative effects of sulfur on SOFC anodes could be minimized through desulfurization of AC


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Figure 29. Cell voltage drop with various additives impregnated in porous anode layers at 200 mA cm−2 (800 °C, H2S conc. = 20 ppm, H2/CO = 100:0, electrolyte|SSZ, anode|Ni−YSZ + impregnated additives). Reprinted with permission from ref 296. Copyright 2006 The Electrochemical Society.

stabilized zirconia as the ceramic phase in a composite anode. Hagen et al. have compared the durability of Ni−YSZ, Ni− scandia, and yttria-stabilized zirconia (ScYSZ) anodes in the presence of sulfur-containing fuels.323 In a fuel composed of 13% H2, 29% CH4, 58% H2O, and 2 ppm H2S, the Ni−ScYSZ anode showed excellent stability for 500 h, whereas a Ni−YSZ anode decayed rapidly, mainly due to an irreversible increase in the ohmic resistance. When the H2S concentrations were increased into the range of 10−20 ppm, a degradation of the cell voltage was also observed for Ni−ScYSZ anode. These results suggest that the Ni−ScYSZ anode is still not a practical solution to sulfur poisoning when the fuel gas contains high concentrations of H2S. Some researchers have also devoted themselves to the development of sulfur-tolerant anodes through modification of conventional Ni-based anodes. Cobalt was observed to exhibit better sulfur tolerance than Ni.324 Grgicak et al. compared the behavior of Ni−YSZ anodes modified with cobalt and copper for fuel cells operated in H2, CH4, and H2S/CH4 gas mixtures.325 The incorporation of Cu2+ into the NiO lattice resulted in large metal particle sizes, and a Ni0.92Co0.08O−YSZ anode, which was formed through doping a proper amount of Co2+ into the NiO lattice, exhibited a large BET surface area and active metal surface area, which could increase the hydrogen oxidation rate. The performance of both bimetallics was quickly degraded in dry methane due to the deposition of carbon. However, Ni0.69Co0.31−YSZ showed superior activity in H2S/CH4 gas mixtures and exhibited better cell performance than the same anode operated on H2 fuel. The authors hypothesized that the active anode became a Ni−Co−S-type alloy under the fuel-cell operating conditions and that a synergistic effect was created in the Ni−Co−S anode. However, the durability of this anode in the presence of H2S/CH4 gas mixtures needs to be further studied. Zheng et al. have modified Ni−GDC anodes by impregnating them with Pd nanoparticles to improve their sulfur tolerance when operated on H2−H2S fuels at 800 °C.326 The anodes were tested at a current density of 200 mA cm−2 in H2 and H2S/H2 fuels with a H2S concentration that was gradually increased from 5 to 700 ppm. The degradation ratio in various H2S/H2 fuels, particularly under low H2S concentration conditions, was substantially smaller on Pd-impregnated Ni/ GDC cermet anodes than on pure Ni−GDC anodes. The results showed that Pd impregnation significantly enhanced the

results suggest that the operation of a SOFC on partial oxidation is preferable to steam reforming for fuels that contain sulfur. Wang et al. have studied the surface regeneration process (by O2 and H2O) for a sulfur-poisoned nickel surface using firstprinciples calculations with proper thermodynamic corrections.321 They found that O2 is more effective than H2O in removing the sulfur atoms adsorbed onto the nickel surface; however, the O2 readily reacted with the regenerated Ni surface, which led to an overoxidization of Ni. Thus, H2O appears to be a better choice for the surface-regeneration process. In reality, however, both O2 and H2O may be present in the system under fuel-cell operating conditions. Accordingly, the effects of the partial pressures of O2 [pO2] and H2O [pH2O] as well as that of the pO2/pH2O ratio on the regeneration of a sulfur-covered Ni surfaces without overoxidization at different temperatures are systematically examined to identify the best conditions for the regeneration of Ni-based anodes under practical conditions. The recovery of the performance of a fuel cell’s anode after it has been poisoned with sulfur is also closely related to the operating conditions. Zha et al. have studied the regeneration behavior of Ni-based anodes in SOFCs after they were poisoned with sulfur.322 After H2S was removed from the fuel stream, the anode performance recovered fully or partially, depending on the operating conditions and the time of H2S exposure. The rate of the recovery process increased as the operating temperature and polarization current were increased, which suggests that the current passing through the fuel cell can enhance the electrochemical oxidation of the sulfur adsorbed onto the anode surfaces. 4.3.3. Sulfur-Tolerant Anode Materials. Significant progress has been made in the development of alternative anode materials that are more tolerant to sulfur impurities; advances in this area have been reviewed by Gong et al.303 However, most researchers have focused on the improvement of Ni-based anodes for hydrogen that contains H2S to simplify the system, whereas few detailed investigations of SOFCs that operate on sulfur-containing methane fuel have been reported. Sulfur tolerance can be improved through the use of ScSZ as the electrolyte component in the anode because conductivity is greater than that of YSZ;296 however, ScSZ suffers the shortcoming of being significantly more expensive than other similar alternatives. One solution is to use Sc- and Y-coAD


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sites, respectively. Potential energy profiles for the H2S−CeO2 (111) interactions along the three product channels that produce H2, H2O, and SO2 were constructed using the nudged elastic band (NEB) method. The calculated results showed that H2S weakly bound to CeO2 (111) with a small binding energy, followed by dehydrogenation processes, which formed surface sulfur species with an exothermicity of 29.9 kcal mol−1. Molecular-level calculations demonstrated that the SO2forming pathway is most energetically favorable. Kurokawa et al. have reported that a cathode-supported fuel cell with a CeO2-infiltrated Ni−YSZ anode delivered PPDs of 200−240 mW cm−2 for 500 h under a current density of 0.4 A cm−2 at 700 °C when operated on H2 fuel that contained 40 ppm H2S, whereas the voltage of the cell with a conventional Ni−YSZ anode decreased to 0.0 V within 13 min.330 These results indicated that the addition of ceria nanoparticles increased the sulfur tolerance of the Ni-based electrodes. Similar results have also been reported by Yun et al.331 They found that the ceria reacted with H2S to form Ce2O2S according to reaction 25, which acted as a sulfur sorbent, whereas Ni3S2 was barely formed in the ceria-coated Ni−YSZ anode. Thus, the performance of the CeO2-infiltrated Ni−YSZ anode decreased to a lesser degree relative to that of the conventional Ni−YSZ anode. Yun et al. have also modified Ni-based anode with a thin coating of SDC on the pore-wall surface of the anode and compared its performance with that of pure ceria.332 The anode-supported fuel cells were tested in hydrogen with various H2S concentrations (0−100 ppm) at 600 and 700 °C. The cell performance improved by 20% and 50%, respectively, in the ceria- and SDC-modified anodes due to the introduction of additional TPB area in the anodes. Under the condition of 100 ppm H2S in the fuel gas, the performance decreased by 43% for the fuel cell with an unmodified anode, whereas it decreased by 30% for the fuel cell with the ceria-modified anode and only 14% for the fuel cell with the SDC-modified anode. A porous SDC layer on the Ni−YSZ anode pore wall acted as a sulfur sorbent and created additional TPB length, whereas ceria acted mainly as a sulfur sorbent at high concentrations of H2S (>60 ppm). Chen et al. have modified a conventional Ni−YSZ anode with 10 mol % Mo-doped CeO2 (MDC) using an impregnation method for SOFCs that operated on methane and H2S−H2 fuels.333 The anode modified with MDC showed very good sulfur tolerance with H2 fuel that contained 50 ppm H2S and good coking resistance toward methane because the cell stability for both methane and H2S-containing fuels was improved; the authors attributed this improvement to the enhanced electrical conductivity and TPB length. Zhou et al. have reported a CeO2-based anode doubly doped with Y and Yb for SOFCs and investigated its performance when operated on hydrogen that contained 5 ppm H2S.267 The authors found that the Ni−Ce0.8Y0.1Yb0.1O1.9 anode was tolerant toward H2S contamination, which led to good anode stability in hydrogen that contained 5 ppm H2S. The two rareearth-element dopants were believed to create a synergistic effect that improved the ionic conductivity as well as the possible conversion of H2S to SO2. These results suggest that co-doped CeO2 is an attractive anode for SOFCs operated on sulfur-containing fuels. This new material may also be used as a catalyst for the removal of fuel-gas contaminants, such as sulfur, before the fuel enters the fuel-cell reactor. Xu et al. have studied Ni−GDC anodes exposed to synthesized coal-syngas and H2 fuels with various concentrations of H2S at 800 °C under a constant current density with

sulfur tolerance of Ni−GDC cermet anodes, particularly when the H2S concentration was less than 100 ppm. The authors hypothesized that the enhanced sulfur tolerance of Pdimpregnated Ni−GDC anodes was most likely related to the weakened adsorption of sulfur on Ni and GDC phases due to the presence of Pd nanoparticles. However, the authors also showed that Pd impregnation did not completely inhibit the sulfur poisoning effect on the Ni−GDC anodes of SOFCs. The aggregation and growth of Pd nanoparticles were most likely the cause of the deterioration in the stability and the decrease in the sulfur tolerance of PdNi−GDC anodes. Sasaki et al. have performed calculations for materials that were thermochemically stable under a sulfur-containing reducing atmosphere (H2 gas with 5% H2O and 20 ppm H2S) at 800 °C.296 They found that select elements, including Ce, Y, La, Mg, Ti, Nb, Sc, Zr, Ca, or Al, were stable as oxides under such harsh operational conditions. Detailed information about the decrease in cell voltage that resulted from the H2S addition is shown in Figure 29. Some oxides were effective in reducing the decrease in cell voltage; however, some other oxides exhibited almost no influence. Although the effect should also depend on the morphology of the impregnated particles in the porous anodes, these results at least suggested that sulfur poisoning could be controlled to some extent through anode modifications via impregnation with appropriate oxides. Unfortunately, both the anode performance after the impregnation and the sulfur tolerance mechanism were not thoroughly investigated. Only the doped or undoped ceria oxides have been commonly used as sulfur-tolerant components in cermet anodes because of their good performance and low cost relative to those of the available alternatives. The use of CeO2 with Ni to form a cermet anode also effectively suppressed the sulfur poisoning of the nickel catalyst. Here, CeO2 acted as a H2S absorbent in fuel cells.317,327 The proposed possible reactions between the ceria and sulfur under a certain current density are listed as follows:328 CeO2 (s) + x H 2 → CeO2 − x (s) + x H 2O(g)

(24)

2CeO2 − x (s) + H 2S(g) + (1 − 2x)H 2(g) → Ce2O2 S(s) + 2(1 − x)H 2O(g)

(25)

CeO2 − x (s) + xO2 − → 2CeO2 (s) + 2x e−

(26)

Ce2O2 S(s) + 2O2 − → 2CeO2 (s) + SO2 (g) + 2e−

(27)

2Ce2O2 S(s) + 10O2 − → 2Ce(SO4 )2 (s) + 3CeO2 (s) + 20e−

(28)

Ce(SO4 )2 (s) + 2H 2(g) → CeO2 (s) + 2SO2 (g) + 2H 2O(g)

(29)

Chen et al. have studied the mechanism for H2S−CeO2 (111) interactions in SOFCs using periodic DFT calculations.329 DFT+U calculations were applied to properly characterize the effect on the interactions of the localization of the Ce4f states. The applied DFT plane wave calculations were implemented in the VASP with the PAW method. The GGA with the PW91 exchange-correlation functional was used. Adsorptions of H2S, SH, and atomic S were initially examined to probe energetically favorable intermediates. The species adsorb favorably at the Ce−top, O−top, and Ce−O bridging AE


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Figure 30. A conceptualization of the cell anode and electrolyte active interface showing that (a) the Ni is slowly oxidized by O2− at the active interface and (b) the GDC layer is suppressing the NiO formation at the active interface. Reprinted with permission from ref 328. Copyright 2011 The Electrochemical Society.

a ∼5 μm thick GDC barrier layer between the anode and the YSZ electrolyte.328 The results showed that the fuel cell with the GDC barrier layer was resistant to H2S at concentrations up to 1000 ppm in wet H2 and 100 ppm in syngas during longterm tests, whereas the fuel cell without the barrier layer exhibited a lower tolerance for the H2S impurity. Figure 30 shows a schematic diagram of the active interface between the anode and the electrolyte of the fuel cell. The use of H2 fuel with a certain level of H2S impurities resulted in the chemisorptions of sulfur onto the Ni particles of the anode. The diagram in Figure 30a shows that the Ni particles were oxidized by O2− at the TPB when sulfur was partially blocking the active sites on the Ni particles. If the H2 did not quickly reverse the formation of NiO, more NiO would be produced at the local area near the active interface. When the H2S impurity was present, H2 could only be adsorbed onto certain planes of nickel with a relatively high adsorption coefficient. On these sites, reaction 29 was still active; however, some active Ni sites were occupied by sulfur. Under a certain current density of 200 mA cm−2, the O2− can oxidize the Ni particles at the TPB, especially at the interface of the electrolyte and Ni−YSZ anode (Figure 30a). After the NiO is formed and accumulates to a certain level, the TPB could be damaged by the microstructural changes that result from Ni oxidation. The low conductivity of NiO and the loss of electrical and ionic contact between the Ni and the YSZ electrolyte could cause permanent damage and thereby increase the cell’s ohmic resistance. However, the GDC barrier layer (Figure 30b) can suppress the direct transport of O2− to the active interfacial area of the Ni or the TPB by partially consuming O2−. Under highly reducing conditions, CeO2 in the Ni−GDC and the GDC barrier layer was reduced to CeO2−x, which reacted with H2S to become cerium oxysulfide, Ce2O2S, at 800 °C via reaction 25.334,335 The adsorbed sulfur can partially block the local H2 oxidation active sites in the Ni−GDC anode; however, the electrochemical reactions, that is, reactions 26−28, can occur in the GDC barrier layer (Figure 30b), and then enhanced sulfur resistance would be obtained with the addition of the GDC barrier layer. Moreover, the H2 concentration could be further reduced via the hydrogen oxidation reaction in the GDC barrier layer. Choi et al. have modified the Ni−YSZ anode surface with niobium oxide (Nb2O5) to improve its sulfur tolerance.336 They found that Nb2O5 was reduced to NbO2 under SOFC operating conditions and that the NbO2 exhibited high electrical conductivity and catalytic activity toward hydrogen oxidation. The addition of NbOx improved the cell performance on dry

hydrogen fuel. The NbOx-coated Ni−YSZ cermet anode showed good sulfur tolerance when exposed to 50 ppm H2S at 700 °C over a period of 12 h. They found that different phases of niobium sulfides (NbSx) were formed on the surfaces of the niobium oxides. They attributed the observed excellent sulfur tolerance of the Nb2O5-coated Ni−YSZ anode to the good catalytic activity of the sulfides formed on the surface of the niobium oxide toward H2 oxidation in low concentrations of H2S. There are two possible mechanisms of sulfur poisoning, including the sulfur adsorption/desorption mechanisms and Ni oxidation/Ni3S2 formation mechanisms, which could be assigned to the different sulfur contents of 0−20 and >20 ppm. The combined use of new sulfur-tolerant anode materials and optimized operating conditions is effective to solve the sulfur poisoning problems assigned to the two different mechanisms. For the fuel with higher sulfur content, the desulfurization could be used before the fuel entered the fuel cell system, which could reduce the demand of the design of new sulfur-tolerant anode materials. On the basis of the advantages and disadvantages of current sulfur-tolerant materials, the combined application of available materials as different functional components in anodes through proper design may be effective to achieve a balance between stability and performance. Addition of more ionic conductive and sulfurreactive phases into Ni-based cermet anodes tends to be an easy and effective way to promote sulfur tolerance, but due to activity limitations of ionic phases, the composition and the contact of different phases need to be adjusted to mediate between sulfur tolerance and fuel-oxidation performance.

5. REDOX STABILITY During their long-term operation, SOFCs are expected to go through numerous redox cycles. As long as the fuel is supplied to the anode, the anode is expected to maintain a reduced state. However, metallic nickel in the anode could be reoxidized if air leaks through the electrolyte or if seals to the anode chamber of the fuel cell are imperfect because interruption of the fuel supply might occur accidentally as a result of an error in the system control or intentionally upon system shutdown and lead to a sponge-like NiO structure that occupies more volume than in the original oxidized state.337 An SOFC operated with undesirably high fuel utilization could also cause the oxygen activity to increase to greater than that at the equilibrium condition between Ni and NiO and result in the formation of NiO.338 The structural changes in the substrate and anode AF


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Figure 31. SEM study on microstructural changes in anode and substrate due to reduction/reoxidation cycles: (a) initial cofired state, (b) reduced, (c) reoxidized, (d) rereduced, and (e) rereoxidized. Reprinted with permission from ref 348. Copyright 2007 Elsevier.

Figure 32. Model illustrating the redox mechanism in Ni−YSZ cermets. The cartoon illustrates (a) the sintered state, (b) the short-term reduced state, (c) the long-term reduced state, and (d) the reoxidized state. The arrows in (d) point to a crack in the electrolyte and a failure in the ceramic network of the cermet, respectively. Reprinted with permission from ref 342. Copyright 2005 The Electrochemical Society.

based anodes because of their composite structure. Redox cycling on nickel-based cermet anodes for SOFCs has been studied through the testing of bars, discs, and powders, and the changes in weight, dimensions, mechanical integrity, and microstructure have been investigated using thermogravimetry, dilatometry, and microscopy.342−347 Among the various measurements, the change in microstructure is the most direct way to study the redox behavior of Ni-based cermet anodes. Malzbender et al. used micrographs taken from exactly the

upon reoxidation could lead to dimensional changes that generate stresses in the substrate, anode, and other cell components, which could potentially result in damage in all layers of the cell and therefore degrade the cell performance or even cause the complete failure of the fuel cell.338−341 5.1. Redox Behavior of Ni-Based Cermet Anodes

Although nickel oxidation and reduction have been extensively investigated, these results cannot be directly applied to nickelAG


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changed NiO network was believed to induce local tensile stresses in the YSZ network and to have subsequently caused failures in some places. The local failures in the network were accounted for the observed bulk expansion and increased porosity. These failures again gave rise to cracks in the electrolyte, which are illustrated in Figure 32d.

same location of an anode-supported half-cell in states that ranged from the cofired state to the state after the second reoxidation to study the reoxidation process.348 The initial dense NiO particles (Figure 31a) showed shrinkage upon reduction (Figure 31b). However, the microstructure after reoxidation was different from that of the as-prepared microstructure. The particles were fragmented in the reoxidized state and exhibited a spongy microstructure (Figure 31c). The reoxidized particles exhibited greater porosity (∼30%) and a larger volume than those in the as-prepared state. Furthermore, microcracks in the anode and electrolyte were visible after the reoxidation (Figure 31c). The re-reduction resulted in dense particles (Figure 31d), which appeared to occupy a larger volume than in the initial reduced state (Figure 31b) and therefore exhibited a coarser structure. After the second reoxidation, the NiO particles appeared to be even more fragmented (Figure 31e), and the electrolyte crack became larger. Macroscopically, the crack density was larger after the second reoxidation than after the first, which may be the result of the coarser structure of the Ni in the re-reduced state. Hatae et al. have also studied the initial damage to the anode microstructure of Ni−ScSZ caused by partial redox cycles during electrochemical oxidation.349 Partial electrochemical oxidation was used in the redox treatment and was tightly controlled at a very low degree to prevent catastrophic destruction of the cell. The authors found that the cell performance degraded and that microcracks formed at the interfaces between ScSZ and Ni in the anode layer, although no decrease in the OCV or mechanical damage to the electrolyte was observed. These results indicated that the microcracks in the anode layer were the initial damage caused by the partial redox cycles, which then led to the degradation of the cell performance. Heo et al. have also studied the redox behavior of a Ni−YSZ anode and the performance degradation under complete oxidation and reduction conditions.350 They found that the exposure time in oxidizing and reducing atmospheres played a critical role in the degradation of the structures and in the physical properties of the anode support. In particular, the redox cycling with 8 h of exposure time resulted in a destroyed YSZ network, which led to a significant degradation in the mechanical strength. The polarization experiments on the redox-cycled cell showed serious degradation of both the power output and the OCV. The authors hypothesized that the performance degradation resulted mainly from the microcracks across the whole cell and that interface delamination occurred in the anode. The redox mechanism of Ni−YSZ cermets was illustrated by Klemensø, as shown in Figure 32.342 The processes are illustrated for the 2D structure that consists of unisized particles shown in Figure 32a, which represents the as-sintered structure. In Figure 32, the YSZ particles are indicated by a hatched pattern, and the NiO particles are indicated by a plain and palegray color. After reduction, the phase change of NiO to metallic Ni occurred and resulted in a 25% volume reduction of the Ni phase. This reduction led to decreased nickel particle sizes and to reorganization of the phase. The changes in the nickel phase are illustrated in Figure 32b. However, the YSZ phase was believed to be unaffected by the first reduction process, as shown in Figure 32b. Nickel sintering resulted in increased particle sizes and the coarsening of grains, as shown in Figure 32c. After the reoxidization, the nickel phase expanded back to the as-sintered volume; however, the structural changes in the reduced state resulted in a different volume structure. The

5.2. Possible Solutions

On the basis of the previously discussed analysis, to ensure good geometric integrity of SOFCs with nickel-based anodes, intentional air break-in at the anode side of SOFCs should be avoided. If the ingress of air cannot be avoided due to technical limitations, then the amount of air should be strongly restricted, and the exposure time of the anode side surface should be strictly limited to reduce the amount of air that diffuses into the substrate and the anode to minimize the degree of oxidation of nickel. Vedasri et al. have developed a method to possibly minimize the mechanical degradation of Ni−YSZ anodes caused by redox cycling at high temperatures in air, and the viability of cooling a SOFC during air exposure was investigated as a possible solution.351 To prevent electrolyte and cell cracking, cooling of the anode-supported Ni−YSZ during air exposure from 800 to less than 600 °C at a rate greater than 3 °C min−1 was found to significantly slow the oxidation of nickel and to minimize the volume expansion due to NiO formation. A cell cooling rate of less than 3 °C min−1 resulted in cracking of the thin electrolyte layer because sufficient time was available for substantial formation of NiO. The oxidation during cooldown resulted in more extensive Ni oxidation in the outer regions of the anode layer as compared to the regions closer to the electrolyte. Recently, the development of redox-stable Ni-based anodes has attracted increasing attention. Buyukaksoy et al. have reported a redox-stable Ni−YSZ anode prepared using a polymeric-precursor-infiltration technique, by which a Ni−YSZ cermet microstructure with nanosized Ni particles coating the surface of porous YSZ was fabricated.352 Low-temperature processing of nanostructured Ni−YSZ cermets resulted in a reduction of the internal stresses between the Ni/NiO coating and the YSZ skeleton during the redox cycling process. An electrolyte-supported SOFC prepared with a Ni-infiltrated anode showed a PPD of 0.315 W cm−2 at 800 °C and a highly redox-stable anode (i.e., the reduction of power density was less than 1% after 15 redox cycles) in humidified 10% H2−90% Ar. Busawon et al. have also reported that Ni infiltration could be a potential solution to the redox problem of SOFC anodes.353 Kim et al. have reported a redox-stable NiO-coated YSZ composite anode synthesized via the Pechini process.354 Their microstructures and electrical properties have been investigated with thermal and redox cycling tests. The coverage of NiO crystals on the YSZ surface could be altered through control of the ratio between NiO and YSZ. A PPD of 0.56 W cm−2 was obtained at 800 °C in humidified hydrogen with this anode. This anode exhibited high homogeneity and abundant contact sites between the metallic Ni and the YSZ ceramic, which showed an excellent tolerance against thermal and redox cycling. The coating of Ni with reactive elements, such as Y, Zr, La, and Ce or their oxides, reduced the oxidation rate and changed the morphology of the oxide surface by changing the scale growth from primarily outward to primarily inward.355 Such coatings therefore have the potential to further reduce the oxidation degree of the substrate after an intended redox cycle. However, a suitable technology for coating the outer surface of AH


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most of the fuel and produce electricity. However, the authors’ EPOx device was a SOFC system in which methane was reformed within the anode chamber if the fuel flow rates were sufficiently high such that all of the syngas formed could not be electrochemically oxidized. The system can be designed and operated such that the exhaust stream is predominantly syngas, but some H2O, CO2, and possibly some unreacted hydrocarbons can also be presented. A schematic diagram of the major features of the button-cell experiments is shown in Figure 33, in which fuel supplied via a central tube flowed radically

the anode as well as its inner structure remains to be developed and established. Ju et al. have studied the effect of temperature on the reoxidation behavior of a Ni−Fe-anode-supported SOFC.356 After the oxidation and reduction treatments were performed for 2 h, the voltage did not return to the initial voltage at higher temperatures (500−700 °C). However, after reoxidation at 400 °C, the cell exhibited almost the same OCV as the as-prepared cell. During the reoxidation with air at higher temperatures, the Ni−Fe metal substrate exhibited two different expansion behaviors for the different oxidation rates of Ni and Fe, whereas the volumetric change of the oxidized substrate at 400 °C was negligible. The cell generated a power output of 160 mW cm−2 at 400 °C with 2.8 vol % H2O humidified hydrogen as the fuel, which was slightly decreased after the redox cycles; these results suggest that the addition of Fe could improve the redox behavior of the Ni-based anode, although it worked only at lower temperatures. It is reported that SrTiO3-based anode-supported SOFCs showed excellent redox stability.357,358 Very recently, SrTiO3based materials are also added to the Ni-based anodes to improve the redox stability. For example, Puengjinda et al. fabricated a composite anode that consists of yttria-doped SrTiO3 (YST) and SDC incorporated with NiO to evaluate the stability of the electrocatalytic performance and redox-cycling behavior relative to those of conventional Ni−YSZ anodes.359 The cells with Ni−YST−SDC anodes exhibited more stable performance than conventional Ni-based cermets during longterm operating tests under severe conditions in methane fuel with low steam-to-carbon ratio (S/C = 0.1) and highly humidified hydrogen fuel (40 vol % H2O−H2). Thus, the authors considered the Ni−YST−SDC composite to be a promising candidate as an SOFC anode material that is capable of operating under severe conditions. However, the most useful method to increase the redox-cycling stability is the replacement of Ni-based anodes with pure oxide anodes, such as perovskite-based materials.

Figure 33. Schematic representation of the EPOx button-cell experiment. Reprinted with permission from ref 365. Copyright 2008 Elsevier.

across the anode face of the membrane-electrode assembly (MEA). Air is supplied to the cathode side of the MEA. The button-cell experiments were used to establish representative parameters for the MEA structure that were used in the tubular model.370,371 The authors found that many of the relevant parameters were physical dimensions, which were easily established. Coke formation is an important consideration when Ni surfaces are exposed to hydrocarbons at high temperatures, and a barrier layer was proposed to inhibit carbon deposition. The results show that a tubular cell can be designed to deliver syngas and electricity using methane as the primary fuel. Furthermore, numerous designs and operating alternatives can be used to optimize the performance. Zhang et al. have reported the cogeneration of electricity and syngas in a SOFC system.366 A PPD of 90 mW cm−2 was obtained in an electrochemical cell that consisted of a Ni−SDC anode and LSGM electrolyte for the conversion of methane to syngas at 800 °C. The cell, after being operated at 50 mA cm−2 for 24 h, showed no obvious degradation of performance. The authors found that the syngas cogenerated at a H2/CO ratio of 1.4−2.0 varied with the applied current density, the CH4 flow rate, and the operating temperature. The anode potential was found to dominate the selectivity of products, and a low anode polarization facilitated the formation of syngas at high H2/CO ratios. Sobyanin et al. have also presented a cogeneration system in a SOFC reactor with a Ni- or Pt-based anodes to convert methane. They found that the anodes showed high electrocatalytic activity toward the formation of syngas from methane; however, the fuel cell generated a very low power output.372 Most prior reports of cogeneration systems with SOFCs have reported relatively low power densities and low syngas

6. COGENERATION OF ELECTRIC POWER AND SYNGAS In a direct-methane SOFC, the methane is generally electrochemically oxidized to CO2 during the power generation process. In some cases, however, the methane can also be partially and deliberately oxidized by the O2− transported from the cathode side through the electrolyte to value-added products, such as synthesis gas (syngas). Such a process is called electrochemical partial oxidation (EPOx), and the fuel cell can be considered an electrochemical membrane reactor. Syngas is an important precursor to hydrogen and synthetic liquid chemicals/fuels, including methanol and various hydrocarbons.360−362 The advantages of this cogeneration system are similar to those of conventional ceramic-membrane reactors, including the ability to produce syngas without nitrogen dilution and at reduced cost due to process intensification through the combination of the oxygen-separation and partialoxidation steps. Recently, both the theoretical363−365 and the experimental aspects366−369 of the cogeneration of electricity and syngas production in a SOFC system have attracted increased attention. For example, Zhu et al. have developed a computational model to investigate the characteristics of an EPOx system that was intended to produce both electricity and syngas.365 SOFC systems are typically operated to consume AI


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Figure 34. The fuel cell reactor system used for synthesis gas and electric power cogeneration from methane. Reprinted with permission from ref 369. Copyright 2011 John Wiley and Sons.

integrated with a catalyst (GdNi/Al2O3) for the partial oxidation of methane in the same gas chamber was used for the facile cogeneration of electricity and syngas from methane with zero waste gas emissions (Figure 34). When a methane− oxygen mixture (CH4/O2 molar ratio 2:1) was used as the feed gas, a high cell power output, a high methane conversion, a high H2 and CO selectivity, and an ideal H2-to-CO molar ratio of approximately 2.0 were achieved simultaneously. The authors demonstrated that the fuel cell with a bilayer electrolyte (SDC and YSZ), a tape-cast Ni−YSZ anode, and a Ba0.5Sr0.5Co0.8Fe0.2O3‑δ (BSCF) cathode delivered an OCV of 1.07 V and a PPD of 1500 mW cm−2 at 700 °C when operated on methane−oxygen gas mixtures with a ratio of 2:1, and no coking was observed. The authors found that both the syngas formation rate and the H2/CO molar ratio were unaffected by the polarization current density, which is a significant advantage with respect to practical applications.369 However, the concept is flawed because introducing the oxygen or steam into the methane stream turns it into a catalytic membrane reactor, and the electricity generated by the SOFC is just a minor side reaction. On the other hand, this concept can be considered as an internal reforming SOFC with extremely low fuel utilization. Furthermore, the performance results are exaggerated by the heating of the cell that occurs due to the combustion of the CH4−O2 gas mixtures and that the real temperature of the fuel cell is not well-controllable. The stability of the system should be investigated in the future.

production rates. Zhan et al. have demonstrated that a Ni− YSZ-anode-supported SOFC operated on pure methane fuel could generate high power output and high syngas production rates.367 Fuel cells operated at T ≈ 750 °C, V ≈ 0.4 V, and O2−/CH4 ≈ 1.2 yielded an electrical power output of ∼0.7 W cm−2 and a maximum measured syngas production rate of ∼20 mL min−1 cm−2. The authors also found that the H2-to-CO ratio of the produced syngas was strongly dependent on the polarization current density. A stable SOFC power output for up to 300 h was reported; however, the chemical products as a function of time were not measured. Pillai et al. have also reported a direct-methane SOFC for the simultaneous production of electricity and syngas.373 The fuel cells produced a PPD of 0.9 W cm−2, and a methane conversion of ∼90% and a syngas formation rate of 30 mL min−1 cm−2 when operated on methane fuel at 750 °C. However, the methane conversion decreased continuously during the first 30−40 h, even though the electrochemical performance of the SOFC was stable. This decay in methane conversion resulted from the reduction in the reforming catalytic activity of the Ni−YSZ anode due to Ni sintering. Initial tests showed that an added catalyst could improve the methane conversion stability; however, additional research is needed to demonstrate long-term-stability of EPOx in SOFCs. The previously mentioned syngas−electricity cogeneration systems were all based on a dual-chamber configuration, where pure methane was fed to the anode chamber and oxygen for the partial oxidation was transported from the cathode.374,375 The syngas formation rate and the methane conversion are all strongly influenced by the polarization current.367,368 Furthermore, coke was easily formed over the nickel-based cermet anodes, especially under a low polarization current density. Given the drawbacks of the dual-chamber configuration, SCSOFCs should be an ideal way to cogenerate syngas and electricity because carbon deposition is not a serious problem due to the high oxygen content in the gas mixtures. A previous study showed that the cogeneration of electric power and syngas can be realized in a SC-SOFC constructed with Ni− YSZ|YSZ|Au.376 However, an extremely low PPD was obtained, and the conversion of methane was also inevitably influenced by the polarization current. Recently, Shao et al. have reported an energy-efficient combination of a SC-SOFC and a downstream catalytic equilibration system.369 A SC-SOFC

7. SOFCs OPERATING ON OXYGENATED METHANE FUELS As the simplest hydrocarbon, methane possesses only C−H bonds. From a structural point of view, DME and methanol can be considered oxygenated methane because it also contains no C−C bonds. Although neither DME nor methanol can be synthesized from the direct oxidation of methane, they can be synthesized indirectly via syngas, a mixture of CO and H2. Because methane can be converted to syngas through steam reforming, methanol could then be synthesized from the syngas through a catalytic reaction, whereas DME can be further synthesized from the catalytic dehydration of methanol. Methanol is a liquid at room temperature and ambient pressure, whereas DME can also be easily liquefied at room temperature under moderate pressure. Thus, the transportation AJ


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Figure 35. SEM photos and EDX profiles of a fresh Ni−YSZ anode (a and b), after treatment in pure DME (c and d), and after treatment in DME− CO2 (DME:CO2 = 1:1) (e and f) under open circuit conditions at 850 °C for 30 min. Reprinted with permission from ref 377. Copyright 2011 Elsevier.

fuels for PEMFCs via direct oxidation at low temperatures. Because of the presence of oxygen in the DME and methanol molecular structures, significant advantages, such as very low concentrations of sulfur impurities and a tendency to not form coke (based on thermodynamic calculations), can be achieved when DME and methanol are used as fuels in fuel cells. Therefore, the use of methanol and DME as fuels for SOFCs has also attracted attention.

and storage of DME and methanol are much easier than the transportation and storage of methane. Because both methanol and DME can be synthesized from syngas, and because syngas can be obtained from a wide range of chemicals, including natural gas, coal, and biomass, the wide accessibility of methanol and DME is expected if mass production could be realized. Therefore, the application of methanol and DME as fuels has also received considerable attention recently. To this point, methanol and DME have been investigated primarily as AK


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not sufficient at temperatures less than 750 °C, which implies that the operation of a cell on DME−CO2 gas mixtures is possible only in a narrow temperature window of 750−850 °C. At present, there is increasing interest in the reduction of the operating temperatures of SOFCs to extend the cell lifetime and achieve cost effectiveness. As compared to CO2, O2 is much more reactive and oxidative; thus, Su et al. further investigated the operation of SOFCs on a DME−O2 mixture through internal partial oxidation with Ni-based anodes.379 They demonstrated that the presence of oxygen promoted the decomposition of DME into primarily H2, CO, and CH4. This promoting effect was due to the reduced activation energy (Ea) required for DME decomposition in the presence of oxygen. This viewpoint is congruent with the primary results by Murray et al.382 Carbon deposition did occur on the anode’s outer surface during the stability test when the cell was operated on a DME−O2 mixture at 700 °C; however, no carbon was formed inside the anode layer, which ensured stable cell performance. They also found that maximum carbon deposition occurred at 700 °C in the investigated temperature range of 600−800 °C; therefore, the elimination of deposited carbon at 700 °C is still a major challenge for increasing the performance of SOFC using DME as the fuel. In addition to thermodynamics, reaction kinetics also affected the coke formation of cells that operate on DME. This finding suggests that modification of reaction kinetics may be another route to improve the coking resistance as well as the cell performance of SOFCs that use DME as fuel. Some supported catalysts have been reported to show very high catalytic activity toward the partial oxidation of DME to syngas at intermediate temperatures. 383−385 For example, Ni supported on La0.8Sr0.2Ga0.8Mg0.15Co0.05O3 (LSGMC5) was found to be a good catalyst for the partial oxidation of DME to CO and H2.384 Further alloying of the nickel in the catalyst with iron in the development of Ni−Fe−La0.8Sr0.2Ga0.8Mg0.115Co0.085O3 (LSGMC8.5) alloy anodes resulted in a cell that demonstrated improved coking resistance when fed DME fuel directly.232 Among the various supported catalysts, a mixture of Pt/Al2O3 and Ni/MgO showed particularly high catalytic activity toward the partial oxidation and reforming of DME between 600 and 750 °C.385 Recently, Su et al. directly deposited such a catalyst onto a nickel cermet anode surface to function as a functional catalyst layer for improving the performance of SOFCs that operate on DME−O2 gas mixtures.386 The results demonstrated that the Pt/Al2O3−Ni/MgO catalyst exhibited significantly greater resistance against coke formation than a sintered Ni−YSZ anode when exposed to DME−O2 gas mixtures; in particular, almost no coke formation was observed on the mixed catalyst when a stoichiometric ratio between DME and O2 (2:1) was used. Meanwhile, the mixed catalyst also exhibited high activity that was comparable to that of a Ni−YSZ anode for the partial oxidation of DME. Therefore, the SOFC with a Pt/Al2O3−Ni/MgO anode catalyst layer exhibited favorable power outputs when DME−O2 mixtures were used as the fuel; these power outputs were only slightly lower than those achieved during the operation on hydrogen. Similarly, Yano et al. have applied a Ru/Ni/SDC catalyst layer over an anode to promote the partial oxidation of DME in a SOFC that operates in a single-chamber mode; effective increases in the OCV and in the power output of the fuel cell were observed after the catalyst layer was introduced.387 On the basis of the previously discussed fundamental research, the direct utilization of DME for SOFCs with

7.1. Dimethyl Ether

The direct operation of SOFCs with DME as fuel was first reported by Murray et al. In contrast to methane, DME is unstable under the operating temperatures of SOFCs because thermal decomposition of DME begins at temperatures greater than 600 °C.40 Thus, the oxidation of DME over a SOFC anode likely occurs via the indirect oxidation mechanism, even with perovskite oxide as the anode material. Murray et al. predicted the decomposition reaction of DME fuel using kinetic modeling and verified the reaction experimentally; H2, CH4, and CO were found to be the major decomposition species.40 At operating temperatures greater than 650 °C, most DME has thermally decomposed before reaching the anode catalyst; therefore, in such cases, the cell actually operates on a mixture of H2, CO, and CH4 when DME is used as the fuel. The nickelbased anode further catalyzed the decomposition of DME at lower temperatures (550−650 °C). However, in the decomposition products, the molar fractions of both H2 and CO were found to be higher than that of CH4 at T < 700 °C, which suggests that solid carbon formed and most likely came from CH4.40 Interestingly, carbon deposition was not observed over the anode of the tested cell. The authors hypothesized that the oxygen-ion current was sufficiently high to oxidize any carbon deposited at the anode while the cell operated.40 In principle, under OCV conditions, no oxygen transports from the cathode to the anode, and then the accumulation of solid carbon is likely to occur. Especially for long-term operation, coke formation continues to be a practical problem. In practice, many researchers have found that the formation of coke occurs on Ni-based anodes when DME is directly used as the fuel in SOFCs.377,378 For example, Su et al. have reported the apparent formation of coke on a Ni−YSZ anode after it was treated in a pure DME atmosphere under OCV conditions, and greater carbon deposition was observed at higher operating temperatures.377 Although the cell delivered comparable power densities when operated on H2 and DME, the cell performance obviously degraded during the stability test because of the carbon deposition when the cell was operated on pure DME. As previously mentioned, coke formation is still a major problem when DME is used directly as the fuel in SOFCs with Ni-based anodes. The coke formation was suppressed thermodynamically through the addition of an oxidant to DME.377,379,380 Faungnawakij et al. calculated the thermodynamic equilibrium of the steam reforming of DME at the operating temperature of SOFCs using a Gibbs free energy minimization technique and theoretically studied the carbon formation boundary.381 During the calculation, they assumed that the carbon product was graphite. The authors demonstrated that the coke formation could be avoided by an increase in the S/C ratio and/or the reforming temperature. When DME, MeOH, H2O, H2, CO, CO2, and C were considered as the products, the total conversion of DME without carbon deposition was realized when the S/C ratio was greater than 2.5 at 900 °C, whereas the coke formation was greatly suppressed when CH4 was considered to be present. Su et al. have studied the effect of the introduction of CO2 into the DME fuel for SOFCs with Ni-based anodes.377 They found that the introduction of CO2 could effectively suppress the coke formation, especially at high temperatures (>750 °C), as shown in Figure 35. This suppression is due to the high catalytic activity of the Ni-based anode for the CO2 reforming of DME and CH4 and the gasification of deposited solid carbon. However, the catalytic activity of Ni-based anodes is AL


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previous authors differ from each other, and these differences affect the methanol catalytic decomposition and the rate of electrochemical oxidation of the decomposed products (CO +H2). More systematic work is thus needed to clarify the performance−microstructure relationship of SOFCs operated on methanol fuel. Coking is a general problem for SOFCs with nickel cermet anodes that operate on carbon-containing fuels. Because of the high oxygen content in the methanol structure, SOFCs with conventional nickel cermet anodes can still be operated stably on methanol fuel. For example, Liu et al. demonstrated that a cell with a nickel cermet anode operating on methanol fuel under a constant polarization voltage of 0.5 V did not exhibit a noticeable decline in cell voltage for a period of 60 h.393 Sun et al. have also reported the stable operation of a SOFC on methanol fuel with no carbon deposition over the anode in a direct-flame fuel cell configuration for a period of 30 h.396 Notably, coke formation is still favorable at the operating temperatures of SOFCs when methanol is used as the fuel thermodynamically;397−400 however, the coking may be effectively prevented through the addition of water to methanol, through an increase in the operating temperature, or under conditions of polarization current.41,397 For example, Saunders et al. demonstrated that a small amount of carbon deposition was produced on a Ni−YSZ anode when an SOFC was operated on methanol.401 Carbon deposition increased at lower operating temperatures in the range of 700−850 °C; however, such deposited carbon could be inhibited by the introduction of water or air into the methanol fuel. Assabumrungrat, Laosiripojana, and their co-workers have also devoted themselves to the operation of SOFCs on methanol fuel through internal reforming.189,400,402,403 Assabumrungrat et al. compared the H+- and O2−-conducting electrolyte systems and showed that coke formation in the latter system was less likely because of the excess H2O produced by the electrochemical reaction of H2 in the anode chamber.400 With respect to the methanol-fueled SOFC with an H+-conducting electrolyte, the cell performance can be maximized through the use of a coflow pattern and a high operating pressure. The authors also found that high-surfacearea nanoparticulate CeO2403 and Ni/Ce−ZrO2189 are good catalysts for the decomposition and reforming of methanol in SOFC applications. Leone et al. studied the feasibility of operating a cell with a Ni−YSZ anode under the direct reforming of methanol from an experimental point of view and observed that the carbon deposition did not occur in an SOFC operated at 800 °C with a methanol−H2O mixture when the ratio of methanol to water was 1.5 or less; good cell performance was also observed.404 Gao et al. used steam/ methanol in a ratio of 1/1 for the low-temperature operation of an SOFC based on the thermodynamic calculations, and their fuel cell with a lithiated Ni−Cu−ZnO anode exhibited excellent performance, with a PPD of 431 mW cm−2 at 500 °C.399 Similar to fuel cells that operate on other hydrocarbon fuels, the coking from methanol fuel can also be suppressed through the use of non-nickel-based anode materials.405−408 For example, Brett et al. found that SOFCs with ceria-impregnated Cu anodes provided reasonable power outputs when operated on dry methanol and exhibited no detectable carbon deposition at intermediate temperatures (500−600 °C).405 However, the Cu-based catalysts exhibited poor activity toward H2 oxidation and were prone to thermal sintering. Zhu and co-workers subsequently developed a series of Ni−Cu-alloy anode catalysts

nickel-based anodes is possible if the coke formation is effectively suppressed, which can be realized via various methods. However, all of the previously discussed investigations were conducted on single cells, whereas practical SOFCs should be constructed in stacked configuration to allow sufficient power and voltage output. Recently Sato et al. focused on the power generation properties of SOFC stacks that operated on DME. The fuel was supplied to a conventional SOFC with a nickel-based anode via external steam reforming.388 They used Ni/Al2O3 as the reforming catalyst, and carbon deposition over the surface of the catalyst was not observed when DME steam reforming was performed at a S/C ratio of 1.5. Although the authors used external reforming of DME in their study, the deposition of Ni/Al2O3 directly onto the anode surface as the catalyst layer for internal reforming of DME allows the SOFC stack to also be directly operated on DME. Thus, the use of DME as a fuel for SOFCs is promising for practical applications. 7.2. Methanol

Methanol has been widely investigated as a fuel for lowtemperature fuel cells, which are considered portable power sources because of their high energy density, ease of storage and transportation, and wide availability. Recently, the direct operation of SOFCs on methanol has also attracted attention. As compared to DME, the oxygen content in the molecular structure of methanol is greater; thus the nickel cermet anode is less susceptible to coke formation when operated on methanol fuel. Similar to that of DME, the mechanism of methanol electrochemical oxidation over an SOFC anode likely involves an indirect pathway due to the ease with which methanol undergoes thermal decomposition. Because of its sulfur-free nature, the problem of sulfur poisoning of the anode is avoided. Sahibzada et al. have reported the fabrication of SOFCs with a Ni−YSZ-anode-supported GDC electrolyte that operate on methanol at intermediate-to-low temperatures; their primary results showed a PPD of only 25 mW cm−2 at 600 °C with methanol fuel, which was approximately one-half of that achieved with H2.389 The addition of some Pd catalyst to the Ni−YSZ anode improved the cell performance somewhat,390 which suggests that the anode might have some kinetic limitations for the internal reforming/decomposition of methanol or the electrochemical oxidation of CO and H2. Similar results were also reported by Babaei et al.391 Interestingly, even without any modification, very high power output was sometimes obtained from cells with conventional nickel cermet anodes that were operated on methanol fuel. For example, Jiang et al. reported an SOFC with high performance, with a PPD of 600 mW cm−2 at 650 °C;392 Liu et al. also reported very high power output for an SOFC with a Ni−SDC anode operated on methanol fuel at intermediate temperatures and obtained a PPD of 698 mW cm−2 at 650 °C.393 Recently, Meng et al. have reported that anode-supported SOFCs with graded porous Ni−SDC anodes delivered a PPD of 0.82 W cm−2 at 600 °C when operated on methanol fuel.394 Lo Faro et al. applied a new anode catalyst, Ni-modified La 0.6Sr0.4Fe0.8Co0.2O3 (LSCF), and obtained a PPD of approximately 350 mW cm−2 with an electrolyte-supported SOFC when they used both syngas and pure methanol at 800 °C.395 These results suggest that the performance of SOFCs that operate on methanol depends not only on the anode materials but also on the anode microstructure and/or cell configuration. The anode microstructures reported by the AM


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With respect to the electrochemical oxidation of methane over SOFC anodes at elevated temperatures, two mechanisms are available: direct electrocatalytic oxidation to form CO2 and H2O and indirect oxidation via syngas (CO+H2). When nickel cermet anodes are used, the reaction is conducted mainly through the indirect mechanism. However, the actual reaction pathways are still far from being well understood, and more research is needed to obtain a clearer picture. Because of the complicated electrode microstructures, which complicate the reaction process, model electrodes based on pattern or point electrodes or thin-film electrodes are preferred to obtain more in-depth understanding of the electrode reactions. For example, numerical modeling may provide useful information about the methane electrochemical reactions over the nickel cermet anodes and the effects of electrode modification on the reaction pathways; such modeling therefore deserves greater attention in future studies. Although methane is the simplest hydrocarbon, the knowledge gained from direct-methane SOFCs may provide useful guidance for understanding the mechanism and kinetics of the electro-oxidation of other hydrocarbons or oxygenated hydrocarbons over SOFC anodes. Coking is the main limitation of nickel cermet anodes that operate on methane or related fuels, and numerous strategies have been developed to improve their coking resistance. Although the addition of O2, CO2, and H2O to methane in the fuel gas can increase the O/C ratio to suppress coke formation thermodynamically, notably, a large amount of oxidant may decrease the power density and increase the probability of oxidizing the SOFC anode. Therefore, the amount of such gas(es) should be minimized. The addition of hydrogen can also mediate the problem of coking; however, a high H2-tomethane ratio is required to completely avoid coke formation thermodynamically. This approach is therefore not practical. Instead, we believe that the addition of a basic gas, such as ammonia, may be a potential way to suppress coke formation. The acidic sites on the anodes are well-known to promote the hydrocarbon cracking reactions. A basic gas could neutralize the acidity of the Ni-based anodes through preferential adsorption and thereby reduce the formation of coke. This approach is particularly interesting because SOFCs with conventional nickel cermet anodes may still be used without significant modification of the cell materials. Furthermore, the addition of ammonia would not decrease the efficiency of the fuel cells because ammonia itself could also be used as a fuel for SOFCs. The fabrication of alloy anodes and the modification of nickel cermet anodes with other metals or oxides have provided some potential approaches to suppress coke formation. We expect that an increasing number of new potential anodes based on these strategies will appear in the future. In addition, the composition and microstructure of these new anodes also need further optimization to maximize their electrochemical performance. Because Sn is a good promoter for Ni-based anodes to suppress coke formation, some other group-IVA metals with similar properties could also be used as promoters to form alloys. Basic supports are known to reduce coke formation; thus, the addition of alkali elements such as Li, Na, and K to Ni-based anodes is expected to suppress coke formation by increasing the basicity of the anode. However, such alkali elements could be easily lost through vaporization at elevated temperatures; thus, performance degradation is expected during long-term operation. Some composite oxides are well-known to exhibit Li+, Na+, or K+ conductivity. Some oxides even exhibit high lithium-ion conductivity at room

with improved sintering resistance, which were not only active toward methanol decomposition but were also resistant to coke formation.399,409−412 For example, Feng et al. synthesized a novel anode material, C−MO−SDC (C = activated carbon (AC)/carbon black, M = Cu, Ni, or Co, SDC = Ce0.9Sm0.1O1.95);409 the carbon materials improved the microstructure, electronic conductivity, and catalytic activity of the anode. An SOFC with an AC−MO−SDC anode and a ceria− salt composite (CSC) electrolyte achieved a PPD of 250 mW cm−2 during direct operation on methanol fuel at 560 °C. Xu et al. used vapor grown carbon fiber (VGCF) as the carbon material in a C−MO−SDC anode.410 A cell with a 1.25 wt % VGCF−MO−SDC anode obtained a PPD of 258 mW cm−2 when fed 3 vol % H2O−methanol; this result was similar to that of a cell that contained a 1.0 wt % AC−MO−SDC anode, which produced a PPD of 257 mW cm−2. However, the cell with the VGCF−MO−SDC anode performed better in the medium-current-density range. Feng et al. and Xu et al. all noted that the operational stability of SOFCs with C−MO− SDC was not sufficient, possibly because the carbon materials were destroyed at temperatures greater than 500 °C during long-term operation.409,410 Thus, Raza et al. recently developed an advanced SOFC with a functional nanocomposite Li, Cu, Ni, and Zn electrode (LCNZ) in a molar ratio of 4:6:15:7 and a nanocomposite electrolyte of Na2CO3 and Sm-doped CeO2 (NSDC).411 A symmetric electrolyte-supported cell with a configuration of LCNZ+NSDC| NSDC| LCNZ +NSDC exhibited a high PPD of 600 mW cm−2 at 550 °C when operated on biomethanol. This result demonstrates that the advanced SOFCs exhibit good chemical stability within the selected temperature range of 500−530 °C when fueled with biomethanol. No coke formation was observed on the anode after the SOFC was operated for 72 h. The utilization of functional nanocomposite anodes is a strong contribution to the development of intermediate-to-low-temperature SOFCs. In summary, coke formation is not a serious problem for SOFCs that operate on methanol, as compared to those that operate on DME and other hydrocarbons, because of the high oxygen content in methanol’s molecular structure. The ease of decomposition of methanol also allows a high cell power output at intermediate temperatures, even for cells with conventional nickel cermet anodes. The storage and transportation of liquid methanol in its liquid state at ambient conditions is facile; however, a gasification system is generally required to deliver the methanol fuel to the SOFC reactor. A carrier gas is also required for feeding liquid fuels into a fuel-cell reactor, which will dilute the fuel gas and reduce its volumetric energy density. In contrast, DME exists as a gas at normal conditions and is easily liquefied at room temperature. Therefore, the use of a fuel in the gaseous state (e.g., CH4 or DME) as a carrier gas to bring a fuel in the liquid state to the fuel cell anode may be an alternative and facile way to further improve the performance of fuel cells.

8. CONCLUSIONS AND PROSPECTS In this Review, we have highlighted recent studies that describe the mechanisms, challenges, and new approaches for improving the performance of SOFCs with nickel-based anodes that operate on methane and related fuels. Although great progress has been made during the past few decades, considerable research is still needed to realize the widespread application of this technology in our daily lives. AN


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the Ni-based anodes is still required to improve the sulfur tolerance. Redox instability is also a potential problem with nickel cermet anodes. Methods similar to those used to modify anodes to increase their coking resistance can also be applied to increase the sulfur tolerance of nickel-based anodes and their redox stability. However, to this point, no studies have been conducted that have addressed all three problems simultaneously. In the future, anode modification strategies that could result in simultaneous increases in coking resistance, sulfur tolerance, and redox stability should be developed to advance direct-methane SOFCs. Oxygenated forms of methane, such as methanol and DME, are also potential fuels of the future and are superior to methane because of their high volumetric energy density and easier storage and transport. The presence of oxygen in the molecular structure of these hydrocarbons facilitates the thermodynamic suppression of coke formation. We expect that additional research will be reported in the near future concerning the generation of power from these fuels via SOFC technology. One substantial challenge with respect to the development of SOFCs with nickel-based anodes that operate on these fuels is the redox stability of the anodes because the presence of oxygen in the molecules can increase the risk of reoxidation of metallic nickel. Special attention is therefore required to improve the redox stability of the anodes through modification.

temperature, and these materials have been widely used in lithium-ion batteries as electrodes or electrolyte materials. Similar materials could be developed as anode components of SOFCs. The loss of surface alkali elements can be compensated by the diffusion of those cations from the bulk of the oxides, thereby ensuring a long operational stability under fuel-cell operating conditions. The substitution or doping of ceramic phases into Ni-based anodes was found to be another effective way to improve the coking resistance, although the literature contains few reports about this approach. The challenge is that the ceramic phase should typically exhibit a high ionic conductivity to maximize the TPB length for the electrode reaction. Thus, the selection of ceramic phases is highly limited. Interestingly, some single oxides, such as Sm2O3, have been reported to show good electrode performance, which was believed to result from their capability to spillover hydrogen. However, more research from different groups is needed to achieve more in-depth understanding of the mechanism. The easy hydration of lanthanide oxides and their reactions with conventional electrolytes cast doubt on the long-term operational stability of such systems. In addition, alkali metals could react with other fuel cell components and then lead to a cell degradation, which is a disadvantage of this approach. The substitution of a ceramic phase in the anode (stabilized zirconia or doped ceria) for oxygen-ion-conducting SOFCs with protonic conductors, such as barium zirconate, with a high water-absorption capacity may be a good approach to reduce coke formation. This is because the strong water-absorption capability of the protonic conductors may increase the local O/ C ratio on the surface of the anode catalyst and thereby suppress the coke formation thermodynamically. Interestingly, such proton-conducting oxides also exhibit high basicity and strong interactions with the nickel phase; thus, the coking resistance could be further improved. Furthermore, Su et al. recently demonstrated that such proton conductors can react with doped ceria electrolytes to form an electron blocking layer over a SDC surface. As a result, the OCV of cells with thin-film SDC electrolytes was also improved.413 We expect that an increasing number of anodes for direct-hydrocarbon SOFCs and catalysts for the steam reforming of hydrocarbons with protonic conductors will be reported in the near future. Fuel cells with an anode catalyst layer could operate on methane directly without any external reforming. This concept can be applied to any type of hydrocarbon fuel through the selection of an appropriate catalyst for use in the catalyst layer and the selection of an appropriate operating temperature, which would allow the design of flexible SOFC systems for various applications. In addition, conventional SOFC materials and configurations could still be used, thus offering great economical attractiveness. This approach therefore requires more attention in the future. The challenge is that the catalyst layer should possess sufficient electrical conductivity to ensure good current collection. The primary results from Wang et al. suggest that the addition of inert copper into the catalyst layer may be a potential solution.256 Sulfur poisoning is another important issue for the operation of SOFCs on hydrocarbons because many of the fuel sources contain sulfur impurities. Pre-desulfurization is the most convenient solution to this problem; however, it is somewhat difficult to decrease the sulfur content of fuels to less than 10 ppm, which is required for stable operation of a SOFC with conventional nickel-cermet anodes. Thus, the modification of

AUTHOR INFORMATION Corresponding Author

*Phone: +86 25 83172256. Fax: +86 25 83172242. E-mail: [email protected]. Notes

The authors declare no competing financial interest. Biographies

Wei Wang was born in Yangzhou, China in July, 1986. He completed his Ph.D. in Chemical Engineering at Nanjing University of Technology in June, 2013. He will go to Curtin University, Australia, as a postdoctoral fellow in 2013. His research interests include anode materials and anode catalyst layer for SOFCs operating on hydrocarbons, the study of fuel varieties for SOFCs, and the hydrogen production from hydrocarbons reforming for fuel cells. AO


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Chinese Academy of Sciences, in 2003. She was a visiting scholar at the Petroleum Institute of Catalysis Madrid, Spain. She joined the College of Chemistry and Chemical Engineering at Nanjing University of Technology in 2005. Her main research interests include electrode materials of lithium-ion batteries, hydrogen production process from hydrocarbon, and anode catalytic materials for SOFCs.

Chao Su was born in Dalian, China in 1987. She completed her Ph.D. in Chemical Engineering at Nanjing University of Technology in 2012. She is currently a postdoctoral fellow at Curtin University, Australia. Her research activities are mainly focused on the systematic study of SOFCs utilizing oxygenated chemicals fuels, mechanism investigation of carbon deposition over the anode of cells, development and characterization of new anode materials, and optimization of performance for SOFCs based on doped-ceria electrolytes.

Zongping Shao received his Ph.D. from Dalian Institute of Chemical Physics, Chinese Academy of Sciences, in 2000; after that he was a visiting scholar at the Institut de Researches Sur La Catalyse, CNRS, France and a postdoc at Materials Science, California Institute of Technology, CA from March 2002 to June 2005. In July 2005, he joined Nanjing University of Technology where he was promoted to professor. Since then he has been the director of institute of new energy materials and technology. His research interests include SOFCs, oxygen permeable membranes, polymer electrolyte membrane fuel cells, lithium-ion batteries and supercapacitors. He has published over 230 international journal papers with total citations >5000.

ACKNOWLEDGMENTS This work was supported by the “National Science Foundation for Distinguished Young Scholars of China” under contract no. 51025209.

Yuzhou Wu was born in Taizhou, China in July, 1986. She started her graduate research on carbon fueled SOFCs and the optimization of SOFC anode materials in Nanjing University of Technology from 2008, under the supervision of Prof. Zongping Shao. She carried out her research on the optimization of SOFC anode structures in Prof. Huanting Wang’s group, Monash University, as a visiting student from 2010. She is now working on the fabrication and optimization of microscale SOFCs for portable power generation as a Ph.D. candidate at Monash University, under the supervision of Prof. Huanting Wang.

ABBREVIATIONS BET Brunauer−Emmett−Teller BPG biomass-produced gas BSCF Ba0.5Sr0.5Co0.8Fe0.2O3‑δ BZY yttrium-doped barium zirconate CCSD coupled cluster singles and doubles COS carbonyl sulfide CSC ceria−salt composite CZO Ce0.75Zr0.25O2 DFT density functional theory DME dimethyl ether DRIFTS diffuse reflectance Fourier transform infrared spectroscopy Ea activation energy EDS energy-dispersive X-ray spectroscopy EIS electrochemical impedance spectroscopy EPD electrophoretic deposition EPMA electron-probe microanalyzer EPOx electrochemical partial oxidation GDC Gd2O3-doped CeO2 GGA generalized gradient approximation GNP glycine nitrate process H2-TPR hydrogen-temperature programmed reduction IEA International Energy Agency

Ran Ran is an associate professor at Nanjing University of Technology. She received her Ph.D. from Dalian Institute of Chemical Physics, AP


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(7) Shindell, D. T.; Faluvegi, G.; Koch, D. M.; Schmidt, G. A.; Unger, N.; Bauer, S. E. Science 2009, 326, 716. (8) Granovskii, M.; Dincer, I.; Rosen, M. A. J. Power Sources 2007, 165, 307. (9) Wright, S.; Pinkelman, A. Int. J. Energy Res. 2008, 32, 612. (10) Matsumura, Y.; Nakamori, T. Appl. Catal., A: Gen. 2004, 258, 107. (11) Levent, M.; Gunn, D. J.; El-Bousiffi, M. A. Int. J. Hydrogen Energy 2003, 28, 945. (12) Chen, Y. Z.; Wang, Y. Z.; Xu, H. Y.; Xiong, G. X. Appl. Catal., B: Environ. 2008, 80, 283. (13) Xu, J. H.; Yeung, C. M. Y.; Ni, J.; Meunier, F.; Acerbi, N.; Fowles, M.; Tsang, S. C. Appl. Catal., A: Gen. 2008, 345, 119. (14) Roh, H.-S.; Eum, I.-H.; Jeong, D.-W. Renewable Energy 2012, 42, 212. (15) Steele, B. C. H.; Heinzel, A. Nature 2001, 414, 345. (16) Minh, N. Q. J. Am. Ceram. Soc. 1993, 76, 563. (17) Sangtongkitcharoen, W.; Assabumrungrat, S.; Pavarajarn, V.; Laosiripojana, N.; Praserthdam, P. J. Power Sources 2005, 142, 75. (18) Dicks, A. L. J. Power Sources 1998, 71, 111. (19) Liu, J.; Zhang, S.; Wang, W.; Gao, J.; Liu, W.; Chen, C. J. Power Sources 2012, 217, 287. (20) Adachi, H.; Ahmed, S.; Lee, S. H. D.; Papadias, D.; Ahluwalia, R. K.; Bendert, J. C.; Kanner, S. A.; Yamazaki, Y. J. Power Sources 2009, 188, 244. (21) Halabi, M. H.; de Croon, M. H. J. M.; van der Schaaf, J.; Cobden, P. D.; Schouten, J. C. Appl. Catal., A: Gen. 2010, 389, 68. (22) Wang, Y.-F.; Tsai, C.-H.; Chang, W.-Y.; Kuo, Y.-M. Int. J. Hydrogen Energy 2010, 35, 135. (23) Simsek, E.; Avci, A. K.; Ö nsan, Z. I. Catal. Today 2011, 178, 157. (24) Armor, J. N. Appl. Catal., A: Gen. 1999, 176, 159. (25) Liese, E. A.; Gemmen, R. S. J. Eng. Gas Turbines Power 2005, 127, 86. (26) McLean, G. F.; Niet, T.; Prince-Richard, S.; Djilali, N. Int. J. Hydrogen Energy 2002, 27, 507. (27) Mehta, V.; Cooper, J. S. J. Power Sources 2003, 114, 32. (28) Wasmus, S.; Küver, A. J. Electroanal. Chem. 1999, 461, 14. (29) Watanabe, M.; Tsurumi, K.; Mizukami, T.; Nakamura, T.; Stonehart, P. J. Electrochem. Soc. 1994, 141, 2659. (30) Cavallaro, S.; Mondello, N.; Freni, S. J. Power Sources 2001, 102, 198. (31) Suzuki, T.; Hasan, Z.; Funahashi, Y.; Yamaguchi, T.; Fujishiro, Y.; Awano, M. Science 2009, 325, 852. (32) Oetjen, H.-F.; Schmidt, V. M.; Stimming, U.; Trila, F. J. Electrochem. Soc. 1996, 143, 3838. (33) Jacobson, A. J. Chem. Mater. 2010, 22, 660. (34) Zhang, P.; Huang, Y. H.; Cheng, J. G.; Mao, Z. Q.; Goodenough, J. B. J. Power Sources 2011, 196, 1738. (35) Lai, B.-K.; Kerman, K.; Ramanathan, S. J. Power Sources 2011, 196, 6299. (36) Chen, G.; Guan, G. Q.; Abliz, S.; Kasai, Y.; Abudula, A. Electrochim. Acta 2011, 56, 9868. (37) Liang, B.; Suzuki, T.; Hamamoto, K.; Yamaguchi, T.; Sumi, H.; Fujishiro, Y.; Ingram, B. J.; Carter, J. D. J. Power Sources 2012, 202, 225. (38) Vert, V. B.; Melo, F. V.; Navarrete, L.; Serra, J. M. Appl. Catal., B: Environ. 2012, 115−116, 346. (39) Caillot, T.; Gauthier, G.; Delichère, P.; Cayron, C.; Cadete Santos Aires, F. J. J. Catal. 2012, 290, 158. (40) Murray, E. P.; Harris, S. J.; Jen, H. J. Electrochem. Soc. 2002, 149, A1127. (41) Cimenti, M.; Alzate-Restrepo, V.; Hill, J. M. J. Power Sources 2010, 195, 4002. (42) Van herle, J.; Maréchal, F.; Leuenberger, S.; Membrez, Y.; Bucheli, O.; Favrat, D. J. Power Sources 2004, 131, 127. (43) Fagbenle, R. L.; Oguaka, A. B. C.; Olakoyejo, O. T. Appl. Therm. Eng. 2007, 27, 2220. (44) Jiang, S. P.; Chan, S. H. J. Mater. Sci. 2004, 39, 4405. (45) Tao, S. W.; Irvine, J. T. S. Nat. Mater. 2003, 2, 320.

IT LSC LSCr LSCF LSGM LSGMC5 LSGMC8.5 LSM LSSM LT MDC MEA MP2

intermediate temperature La0.8Sr0.2CoO3 La0.8Sr0.2CrO3 La0.6Sr0.4Fe0.8Co0.2O3 La0.9Sr0.1Ga0.8Mg0.2O3‑δ La0.8Sr0.2Ga0.8Mg0.15Co0.05O3 La0.8Sr0.2Ga0.8Mg0.115Co0.085O3 La0.8Sr0.2MnO3 La0.8Sr0.2Sc0.1Mn0.9O3‑δ low-temperature Mo-doped CeO2 membrane−electrode assembly second-order approximation of Møller−Plesset perturbation theory MVK Mars−van Krevelen NEB nudged elastic band NSDC Na2CO3 and Sm-doped CeO2 Jc critical current density O/C oxygen/carbon molar ratio O2-TPO oxygen-temperature programmed oxidation OCV(s) open circuit voltage(s) ORR oxygen reduction reaction OSC oxygen storage capacity PAW projector augmented wave PBE Perdew−Burke−Ernzerhof PCM (Pr0.7Ca0.3)0.9MnO3 PEMFCs polymer electrolyte membrane fuel cells PPD peak power density PW91 Perdew−Wang 91 Rp polarization resistance S/C steam to carbon molar ratio SC slurry coating SC-SOFC(s) single-chamber SOFC(s) ScSZ scandia-stabilized zirconia SCTx Sm0.2(Ce1−xTix)0.8O1.9 SCYb SrCe0.95Yb0.05O3−δ ScCSZ Zr0.89Sc0.1Ce0.01O2−δ ScYSZ scandia and yttria-stabilized zirconia SDC Sm2O3-doped CeO2 SEM scanning electron microscope SIMS secondary ion mass spectrometry SOFC(s) solid oxide fuel cell(s) SSC Sm0.5Sr0.5CoO3 SZY SrZr0.95Y0.05O3−δ TEC thermal expansion coefficient TPB three-phase boundary VASP Vienna ab initio Simulation Package VGCF vapor grown carbon fiber WGS water−gas shift YST yttria-doped SrTiO3 YSZ yttria-stabilized zirconia

REFERENCES (1) Kaparaju, P.; Buendia, I.; Ellegaard, L.; Angelidakia, I. Bioresour. Technol. 2008, 99, 4919. (2) Ting, C. H.; Lee, D. J. Int. J. Hydrogen Energy 2007, 32, 677. (3) Fountoulakis, M. S.; Manios, T. Bioresour. Technol. 2009, 100, 3043. (4) Sanphoti, N.; Towprayoon, S.; Chaiprasert, P.; Nopharatana, A. J. Environ. Manage. 2006, 81, 27. (5) Yu, K.-P.; Yu, W.-Y.; Kuo, M.-C.; Liou, Y.-C.; Chien, S.-H. Appl. Catal., B: Environ. 2008, 84, 112. (6) Gogate, M. R.; Davis, R. J. Catal. Commun. 2010, 11, 901. AQ


Chemical Reviews

Review

(46) Danilovic, N.; Luo, J. L.; Chuang, K. T.; Sanger, A. R. J. Power Sources 2009, 192, 247. (47) Lee, S.-I.; Ahn, K.; Vohs, J. M.; Gorte, R. J. Electrochem. SolidState Lett. 2005, 8, A48. (48) Gross, M. D.; Vohs, J. M.; Gorte, R. J. Electrochim. Acta 2007, 52, 1951. (49) McIntosh, S.; Gorte, R. J. Chem. Rev. 2004, 104, 4845. (50) Tao, S. W.; Irvine, J. T. S. Chem. Rec. 2004, 4, 83. (51) Sun, C. W.; Stimming, U. J. Power Sources 2007, 171, 247. (52) Goodenough, J. B.; Huang, Y. H. J. Power Sources 2007, 173, 1. (53) Gorte, R. J.; Vohs, J. M. Curr. Opin. Colloid Interface Sci. 2009, 14, 236. (54) Zhu, W. Z.; Deevi, S. C. Mater. Sci. Eng. 2003, A362, 228. (55) Mogensen, M.; Kammer, K. Annu. Rev. Mater. Res. 2003, 33, 321. (56) Steele, B. C. H.; Kelly, I.; Middleton, H.; Rudkin, R. Solid State Ionics 1988, 28−30, 1547. (57) Marina, O. A.; Mogensen, M. Appl. Catal., A: Gen. 1999, 189, 117. (58) Li, J. H.; Fu, X. Z.; Luo, J. L.; Chuang, K. T.; Sanger, A. R. J. Power Sources 2012, 213, 69. (59) Sumi, H.; Ukai, K.; Mizutani, Y.; Mori, H.; Wen, C.-J.; Takahashi, H.; Yamamoto, O. Solid State Ionics 2004, 174, 151. (60) Homel, M.; Gür, T. M.; Koh, J. H.; Virkar, A. V. J. Power Sources 2010, 195, 6367. (61) Gorte, R. J.; Park, S.; Vohs, J. M.; Wang, C. H. Adv. Mater. 2000, 12, 1465. (62) van den Bossche, M.; McIntosh, S. J. Catal. 2008, 255, 313. (63) Raj, E. S.; Kilner, J. A.; Irvine, J. T. S. Solid State Ionics 2006, 177, 1747. (64) Mogensen, M.; Skaarup, S. Solid State Ionics 1996, 86−88, 1151. (65) Jiang, S. P.; Badwal, S. P. S. Solid State Ionics 1999, 123, 209. (66) Brown, M.; Primdahl, S.; Mogensen, M. J. Electrochem. Soc. 2000, 147, 475. (67) Bieberle, A.; Meier, L. P.; Gauckler, L. J. J. Electrochem. Soc. 2001, 148, A646. (68) Ferri, D.; Forni, L. Appl. Catal., B: Environ. 1998, 16, 119. (69) Saracco, G.; Geobaldo, F.; Baldi, G. Appl. Catal., B: Environ. 1999, 20, 277. (70) Sfeir, J.; Buffat, P. A.; Möckli, P.; Xanthopoulos, N.; Vasquez, R.; Mathieu, H. J.; Van herle, J.; Thampi, K. R. J. Catal. 2001, 202, 229. (71) Falcón, H.; Barbero, J. A.; Alonso, J. A.; Martínez-Lope, M. J.; Fierro, J. L. G. Chem. Mater. 2002, 14, 2325. (72) Tao, S. W.; Irvine, J. T. S.; Plint, S. M. J. Phys. Chem. B 2006, 110, 21771. (73) Hu, R. S.; Ding, R. R.; Chen, J.; Hu, J. N.; Zhang, Y. L. Catal. Commun. 2012, 21, 38. (74) Song, K. S.; Klvana, D.; Kirchnerova, J. Appl. Catal., A: Gen. 2001, 213, 113. (75) Belessi, V. C.; Ladavos, A. K.; Armatas, G. S.; Pomonis, P. J. Phys. Chem. Chem. Phys. 2001, 3, 3856. (76) van Santen, R. A.; Neurock, M. Molecular Heterogenous Catalysis; Wiley-VCH: Weinheim, 2006; p 244. (77) McIntosh, S.; van den Bossche, M. Solid State Ionics 2011, 192, 453. (78) Holtappels, P.; Vinke, I. C.; de Haart, L. G. J.; Stimming, U. J. Electrochem. Soc. 1999, 146, 2976. (79) Horita, T.; Kishimoto, H.; Yamaji, K.; Xiong, Y. P.; Sakai, N.; Brito, M. E.; Yokokawa, H. Solid State Ionics 2006, 177, 1941. (80) Holtappels, P.; de Haart, L. G. J.; Stimming, U. J. Electrochem. Soc. 1999, 146, 1620. (81) Jiang, S. P.; Badwal, S. P. S. J. Electrochem. Soc. 1997, 144, 3777. (82) Matsuzaki, Y.; Yasuda, I. J. Electrochem. Soc. 2000, 147, 1630. (83) Aguiar, P.; Adjiman, C. S.; Brandon, N. P. J. Power Sources 2004, 138, 120. (84) Bessler, W. G.; Vogler, M.; Störmer, H.; Gerthsen, D.; Utz, A.; Weber, A.; Ivers-Tiffée, E. Phys. Chem. Chem. Phys. 2010, 12, 13888.

(85) Mizusaki, J.; Tagawa, H.; Saito, T.; Yamamura, T.; Kamitani, K.; Hirano, K.; Ehara, S.; Takagi, T.; Hikita, T.; Ippommatsu, M.; Nakagawa, S.; Hashimoto, K. Solid State Ionics 1994, 70−71, 52. (86) Sukeshini, A. M.; Habibzadeh, B.; Becker, B. P.; Stoltz, C. A.; Eichhorn, B. W.; Jackson, G. S. J. Electrochem. Soc. 2006, 153, A705. (87) Utz, A.; Störmer, H.; Leonide, A.; Weber, A.; Ivers-Tiffée, E. J. Electrochem. Soc. 2010, 157, B920. (88) Jensen, K. V.; Primdahl, S.; Chorkendorff, I.; Mogensen, M. Solid State Ionics 2001, 144, 197. (89) Lauvstad, G. Ø.; Tunold, R.; Sunde, S. J. Electrochem. Soc. 2002, 149, E497. (90) Hansen, K. V.; Norrman, K.; Mogensen, M. J. Electrochem. Soc. 2004, 151, A1436. (91) Schmidt, M. S.; Hansen, K. V.; Norrman, K.; Mogensen, M. Solid State Ionics 2009, 180, 431. (92) Ioroi, T.; Hara, T.; Uchimoto, Y.; Ogumi, Z.; Takehara, Z. J. Electrochem. Soc. 1998, 145, 1999. (93) Horita, T.; Yamaji, K.; Sakai, N.; Yokokawa, H.; Kawada, T.; Kato, T. Solid State Ionics 2000, 127, 55. (94) Adler, S. B. Chem. Rev. 2004, 104, 4791. (95) Mukherjee, J.; Linic, S. J. Electrochem. Soc. 2007, 154, B919. (96) Ingram, D. B.; Linic, S. J. Electrochem. Soc. 2009, 156, B1457. (97) Perdew, J. P.; Wang, Y. Phys. Rev. B 1992, 45, 13244. (98) Anderson, A. B.; Vayner, E. Solid State Ionics 2006, 177, 1355. (99) Shishkin, M.; Ziegler, T. J. Phys. Chem. C 2009, 113, 21667. (100) Perdew, J. P.; Burke, K.; Ernzerhof, M. Phys. Rev. Lett. 1996, 77, 3865. (101) Orio, M.; Pantazis, D. A.; Neese, F. Photosynth. Res. 2009, 102, 443. (102) Bessler, W. G. Solid State Ionics 2005, 176, 997. (103) Goodwin, D. G.; Zhu, H. Y.; Colclasure, A. M.; Kee, R. J. J. Electrochem. Soc. 2009, 156, B1004. (104) Vogler, M.; Bessler, W. G. ECS Trans. 2009, 25, 1957. (105) Yurkiv, V.; Utz, A.; Weber, A.; Ivers-Tiffée, E.; Volpp, H.-R.; Bessler, W. G. Electrochim. Acta 2012, 59, 573. (106) Utz, A.; Leonide, A.; Weber, A.; Ivers-Tiffée, E. J. Power Sources 2011, 196, 7217. (107) Rao, M. V.; Fleig, J.; Zinkevich, M.; Aldinger, F. Solid State Ionics 2010, 181, 1170. (108) Yamazaki, O.; Tomishige, K.; Fujimoto, K. Appl. Catal., A: Gen. 1996, 136, 49. (109) Tomishige, K.; Chen, Y. G.; Fujimoto, K. J. Catal. 1999, 181, 91. (110) Luo, J. Z.; Yu, Z. L.; Ng, C. F.; Au, C. T. J. Catal. 2000, 194, 198. (111) Xu, S.; Zhao, R.; Wang, X. L. Fuel Process. Technol. 2004, 86, 123. (112) Shang, R. J.; Sun, W. Z.; Wang, Y. Y.; Jin, G. Q.; Guo, X. Y. Catal. Commun. 2008, 9, 2103. (113) Trimm, D. L. Catal. Today 1999, 49, 3. (114) Mccarty, J. G.; Wise, H. J. Catal. 1979, 57, 406. (115) Bartholomew, C. H. Catal. Rev.: Sci. Eng. 1982, 24, 67. (116) Guo, J. J.; Lou, H.; Zheng, X. M. Carbon 2007, 45, 1314. (117) Rostrup-Nielsen, J.; Trimm, D. L. J. Catal. 1977, 48, 155. (118) Kock, A. J. H. M.; de Bokx, P. K.; Boellaard, E.; Klop, W.; Geus, J. W. J. Catal. 1985, 96, 468. (119) Holstein, W. L. J. Catal. 1995, 152, 42. (120) Sehested, J. Catal. Today 2006, 111, 103. (121) Liu, C. J.; Ye, J. Y.; Jiang, J. J.; Pan, Y. X. ChemCatChem 2011, 3, 529. (122) Asai, K.; Nagayasu, Y.; Takane, K.; Iwamoto, S.; Yagasaki, E.; Ishii, K.; Inoue, M. J. Jpn. Petrol. Inst. 2008, 51, 42. (123) Inoue, M.; Asai, K.; Nagayasu, Y.; Takane, K.; Iwamoto, S.; Yagasaki, E.; Ishii, K. Diamond Relat. Mater. 2008, 17, 1471. (124) Sasaki, K.; Teraoka, Y. J. Electrochem. Soc. 2003, 150, A878. (125) Sasaki, K.; Teraoka, Y. J. Electrochem. Soc. 2003, 150, A885. (126) Chen, D.; Christensen, K. O.; Ochoa-Fernández, E.; Yu, Z. X.; Tøtdal, B.; Latorre, N.; Monzón, A.; Holmen, A. J. Catal. 2005, 229, 82. AR


Chemical Reviews

Review

(162) Luna, A. E. C.; Iriarte, M. E. Appl. Catal., A: Gen. 2008, 343, 10. (163) Sutthiumporn, K.; Kawi, S. Int. J. Hydrogen Energy 2011, 36, 14435. (164) Wang, S. B.; Lu, G. Q. Appl. Catal., B: Environ. 1998, 19, 267. (165) Guo, J. Z.; Hou, Z. Y.; Gao, J.; Zheng, X. M. Fuel 2008, 87, 1348. (166) Yang, R. Q.; Xing, C.; Lv, C. X.; Shi, L.; Tsubaki, N. Appl. Catal., A: Gen. 2010, 385, 92. (167) Wang, W.; Su, C.; Ran, R.; Shao, Z. P. J. Power Sources 2011, 196, 3855. (168) Amin, M. H.; Mantri, K.; Newnham, J.; Tardio, J.; Bhargava, S. K. Appl. Catal., B: Environ. 2012, 119−120, 217. (169) Pan, Y. X.; Kuai, P. Y.; Liu, Y.; Ge, Q. F.; Liu, C. J. Energy Environ. Sci. 2010, 3, 1322. (170) Pan, Y. X.; Liu, C. J.; Mei, D. H.; Ge, Q. F. Langmuir 2010, 26, 5551. (171) Watwe, R. M.; Bengaard, H. S.; Rostrup-Nielsen, J. R.; Dumesic, J. A.; Nørskov, J. K. J. Catal. 2000, 189, 16. (172) Bengaard, H. S.; Nørskov, J. K.; Sehested, J.; Clausen, B. S.; Nielsen, L. P.; Molenbroek, A. M.; Rostrup-Nielsen, J. R. J. Catal. 2002, 209, 365. (173) Laosiripojana, N.; Assabumrungrat, S.; Charojrochkul, S. Appl. Catal., A: Gen. 2007, 327, 180. (174) Jacobs, G.; Keogh, R. A.; Davis, B. H. J. Catal. 2007, 245, 326. (175) Laosiripojana, N.; Kiatkittipong, W.; Charojrochkul, S.; Assabumrungrat, S. Appl. Catal., A: Gen. 2010, 383, 50. (176) Cheng, C. K.; Foo, S. Y.; Adesina, A. A. Int. J. Hydrogen Energy 2012, 37, 10101. (177) Nikooyeh, K.; Clemmer, R.; Alzate-Restrepo, V.; Hill, J. M. Appl. Catal., A: Gen. 2008, 347, 106. (178) Eguchi, K.; Kojo, H.; Takeguchi, T.; Kikuchi, R.; Sasaki, K. Solid State Ionics 2002, 152−153, 411. (179) Finnerty, C. M.; Coe, N. J.; Cunningham, R. H.; Ormerod, R. M. Catal. Today 1998, 46, 137. (180) Koh, J.-H.; Yoo, Y.-S.; Park, J.-W.; Lim, H. C. Solid State Ionics 2002, 149, 157. (181) Ke, K.; Gunji, A.; Mori, H.; Tsuchida, S.; Takahashi, H.; Ukai, K.; Mizutani, Y.; Sumi, H.; Yokoyama, M.; Waki, K. Solid State Ionics 2006, 177, 541. (182) He, H. P.; Hill, J. M. Appl. Catal., A: Gen. 2007, 317, 284. (183) Sumi, H.; Puengjinda, P.; Muroyama, H.; Matsui, T.; Eguchi, K. J. Power Sources 2011, 196, 6048. (184) Sumi, H.; Lee, Y.-H.; Muroyama, H.; Matsui, T.; Eguchi, K. J. Electrochem. Soc. 2010, 157, B1118. (185) Gunji, A.; Wen, C.; Otomo, J.; Kobayashi, T.; Ukai, K.; Mizutani, Y.; Takahashi, H. J. Power Sources 2004, 131, 285. (186) Asamoto, M.; Miyake, S.; Itagaki, Y.; Sadaoka, Y.; Yahiro, H. Catal. Today 2008, 139, 77. (187) Chen, X. J.; Khor, K. A.; Chan, S. H. Electrochem. Solid-State Lett. 2005, 8, A79. (188) Timmermann, H.; Fouquet, D.; Weber, A.; Ivers-Tiffée, E.; Hennings, U.; Reimert, R. Fuel Cells 2006, 6, 307. (189) Laosiripojana, N.; Assabumrungrat, S. J. Power Sources 2007, 163, 943. (190) Offer, G. J.; Mermelstein, J.; Brightman, E.; Brandon, N. P. J. Am. Ceram. Soc. 2009, 92, 763. (191) Trovarelli, A. Cat. Rev.: Sci. Eng. 1996, 38, 439. (192) King, D. L.; Strohm, J. J.; Wang, X. Q.; Roh, H.-S.; Wang, C. M.; Chin, Y.-H; Wang, Y.; Lin, Y. B.; Rozmiarek, R.; Singh, P. J. Catal. 2008, 258, 356. (193) Prasad, D. H.; Park, S. Y.; Ji, H.; Kim, H.-R.; Son, J.-W.; Kim, B.-K.; Lee, H.-W.; Lee, J.-H. Appl. Catal., A: Gen. 2012, 411−412, 160. (194) Wang, Z. H.; Lü, Z.; Wei, B.; Chen, K. F.; Huang, X. Q.; Pan, W. P.; Su, W. H. Electrochim. Acta 2011, 56, 6688. (195) Pillai, M.; Lin, Y. B.; Zhu, H. Y.; Kee, R. J.; Barnett, S. A. J. Power Sources 2010, 195, 271. (196) Goula, G.; Kiousis, V.; Nalbandian, L.; Yentekakis, I. V. Solid State Ionics 2006, 177, 2119.

(127) Jones, G.; Jakobsen, J. G.; Shim, S. S.; Kleis, J.; Andersson, M. P.; Rossmeisl, J.; Abild-Pedersen, F.; Bligaard, T.; Helveg, S.; Hinnemann, B.; Rostrup-Nielsen, J. R.; Chorkendorff, I.; Sehested, J.; Nørskov, J. K. J. Catal. 2008, 259, 147. (128) Pan, Y. X; Liu, C. J.; Wiltowski, T. S.; Ge, Q. F. Catal. Today 2009, 147, 68. (129) García-Diéguez, M.; Pieta, I. S.; Herrera, M. C.; Larrubia, M. A.; Alemany, L. J. J. Catal. 2010, 270, 136. (130) Ermakova, M. A.; Ermakov, D. Y.; Kuvshinov, G. G.; Plyasova, L. M. J. Catal. 1999, 187, 77. (131) Ermakova, M. A.; Ermakov, D. Y.; Kuvshinov, G. G. Appl. Catal., A: Gen. 2000, 201, 61. (132) Christensen, K. O.; Chen, D.; Lødeng, R.; Holmen, A. Appl. Catal., A: Gen. 2006, 314, 9. (133) Chen, Y. Z.; Zhou, W.; Shao, Z. P.; Xu, N. P. Catal. Commun. 2008, 9, 1418. (134) Chen, L.; Zhu, Q. S.; Hao, Z. G.; Zhang, T.; Xie, Z. H. Int. J. Hydrogen Energy 2010, 35, 8494. (135) Daza, C. E.; Kiennemann, A.; Moreno, S.; Molina, R. Appl. Catal., A: Gen. 2009, 364, 65. (136) Zhu, X. L.; Huo, P. P.; Zhang, Y. P.; Cheng, D. G.; Liu, C. J. Appl. Catal., B: Environ. 2008, 81, 132. (137) Kim, J.-H.; Suh, D. J.; Park, T.-J.; Kim, K.-L. Appl. Catal., A: Gen. 2000, 197, 191. (138) Hao, Z. G.; Zhu, Q. S.; Jiang, Z.; Hou, B. L.; Li, H. Z. Fuel Process. Technol. 2009, 90, 113. (139) Ribeiro, N. F. P.; Neto, R. C. R.; Moya, S. F.; Souza, M. M. V. M.; Schmal, M. Int. J. Hydrogen Energy 2010, 35, 11725. (140) García-Diéguez, M.; Pieta, I. S.; Herrera, M. C.; Larrubia, M. A.; Alemany, L. J. Appl. Catal., A: Gen. 2010, 377, 191. (141) Enger, B. C.; Lødeng, R.; Walmsley, J.; Holmen, A. Appl. Catal., A: Gen. 2010, 383, 119. (142) Wang, W.; Su, C.; Wu, Y. Z.; Ran, R.; Shao, Z. P. J. Power Sources 2010, 195, 402. (143) Dong, W.-S.; Roh, H.-S.; Jun, K.-W.; Park, S.-E.; Oh, Y.-S. Appl. Catal., A: Gen. 2002, 226, 63. (144) Huang, T.-J.; Huang, M.-C. Chem. Eng. J. 2008, 145, 149. (145) Wang, W.; Ran, R.; Su, C.; Shao, Z. P.; Jung, D. W.; Seo, S.; Lee, S. M. Int. J. Hydrogen Energy 2011, 36, 10958. (146) Kobayashi, Y.; Horiguchi, J.; Kobayashi, S.; Yamazaki, Y.; Omata, K.; Nagao, D.; Konno, M.; Yamada, M. Appl. Catal., A: Gen. 2011, 395, 129. (147) Qiao, B. T.; Wang, A. Q.; Yang, X. F.; Allard, L. F; Jiang, Z.; Cui, Y. T.; Liu, J. Y.; Li, J.; Zhang, T. Nat. Chem. 2011, 3, 634. (148) Rostrup-Nielsen, J. R.; Sehested, J.; Nørskov, J. K. Adv. Catal. 2002, 47, 65. (149) Kim, Y. T.; Um, J. H.; Kim, S. H.; Lim, T.-H.; Lee, H.-I. Appl. Catal., A: Gen. 2010, 384, 10. (150) Navarro, R. M.; Peña, M. A.; Fierro, J. L. G. Chem. Rev. 2007, 107, 3952. (151) Nikolla, E.; Schwank, J.; Linic, S. J. Catal. 2007, 250, 85. (152) Nikolla, E.; Schwank, J.; Linic, S. Catal. Today 2008, 136, 243. (153) Guo, J. J.; Lou, H.; Mo, L. Y.; Zheng, X. M. J. Mol. Catal. A: Chem. 2010, 316, 1. (154) Lemonidou, A. A.; Vasalos, I. A. Appl. Catal., A: Gen. 2002, 228, 227. (155) Hou, Z. Y.; Yokota, O.; Tanaka, T.; Yashima, T. Appl. Catal., A: Gen. 2003, 253, 381. (156) Xu, L. L.; Song, H. L.; Chou, L. J. Appl. Catal., B: Environ. 2011, 108−109, 177. (157) Kang, K.-M.; Kim, H.-W; Shim, I.-W.; Kwak, H.-Y. Fuel Process. Technol. 2011, 92, 1236. (158) Verykios, X. E. Int. J. Hydrogen Energy 2003, 28, 1045. (159) Horiuchi, T.; Sakuma, K.; Fukui, T.; Kubo, Y.; Osaki, T.; Mori, T. Appl. Catal., A: Gen. 1996, 144, 111. (160) Nandini, A.; Pant, K. K.; Dhingra, S. C. Appl. Catal., A: Gen. 2005, 290, 166. (161) Solov’ev, S. A.; Zatelepa, R. N.; Gubaren, E. V.; Strizhak, P. E.; Moroz, E. M. Russ. J. Appl. Chem. 2007, 80, 1883. AS


Chemical Reviews

Review

(197) Park, S.; Vohs, J. M.; Gorte, R. J. Nature 2000, 404, 265. (198) Alzate-Restrepo, V.; Hill, J. M. Appl. Catal., A: Gen. 2008, 342, 49. (199) Lin, Y. B.; Zhan, Z. L.; Liu, J.; Barnett, S. A. Solid State Ionics 2005, 176, 1827. (200) Liu, J.; Barnett, S. A. Solid State Ionics 2003, 158, 11. (201) Horita, T.; Yamaji, K.; Kato, T.; Kishimoto, H.; Xiong, Y. P.; Sakai, N.; Brito, M. E.; Yokokawa, H. J. Power Sources 2005, 145, 133. (202) Rostrup-Nielsen, J. R.; Alstrup, I. Catal. Today 1999, 53, 311. (203) Gavrielatos, I.; Drakopoulos, V.; Neophytides, S. G. J. Catal. 2008, 259, 75. (204) Niakolas, D. K.; Ouweltjes, J. P.; Rietveld, G.; Drakopoulos, V.; Neophytides, S. G. Int. J. Hydrogen Energy 2010, 35, 7898. (205) Besenbacher, F.; Chorkendorff, I.; Clausen, B. S.; Hammer, B.; Molenbroek, A. M.; Nørskov, J. K.; Stensgaard, I. Science 1998, 279, 1913. (206) Triantafyllopoulos, N. C.; Neophytides, S. G. J. Catal. 2006, 239, 187. (207) Wang, Z. C.; Weng, W. J.; Cheng, K.; Du, P. Y.; Shen, G.; Han, G. R. J. Power Sources 2008, 179, 541. (208) Park, E. W.; Moon, H.; Park, M.-S.; Hyun, S. H. Int. J. Hydrogen Energy 2009, 34, 5537. (209) Jung, S.; Gross, M. D.; Gorte, R. J.; Vohs, J. M. J. Electrochem. Soc. 2006, 153, A1539. (210) Islam, S.; Hill, J. M. J. Power Sources 2011, 196, 5091. (211) Maček, J.; Novosel, B.; Marinšek, M. J. Eur. Ceram. Soc. 2007, 27, 487. (212) Takeguchi, T.; Kikuchi, R.; Yano, T.; Eguchi, K.; Murata, K. Catal. Today 2003, 84, 217. (213) Hibino, T.; Hashimoto, A.; Yano, M.; Suzuki, M.; Sano, M. Electrochim. Acta 2003, 48, 2531. (214) Strohm, J. J.; Zheng, J.; Song, C. S. J. Catal. 2006, 238, 309. (215) An, W.; Gatewood, D.; Dunlap, B.; Turner, C. H. J. Power Sources 2011, 196, 4724. (216) Jia, L. C.; Wang, X.; Hua, B.; Li, W. L.; Chi, B.; Pu, J.; Yuan, S. L.; Li, J. Int. J. Hydrogen Energy 2012, 37, 11941. (217) Ringuedé, A.; Fagg, D. P.; Frade, J. R. J. Eur. Ceram. Soc. 2004, 24, 1355. (218) Li, W. Y.; Lü, Z.; Zhu, X. B.; Guan, B.; Wei, B.; Guan, C. Z.; Su, W. H. Electrochim. Acta 2011, 56, 2230. (219) Jung, S.-W.; Vohs, J. M.; Gorte, R. J. J. Electrochem. Soc. 2007, 154, B1270. (220) Horikoshi, S.; Matsubara, A.; Takayama, S.; Sato, M.; Sakai, F.; Kajitani, M.; Abe, M.; Serpone, N. Appl. Catal., B: Environ. 2010, 99, 490. (221) Galletti, A. M. R.; Antonetti, C.; Venezia, A. M.; Giambastiani, G. Appl. Catal., A: Gen. 2010, 386, 124. (222) Yoshida, K.; Gonzalez-Arellano, C.; Luque, R.; Gai, P. L. Appl. Catal., A: Gen. 2010, 379, 38. (223) Zhang, C. M.; Lin, Y.; Ran, R.; Shao, Z. P. Int. J. Hydrogen Energy 2010, 35, 8171. (224) McIntosh, S.; Vohs, J. M.; Gorte, R. J. J. Electrochem. Soc. 2003, 150, A470. (225) Jung, S.; Lu, C.; He, H. P.; Ahn, K.; Gorte, R. J.; Vohs, J. M. J. Power Sources 2006, 154, 42. (226) Kim, H.; Lu, C.; Worrell, W. L.; Vohs, J. M.; Gorte, R. J. J. Electrochem. Soc. 2002, 149, A247. (227) Sin, A.; Kopnin, E.; Dubitsky, Y.; Zaopo, A.; Aricò, A. S.; Rosa, D. L.; Gullo, L. R.; Antonucci, V. J. Power Sources 2007, 164, 300. (228) Hornés, A.; Gamarra, D.; Munuera, G.; Conesa, J. C.; Martínez-Arias, A. J. Power Sources 2007, 169, 9. (229) Xie, Z.; Xia, C. R.; Zhang, M. Y.; Zhu, W.; Wang, H. T. J. Power Sources 2006, 161, 1056. (230) Lee, S.-I.; Vohs, J. M.; Gorte, R. J. J. Electrochem. Soc. 2004, 151, A1319. (231) Kan, H.; Lee, H. Catal. Commun. 2010, 12, 36. (232) Wang, S. Z.; Gao, J. Electrochem. Solid-State Lett. 2006, 9, A395. (233) Lu, X. C.; Zhu, J. H. J. Power Sources 2007, 165, 678.

(234) Ishihara, T.; Yan, J. W.; Shinagawa, M.; Matsumoto, H. Electrochim. Acta 2006, 52, 1645. (235) Zhu, H. Y.; Wang, W.; Ran, R.; Su, C.; Shi, H. G.; Shao, Z. P. Int. J. Hydrogen Energy 2012, 37, 9801. (236) Wu, Y. Z.; Su, C.; Wang, W.; Wang, H. T.; Shao, Z. P. Int. J. Hydrogen Energy 2012, 37, 9287. (237) Nabae, Y.; Yamanaka, I.; Hatano, M.; Otsuka, K. J. Electrochem. Soc. 2006, 153, A140. (238) Nabae, Y.; Yamanaka, I.; Hatano, M.; Otsuka, K. J. Phys. Chem. C 2008, 112, 10308. (239) Nabae, Y.; Yamanaka, I. Appl. Catal., A: Gen. 2009, 369, 119. (240) Nikolla, E.; Holewinski, A.; Schwank, J.; Linic, S. J. Am. Chem. Soc. 2006, 128, 11354. (241) Nikolla, E.; Schwank, J.; Linic, S. J. Am. Chem. Soc. 2009, 131, 2747. (242) Nikolla, E.; Schwank, J.; Linic, S. J. Catal. 2009, 263, 220. (243) Kan, H.; Lee, H. Appl. Catal., B: Environ. 2010, 97, 108. (244) Nikolla, E.; Schwank, J.; Linic, S. J. Electrochem. Soc. 2009, 156, B1312. (245) Singh, A.; Hill, J. M. J. Power Sources 2012, 214, 185. (246) Zhan, Z. L.; Barnett, S. A. Science 2005, 308, 844. (247) Sun, C. W.; Xie, Z.; Xia, C. R.; Li, H.; Chen, L. Q. Electrochem. Commun. 2006, 8, 833. (248) Ye, X. F.; Wang, S. R.; Wang, Z. R.; Xiong, L.; Sun, X. F.; Wen, T. L. J. Power Sources 2008, 177, 419. (249) Liao, M. M.; Wang, W.; Ran, R.; Shao, Z. P. J. Power Sources 2011, 196, 6177. (250) Jin, C.; Yang, C. H.; Zheng, H. H.; Chen, F. L. J. Power Sources 2012, 201, 66. (251) Liu, X. J.; Zhan, Z. L.; Meng, X.; Huang, W. H.; Wang, S. R.; Wen, T. L. J. Power Sources 2012, 199, 138. (252) Wang, W.; Zhou, W.; Ran, R.; Cai, R.; Shao, Z. P. Electrochem. Commun. 2009, 11, 194. (253) Zhu, H. Y.; Colclasure, A. M.; Kee, R. J.; Lin, Y. B.; Barnett, S. A. J. Power Sources 2006, 161, 413. (254) Lin, Y. B.; Zhan, Z. L.; Barnett, S. A. J. Power Sources 2006, 158, 1313. (255) Zhan, Z. L.; Barnett, S. A. Solid State Ionics 2005, 176, 871. (256) Wang, K.; Ran, R.; Shao, Z. P. J. Power Sources 2007, 170, 251. (257) Wang, W.; Ran, R.; Shao, Z. P. Int. J. Hydrogen Energy 2011, 36, 755. (258) Prasad, D. H.; Kim, H.-R.; Son, J.-W.; Kim, B.-K.; Lee, H.-W.; Lee, J.-H. Catal. Commun. 2009, 10, 1334. (259) Wang, W.; Ran, R.; Shao, Z. P. J. Power Sources 2011, 196, 90. (260) Wang, W.; Su, C.; Ran, R.; Park, H. J.; Kwak, C.; Shao, Z. P. Int. J. Hydrogen Energy 2011, 36, 5632. (261) Jin, C.; Yang, C. H.; Zhao, F.; Coffin, A.; Chen, F. L. Electrochem. Commun. 2010, 12, 1450. (262) Suzuki, T.; Yamaguchi, T.; Hamamoto, K.; Fujishiro, Y.; Awano, M.; Sammes, N. Energy Environ. Sci 2011, 4, 940. (263) He, B. B.; Ding, D.; Xia, C. R. J. Power Sources 2010, 195, 1359. (264) He, B. B.; Wang, W. D.; Zhao, L.; Xia, C. R. Electrochim. Acta 2011, 56, 7071. (265) He, B. B.; Zhao, L.; Wang, W. D.; Chen, F. L.; Xia, C. R. Electrochem. Commun. 2011, 13, 194. (266) He, B. B.; Zhao, L.; Song, S. X.; Jiang, Z. Y.; Xia, C. R. Int. J. Hydrogen Energy 2011, 36, 5589. (267) Zhou, X. L.; Zhen, J. M.; Liu, L. M.; Li, X. K.; Zhang, N. Q.; Sun, K. N. J. Power Sources 2012, 201, 128. (268) Nakagawa, N.; Sagara, H.; Kato, K. J. Power Sources 2001, 92, 88. (269) Takeguchi, T.; Kani, Y.; Yano, T.; Kikuchi, R.; Eguchi, K.; Tsujimoto, K.; Uchida, Y.; Ueno, A.; Omoshiki, K.; Aizawa, M. J. Power Sources 2002, 112, 588. (270) Qiao, J. S.; Sun, K. N.; Zhang, N. Q.; Sun, B.; Kong, J. R.; Zhou, D. R. J. Power Sources 2007, 169, 253. (271) Qiao, J.; Zhang, N.; Wang, Z.; Mao, Y.; Sun, K.; Yuan, Y. Fuel Cells 2009, 9, 729. (272) Takeguchi, T.; Furukawa, S.; Inoue, M. J. Catal. 2001, 202, 14. AT


Chemical Reviews

Review

(309) Alfonso, D. R. Surf. Sci. 2008, 602, 2758. (310) Galea, N. M.; Kadantsev, E. S.; Ziegler, T. J. Phys. Chem. C 2007, 111, 14457. (311) Song, C. S.; Ma, X. L. Appl. Catal., B: Environ. 2003, 41, 207. (312) Farag, H.; Mochida, I. J. Colloid Interface Sci. 2012, 372, 121. (313) Sano, Y.; Choi, K.-H.; Korai, Y.; Mochida, I. Appl. Catal., B: Environ. 2004, 49, 219. (314) Cui, H.; Turn, S. Q.; Reese, M. A. Catal. Today 2009, 139, 274. (315) Breysse, M.; Djega-Mariadassou, G.; Pessayre, S.; Geantet, C.; Vrinat, M.; Pérot, G.; Lemaire, M. Catal. Today 2003, 84, 129. (316) Bhandari, V. M.; Ko, C. H.; Park, J. G.; Han, S.-S.; Cho, S.-H.; Kim, J.-N. Chem. Eng. Sci. 2006, 61, 2599. (317) Flytzani-Stephanopoulos, M.; Sakbodin, M.; Wang, Z. Science 2006, 312, 1508. (318) Cheah, S.; Parent, Y. O.; Jablonski, W. S.; Vinzant, T.; Olstad, J. L. Fuel 2012, 97, 612. (319) Brightman, E.; Ivey, D. G.; Brett, D. J. L.; Brandon, N. P. J. Power Sources 2011, 196, 7182. (320) Lakshminarayanan, N.; Ozkan, U. S. Appl. Catal., A: Gen. 2011, 393, 138. (321) Wang, J.-H.; Liu, M. L. J. Power Sources 2008, 176, 23. (322) Zha, S. W.; Cheng, Z.; Liu, M. L. J. Electrochem. Soc. 2007, 154, B201. (323) Hagen, A.; Rasmussen, J. F. B.; Thydén, K. J. Power Sources 2011, 196, 7271. (324) Grgicak, C. M.; Green, R. G.; Giorgi, J. B. J. Power Sources 2008, 179, 317. (325) Grgicak, C. M.; Pakulska, M. M.; O’Brien, J. S.; Giorgi, J. B. J. Power Sources 2008, 183, 26. (326) Zheng, L. L.; Wang, X.; Zhang, L.; Wang, J. Y.; Jiang, S. P. Int. J. Hydrogen Energy 2012, 37, 10299. (327) Devianto, H.; Yoon, S. P.; Nam, S. W.; Han, J.; Lim, T.-H. J. Power Sources 2006, 159, 1147. (328) Xu, C. C.; Gansor, P.; Zondlo, J. W.; Sabolsky, K.; Sabolsky, E. M. J. Electrochem. Soc. 2011, 158, B1405. (329) Chen, H.-T.; Choi, Y. M.; Liu, M. L.; Lin, M. C. J. Phys. Chem. C 2007, 111, 11117. (330) Kurokawa, H.; Sholklapper, T. Z.; Jacobson, C. P.; De Jonghe, L. C.; Visco, S. J. Electrochem. Solid-State Lett. 2007, 10, B135. (331) Yun, J. W.; Yoon, S. P.; Han, J.; Park, S.; Kim, H. S.; Nam, S. W. J. Electrochem. Soc. 2010, 157, B1825. (332) Yun, J. W.; Yoon, S. P.; Park, S.; Kim, H. S.; Nam, S. W. Int. J. Hydrogen Energy 2011, 36, 787. (333) Chen, Y.; Bunch, J.; Jin, C.; Yang, C. H.; Chen, F. L. J. Power Sources 2012, 204, 40. (334) Ferrizz, R. M.; Gorte, R. J.; Vohs, J. M. Appl. Catal., B: Environ. 2003, 43, 273. (335) Karjalainen, H.; Lassi, U.; Rahkamaa-Tolonen, K.; Kröger, V.; Keiski, R. L. Catal. Today 2005, 100, 291. (336) Choi, S.; Wang, J.; Cheng, Z.; Liu, M. L. J. Electrochem. Soc. 2008, 155, B449. (337) Ettler, M.; Timmermann, H.; Malzbender, J.; Weber, A.; Menzler, N. H. J. Power Sources 2010, 195, 5452. (338) Sarantaridis, D.; Atkinson, A. Fuel Cells 2007, 7, 246. (339) Waldbillig, D.; Wood, A.; Ivey, D. G. J. Power Sources 2005, 145, 206. (340) Malzbender, J.; Wessel, E.; Steinbrech, R. W. Solid State Ionics 2005, 176, 2201. (341) Wood, A.; Pastula, M.; Waldbillig, D.; Ivey, D. G. J. Electrochem. Soc. 2006, 153, A1929. (342) Klemensø, T.; Chung, C.; Larsen, P. H.; Mogensen, M. J. Electrochem. Soc. 2005, 152, A2186. (343) Waldbillig, D.; Wood, A.; Ivey, D. G. Solid State Ionics 2005, 176, 847. (344) Klemensø, T.; Appel, C. C.; Mogensen, M. Electrochem. SolidState Lett 2006, 9, A403. (345) Tikekar, N. M.; Armstrong, T. J.; Virkar, A. V. J. Electrochem. Soc. 2006, 153, A654.

(273) Roh, H.-S.; Jun, K.-W.; Dong, W.-S.; Chang, J.-S.; Park, S.-E.; Joe, Y.-I. J. Mol. Catal. A: Chem. 2002, 181, 137. (274) Zhu, W.; Xia, C. R.; Fan, J.; Peng, R. R.; Meng, G. Y. J. Power Sources 2006, 160, 897. (275) Zhang, L. S.; Gao, J. F.; Liu, M. F.; Xia, C. R. J. Alloys Compd. 2009, 482, 168. (276) Liu, Z. B.; Ding, D.; Liu, B. B.; Guo, W. W.; Wang, W. D.; Xia, C. R. J. Power Sources 2011, 196, 8561. (277) Chen, Y.; Chen, F. L.; Wang, W. D.; Ding, D.; Gao, J. F. J. Power Sources 2011, 196, 4987. (278) Jin, Y.; Saito, H.; Yamahara, K.; Ihara, M. Electrochem. SolidState Lett 2009, 12, B8. (279) Jin, Y.; Yasutake, H.; Yamahara, K.; Ihara, M. J. Electrochem. Soc. 2010, 157, B130. (280) Jin, Y.; Yasutake, H.; Yamahara, K.; Ihara, M. Chem. Eng. Sci. 2010, 65, 597. (281) Shimada, H.; Takami, E.; Ohba, F.; Takei, C.; Hagiwara, A.; Ihara, M. J. Electrochem. Soc. 2011, 158, B1341. (282) Yan, A. Y.; Phongaksorn, M.; Nativel, D.; Croiset, E. J. Power Sources 2012, 210, 374. (283) Tu, B. F.; Dong, Y. L.; Liu, B.; Cheng, M. J. J. Power Sources 2007, 165, 120. (284) Shiratori, Y.; Teraoka, Y.; Sasaki, K. Solid State Ionics 2006, 177, 1371. (285) Lay, E.; Metcalfe, C.; Kesler, O. J. Power Sources 2012, 218, 237. (286) Liu, Y.; Bai, Y. H.; Liu, J. J. Power Sources 2011, 196, 9965. (287) Sakai, T.; Zhong, H.; Eto, H.; Ishihara, T. J. Solid State Electrochem. 2010, 14, 1777. (288) Shiratori, Y.; Sasaki, K. J. Power Sources 2008, 180, 738. (289) Asamoto, M.; Miyake, S.; Sugihara, K.; Yahiro, H. Electrochem. Commun. 2009, 11, 1508. (290) La Rosa, D.; Sin, A.; Lo Faro, M.; Monforte, G.; Antonucci, V.; Aricò, A. S. J. Power Sources 2009, 193, 160. (291) Yang, L.; Choi, Y. M.; Qin, W. T.; Chen, H. Y.; Blinn, K.; Liu, M. F.; Liu, P.; Bai, J. M.; Tyson, T. A.; Liu, M. L. Nat. Commun. 2011, 2, 1. (292) Myung, J.; Lee, J.-J.; Hyun, S.-H. Electrochem. Solid-State Lett. 2010, 13, B43. (293) Tavares, A. C.; Kuzin, B. L.; Beresnev, S. M.; Bogdanovich, N. M.; Kurumchin, E. K.; Dubitsky, Y. A.; Zaopo, A. J. Power Sources 2008, 183, 20. (294) Yoon, S. P.; Han, J.; Nam, S. W.; Lim, T.-H.; Hong, S.-A. J. Power Sources 2004, 136, 30. (295) Yun, J. W.; Yoon, S. P.; Kim, H. S.; Han, J.; Nam, S. W. Int. J. Hydrogen Energy 2012, 37, 4356. (296) Sasaki, K.; Susuki, K.; Iyoshi, A.; Uchimura, M.; Imamura, N.; Kusaba, H.; Teraoka, Y.; Fuchino, H.; Tsujimoto, K.; Uchida, Y.; Jingo, N. J. Electrochem. Soc. 2006, 153, A2023. (297) Matsuzaki, Y.; Yasuda, I. Solid State Ionics 2000, 132, 261. (298) Bartholomew, C. H.; Agrawal, P. K.; Katzer, J. R. Adv. Catal. 1982, 31, 135. (299) Bartholomew, C. H. Appl. Catal., A: Gen. 2001, 212, 17. (300) Smith, T. R.; Wood, A.; Birss, V. I. Appl. Catal., A: Gen. 2009, 354, 1. (301) Rasmussen, J. F. B.; Hagen, A. J. Power Sources 2009, 191, 534. (302) Zhang, L.; Jiang, S. P.; He, H. Q.; Chen, X. B.; Ma, J.; Song, X. C. Int. J. Hydrogen Energy 2010, 35, 12359. (303) Gong, M. Y.; Liu, X. B.; Trembly, J.; Johnson, C. J. Power Sources 2007, 168, 289. (304) Dong, J.; Cheng, Z.; Zha, S. W.; Liu, M. L. J. Power Sources 2006, 156, 461. (305) Sasaki, K.; Haga, K.; Yoshizumi, T.; Minematsu, D.; Yuki, E.; Liu, R. R.; Uryu, C.; Oshima, T.; Ogura, T.; Shiratori, Y.; Ito, K.; Koyama, M.; Yokomoto, K. J. Power Sources 2011, 196, 9130. (306) Sasaki, K.; Adachi, S.; Haga, K.; Uchikawa, M.; Yamamoto, J.; Iyoshi, A.; Chou, J.-T.; Shiratori, Y.; Itoh, K. ECS Trans. 2007, 7, 1675. (307) Hansen, J. B. Electrochem. Solid-State Lett. 2008, 11, B178. (308) Wang, J.-H.; Liu, M. L. Electrochem. Commun. 2007, 9, 2212. AU


Chemical Reviews

Review

(382) Murray, E. P.; Harris, S. J.; Liu, J.; Barnett, S. A. Electrochem. Solid-State Lett. 2005, 8, A531. (383) Laosiripojana, N.; Assabumrungrat, S. Appl. Catal., A: Gen. 2007, 320, 105. (384) Wang, S. Z.; Ishihara, T.; Takita, Y. Appl. Catal., A: Gen. 2002, 228, 167. (385) Zhang, Q. J.; Li, X. H.; Fujimoto, K.; Asami, K. Appl. Catal., A: Gen. 2005, 288, 169. (386) Su, C.; Wang, W.; Ran, R.; Zheng, T.; Shao, Z. P. Int. J. Hydrogen Energy 2012, 37, 6844. (387) Yano, M.; Kawai, T.; Okamoto, K.; Nagao, M.; Sano, M.; Tomita, A.; Hibino, T. J. Electrochem. Soc. 2007, 154, B865. (388) Sato, K.; Tanaka, Y.; Negishi, A.; Kato, T. J. Power Sources 2012, 217, 37. (389) Sahibzada, M.; Steele, B. C. H.; Barth, D.; Rudkin, R. A.; Metcalfe, I. S. Fuel 1999, 78, 639. (390) Sahibzada, M.; Steele, B. C. H.; Hellgardt, K.; Barth, D.; Effendi, A.; Mantzavinos, D.; Metcalfe, I. S. Chem. Eng. Sci. 2000, 55, 3077. (391) Babaei, A.; Zhang, L.; Liu, E.; Jiang, S. P. Int. J. Hydrogen Energy 2012, 37, 15301. (392) Jiang, Y.; Virkar, A. V. J. Electrochem. Soc. 2001, 148, A706. (393) Liu, M. F.; Peng, R. R.; Dong, D. H.; Gao, J. F.; Liu, X. Q.; Meng, G. Y. J. Power Sources 2008, 185, 188. (394) Meng, X.; Zhan, Z. L.; Liu, X. J.; Wu, H.; Wang, S. R.; Wen, T. L. J. Power Sources 2011, 196, 9961. (395) Lo Faro, M.; Stassi, A.; Antonucci, V.; Modafferi, V.; Frontera, P.; Antonucci, P.; Aricò, A. S. Int. J. Hydrogen Energy 2011, 36, 9977. (396) Sun, L. L.; Hao, Y.; Zhang, C. M.; Ran, R.; Shao, Z. P. Int. J. Hydrogen Energy 2010, 35, 7971. (397) Cimenti, M.; Hill, J. M. J. Power Sources 2009, 186, 377. (398) Sasaki, K.; Watanabe, K.; Teraoka, Y. J. Electrochem. Soc. 2004, 151, A965. (399) Gao, Z.; Raza, R.; Zhu, B.; Mao, Z. Q. Int. J. Energy Res. 2011, 35, 690. (400) Assabumrungrat, S.; Laosiripojana, N.; Pavarajarn, V.; Sangtongkitcharoen, W.; Tangjitmatee, A.; Praserthdam, P. J. Power Sources 2005, 139, 55. (401) Saunders, G. J.; Preece, J.; Kendall, K. J. Power Sources 2004, 131, 23. (402) Dokmaingam, P.; Assabumrungrat, S.; Soottitantawat, A.; Laosiripojana, N. J. Power Sources 2010, 195, 69. (403) Laosiripojana, N.; Assabumrungrat, S. Chem. Eng. Sci. 2006, 61, 2540. (404) Leone, P.; Lanzini, A.; Ortigoza-Villalba, G. A.; Borchiellini, R. Chem. Eng. J. 2012, 191, 349. (405) Brett, D. J. L.; Atkinson, A.; Cumming, D.; Ramírez-Cabrera, E.; Rudkin, R.; Brandon, N. P. Chem. Eng. Sci. 2005, 60, 5649. (406) Kim, T.; Ahn, K.; Vohs, J. M.; Gorte, R. J. J. Power Sources 2007, 164, 42. (407) Cimenti, M.; Hill, J. M. J. Power Sources 2010, 195, 54. (408) Li, H. J.; Tian, Y.; Wang, Z. M.; Qie, F. C.; Li, Y. D. RSC Adv. 2012, 2, 3857. (409) Feng, B.; Wang, C. Y.; Zhu, B. Electrochem. Solid-State Lett. 2006, 9, A80. (410) Xu, S. H.; Niu, X. M.; Chen, M. M.; Wang, C. Y.; Zhu, B. J. Power Sources 2007, 165, 82. (411) Raza, R.; Qin, H. Y.; Liu, Q. H.; Samavati, M.; Lima, R. B.; Zhu, B. Adv. Energy Mater. 2011, 1, 1225. (412) Mat, M. D.; Liu, X. R.; Zhu, Z. G.; Zhu, B. Int. J. Hydrogen Energy 2007, 32, 796. (413) Su, C.; Shao, Z. P.; Lin, Y.; Wu, Y. Z.; Wang, H. T. Phys. Chem. Chem. Phys. 2012, 14, 12173.

(346) Modena, S.; Ceschini, S.; Tomasi, A.; Montinaro, D.; Sglavo, V. M. J. Fuel Cell Sci. Technol. 2006, 3, 487. (347) Pihlatie, M.; Kaiser, A.; Larsen, P. H.; Mogensen, M. J. Electrochem. Soc. 2009, 156, B322. (348) Malzbender, J.; Steinbrech, R. W. J. Power Sources 2007, 173, 60. (349) Hatae, T.; Matsuzaki, Y.; Yamashita, S.; Yamazaki, Y. J. Electrochem. Soc. 2009, 156, B609. (350) Heo, Y.-H.; Lee, J.-W.; Lee, S.-B.; Lim, T.-H.; Park, S.-J.; Song, R.-H.; Park, C.-O.; Shin, D.-R. Int. J. Hydrogen Energy 2011, 36, 797. (351) Vedasri, V.; Young, J. L; Birss, V. I. J. Power Sources 2010, 195, 5534. (352) Buyukaksoy, A.; Petrovsky, V.; Dogan, F. J. Electrochem. Soc. 2012, 159, B232. (353) Busawon, A. N.; Sarantaridis, D.; Atkinson, A. Electrochem. Solid-State Lett. 2008, 11, B186. (354) Kim, S.-D.; Moon, H.; Hyun, S.-H.; Moon, J.; Kim, J.; Lee, H.W. Solid State Ionics 2006, 177, 931. (355) Jackson, R. W.; Leonard, J. P.; Pettit, F. S.; Meier, G. H. Solid State Ionics 2008, 179, 2111. (356) Ju, Y.-W.; Ida, S.; Inagaki, T.; Ishihara, T. J. Power Sources 2011, 196, 6062. (357) Pillai, M. R.; Kimb, I.; Bierschenka, D. M.; Barnett, S. A. J. Power Sources 2008, 185, 1086. (358) Ma, Q.; Tietz, F.; Leonide, A.; Ivers-Tiffée, E. Electrochem. Commun. 2010, 12, 1326. (359) Puengjinda, P.; Muroyama, H.; Matsui, T.; Eguchi, K. J. Power Sources 2012, 216, 409. (360) Bermúdez, J. M.; Fidalgo, B.; Arenillas, A.; Menéndez, J. A. Fuel 2012, 94, 197. (361) Jun, K.-W.; Roh, H.-S.; Kim, K.-S.; Ryu, J.-S.; Lee, K.-W. Appl. Catal., A: Gen. 2004, 259, 221. (362) Zhang, Q. W.; Liu, P.; Fujiyama, Y.; Chen, C.; Li, X. H. Appl. Catal., A: Gen. 2011, 401, 147. (363) Vollmar, H.-E.; Maier, C.-U.; Nölscher, C.; Merklein, T.; Poppinger, M. J. Power Sources 2000, 86, 90. (364) Leal, E. M.; Brouwer, J. J. Fuel Cell Sci. Technol. 2006, 3, 137. (365) Zhu, H. Y.; Kee, R. J.; Pillai, M. R.; Barnett, S. A. J. Power Sources 2008, 183, 143. (366) Zhang, X. G.; Ohara, S.; Chen, H.; Fukui, T. Fuel 2002, 81, 989. (367) Zhan, Z. L.; Lin, Y. B.; Pillai, M.; Kim, I.; Barnett, S. A. J. Power Sources 2006, 161, 460. (368) Karl, J.; Frank, N.; Karellas, S.; Saule, M.; Hohenwarter, U. J. Fuel Cell Sci. Technol. 2009, 6, 02115−1. (369) Shao, Z. P.; Zhang, C. M.; Wang, W.; Su, C.; Zhou, W.; Zhu, Z. H.; Park, H. J.; Kwak, C. Angew. Chem., Int. Ed. 2011, 50, 1792. (370) Zhu, H. Y.; Kee, R. J. J. Power Sources 2007, 169, 315. (371) Zhu, H. Y.; Kee, R. J.; Janardhanan, V. M.; Deutschmann, O.; Goodwin, D. G. J. Electrochem. Soc. 2005, 152, A2427. (372) Sobyanin, V. A.; Belyaev, V. D. Solid State Ionics 2000, 136− 137, 747. (373) Pillai, M. R.; Bierschenk, D. M.; Barnett, S. A. Catal. Lett. 2008, 121, 19. (374) Ishihara, T.; Yamada, T.; Akbay, T.; Takita, Y. Chem. Eng. Sci. 1999, 54, 1535. (375) Tagawa, T.; Moe, K. K.; Ito, M.; Goto, S. Chem. Eng. Sci. 1999, 54, 1553. (376) Hibino, T.; Iwahara, H. Chem. Lett. 1993, 1131. (377) Su, C.; Ran, R.; Wang, W.; Shao, Z. P. J. Power Sources 2011, 196, 1967. (378) Wang, S. Z.; Ishihara, T.; Takita, Y. Electrochem. Solid-State Lett. 2002, 5, A177. (379) Su, C.; Wang, W.; Shi, H. G.; Ran, R.; Park, H. J.; Kwak, C.; Shao, Z. P. J. Power Sources 2011, 196, 7601. (380) Liu, Y.; Guo, Y. M.; Wang, W.; Su, C.; Ran, R.; Wang, H. T.; Shao, Z. P. J. Power Sources 2011, 196, 9246. (381) Faungnawakij, K.; Kikuchi, R.; Eguchi, K. J. Power Sources 2007, 164, 73. AV


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