14.2.1 Electronic/ionic conductivity

Electronic conductivity in anode materials is considered as the primary selec- tion criterion, with an absolute minimum of 1 S cm^− 1 and an ideal conductivity superior to 10 S cm^− 1(Atkinson et al., 2004). High conductivity in anode materials is necessary for transfer of electrons. Most efficient anode materials are, therefore, highly conductive semiconductors or metallic conductors.

Ionic conductivity in anode materials is achieved through the use of an ionically conductive material in a cermet or through the use of a mixed ionic–electronic conductor (MIEC). Instead of cermets, an ideal solution is to identify a redox stable material which simultaneously exhibits high ionic and electronic conductivities.

The inherent efficiency of a fuel cell is partially derived from the size of the electrode triple phase boundary (TPB) (OØHayre et al, 2005). The TPB is the specific spatial position at which the hydrogen oxidation reaction (HOR) and the oxygen reduction reaction (ORR) occur, where the electrolyte, gas, and catalytically active regions are electrically connected. Figure 14.1 illustrates the TPB regions of different SOFC anode materials, those with a pure electronic conductor (a) and with an O^2_/e' mixed conductor (b). High ionic conductivity in the anode will extend the TPB thus increasing the number of active sites for the HOR and reducing the electrode polarisation resistance thus improving fuel cell power density.

When Cu-CeO₂-YSZ cermet was used as anode, it was observed that gas-phase pyrolysis reactions on anode can lead to tar formation on copper/ceria cermets; however, the compounds that form on copper tend to be polyaromatics, such as naphthalene and anthracene, rather than graphite. It has been suggested that the polyaromatic compounds enhance anode performance by providing additional electronic conductivity in the anode (McIntosh et al., 2003; Atkinson et al., 2004).

14.2.2 Catalytic activity

In theory, pure hydrogen is the ideal fuel for fuel cells; however, current infrastructure, supply and production of high-purity hydrogen gas is limited (Backhaus-Ricoult, 2008). Research into direct use of alternative fuels has been ongoing and much of the current investigation into anode materials utilises materials with catalytic activity for hydrocarbon oxidation. Previously, precious metal catalysts (Pt, Pd, Au) were used for both fuel cell electrodes but the cost of implementation was an obstacle to commercialisation.

Thus, many of the SOFC anode materials currently under development are designed in such a way to incorporate active catalytic sites within their structures, as catalytically active anode materials or as intercalated metals and metal oxides. For example, the introduction of manganese or iron at the B-site of La₁ _ _xSr_xCrO₃ can dramatically increase the catalytic property of the perovskite oxides and make them viable anode materials for SOFCs (Tao and Irvine, 2003, 2004b; Tao et al, 2004). Good catalytic properties are an essential requirement in order to facilitate the anode reaction and reduce the anode polarisation resistance.

14.2.3 Stability and compatibility

As SOFCs operate at very high temperatures, both the electrolyte and the electrodes have to be stable at operating temperatures and also withstand thermal and redox cycling upon start-up and shut-down. This is vital for use of SOFCs as auxiliary power units (APU) for transportation applications (H. Yokokawa, 2003).

An additional consideration is the thermal expansion coefficient (TEC), which is the specific change in volume per change in temperature at a constant pressure. As some materials may expand/contract more than others upon heating/cooling cycles, it is essential to match the TECs of components which contact to avoid stresses building at the interface, possibly causing delamination and cracking.

Chemical reactivity and stability is also an important aspect of SOFC materials as it governs durability and constitutes the most critical challenge for stationary power units (H. Yokokawa, 2003). Ideally, the material will be chemically compatible with contacting components and fuel impurities. However, materials currently employed are subject to inter-diffusion (Yokokawa et al, 2006) and poisoning (Matsuzaki and Yasuda, 2000).

14.2.4 Microstructure

Requirements for SOFC anode materials are not only on the materials themselves, good microstructure is also very important to achieve excellent performance. For a Ni-YSZ anode, a porosity of about 40–50% is ideal to reduce the anode polarisation and achieve the best performance (Sakamoto et al., 2008). Pore formers such as acrylic and starch are normally mixed with anode materials to form the right microstructure (Haslam et al, 2005). Suzuki et al. (2009) reported that in three cells with anode porosities 57% (Cell A), 47% (Cell B) and 37% (Cell C), Cell A exhibits the best performance. The microstructures of anode with different porosities are shown in Fig. 14.2. Not only porosity, but also cell performance is related to the particle size of NiO in the cermet. NiO particle size in Cell A was < 100 nm while >500 nm in Cells B and C due to the different firing temperature. The produced Ni is also smaller from reduction of smaller NiO leading to higher catalytic activities, smaller anode polarisation and high fuel cell performance (Fig. 14.2).

qAtkinson et al. (Busawon et al, 2008) found that infiltration of nickel in a porous YSZ can effectively enhance the dimensional stability of anode and thus the redox stability.

14.4.1 (La₀.₇₅Sr₀.₂₅)Cr₀.₅Mn₀.₅O₃-based anode materials

Much of the research into novel anode materials has focused on the production of perovskite-type compounds (Marina et al., 2002; Tao and Irvine, 2003; Huang et al, 2006). Perovskite is a naturally occurring structure which has the general formula ABO₃, in which the larger A-site cation occupies a 12-coordinate site and the B-site cation occupies a 6-coordinate site (Tao and Irvine, 2004a).

The first example of an efficient redox-stable anode material, (La_0.75Sr_0.25) Cr_0.5Mn_0.5O₃ (LSCM), was developed in 2003 (Tao and Irvine, 2003), building on previous research into interconnect, LaCrO3, and cathode, La₁ _ _xSr_xMnO₃. materials (Minh, 1993). LSCM is a p-type conductor with a conductivity of approximately 38 S cm^− 1 in air and 1.5 S cm^− 1 at 900°C in 5% H₂/Ar (Tao and Irvine, 2004b). Similar performance was found when methane was used, with 3% H₂O, to that found for 5% H₂/Ar. The material also shows a high resistance to coking when using hydrocarbon fuels. Due to similarities between cathode and anode requirements, LSCM has been employed to form symmetrical/reversible fuel cells (Bastidas et al, 2006), which utilise the same compound for both anode and cathode.

Previous reports (Tao and Irvine, 2003) noted that coking was not observed when operating pure LSCM in methane, thus investigation into alternative hydrocarbon fuels was undertaken. When operating in ethanol a minor reduction in fuel cell performance was found when switching between humidified H₂ and a gasified 2:1 ethanol and water mixture (Huang et al., 2007). Stability testing in the gasified 2:1 ethanol and water mixture showed no significant performance degradation over 60 h, suggesting that carbon deposition during operation was minimal.

Kobsiriphat et al. (2010) found that substitution of nickel or ruthenium into LSCM introduced nanoscale metal particles onto the oxide surface upon anode reduction. According to previous reports (Madsen et al, 2007; Kobsiriphat et al, 2009), both nickel and ruthenium nanocluster formation was found to reduce the anodic polarisation resistance, albeit with much greater efficacy for ruthenium. Nickel substitution for manganese was also found to increase the catalytic activity of LSCM for total oxidation of methane, while substitution for either chromium or manganese increased the electronic conductivity in both air and 5% H₂/Ar (Jardiel et al, 2010). A maximum conductivity of 2.4 S cm^− 1 was found for La_0.75Sr_0.25Cr_0.5Mn_0.44Ni_0.06O3 in 5% H₂/Ar at 800°C, with a partial protonic contribution to the total conductivity observed under reducing atmospheres.

Fuel cell testing with a composite LSCRu-GDC anode yielded maximum power densities between 400 and 535 mW cm^− 2 at 800°C in humidified H₂ (Jardiel et al., 2010). Impregnation of LSCM with metallic Ru was also effective at improving the anode properties, with significant increases in the catalytic activity for propane reforming/oxidation (Barison et al, 2010). Additionally, the addition of Ru improved the H₂ and syn-gas (the mixture of H₂ and CO) yield, giving a maximum syn-gas yield of ~ 90% at 800°C, whilst maintaining good chemical stability.

Composite anodes, comprising of GDC and LSCM, exhibited reasonable catalytic activity for the oxidation of humidified H₂ and CH₄ without observed carbon deposition or microstructural alteration (Chen et al, 2007a). Further work by Chen et al. (2007b) demonstrated that impregnation of LSCM with GDC forms an effective anode material for operation in humidified H₂ and CH₄. Reasonable fuel cell performance, of 419 and 158 mW cm^− 2 at 800°C in humidified H₂ and CH₄, respectively, was observed.

Zhu et al. (2010a) discovered reasonable fuel cell performance for a LSCM-impregnated YSZ composite anode, a maximum power density of 567 mW cm^− 2 at 850°C in dry H₂, and little change in fuel cell performance when using dry CH₄, as a fuel. The fuel cell performance was significantly improved without any observed carbon deposition or degradation through impregnation of the composite with 2  wt% Ag and 6  wt% Ni due to the improved catalytic activities. Co-impregnation of a YSZ scaffold with LSCM and Ni improved the fuel cell performance compared with sequen- tial impregnation of YSZ (Zhu et al, 2010c). The enhancement was attributed to the differing microstructures observed and the obvious disparate distribution caused by sequential impregnation. The fuel cell performance of the co-impregnated Ni–LSCM–YSZ anode reached 1151 mW cm^− 2 in dry H₂ at 800°C (Zhu et al, 2010b). The fuel cell performances with LSCM–YSZ anode (Cell I) and Ni–LSCM–YSZ anode (Cell II) are shown in Fig. 14.3. Reasonable fuel cell performance was also observed when operating with dry methane, with no drop in open-circuit voltage (OCV), carbon deposition or compound deterioration noted after 6 h of operation. It should be noted that longer term testing is required for real applications.

Impregnation of Pd nanoparticles in a LSCM/YSZ anode was found to improve fuel cell performance slightly in humidified H₂ and significantly in humidified CH₄ and ethanol-saturated nitrogen (Jiang et al, 2008). The lack of carbon deposition on the anodes and the improved fuel cell performance leads to the conclusion that Pd promotes the partial oxidation reaction of both methane and ethanol. Further increases in fuel cell performance were noted with the addition of a metal catalyst, cerium oxide and the use of a LSCF–YSZ cathode (Zhu et al., 2010d). The addition of 0.5  wt% palladium and 5  wt% cerium oxide produced the highest maximum power density of 520 mW cm^− 2 at 700°C in humidified H₂.

Lu et al. (Lu and Zhu, 2007) produced a series of Cu and Pd impregnated LSCM anodes with optimal Cu loading of 20  wt%. Additional loading of 1.5  wt% Pd showed a slight improvement of the maximum power density in dry H₂. However, a significant increase was noted for dry CH₄. No additional increase in performance was noted for further impregnation of Pd. The maximum power density achieved at 850°C was 890 and 600 mW cm^− 2 in dry H₂ and dry CH₄, respectively. No carbon deposition was observed upon operation, although exposure of the anode to 50 ppm H₂S initiated a significant drop in fuel cell performance, attributed to the low sulphur tolerance of LSCM.

Ye et al. (2008) investigated Cu-LSCM-ScSZ composite anodes, achieving poor maximum power densities albeit with good stability when operating with either H₂ or a 2:1 mixture of C₂H₅OH–steam fuel. Grading of the anode

was used to increase the anodic performance, producing a complicated mul- tilayer anode: Cu–LSCM–YSZ/Ni–YSZ/Ni–ScSZ. An increase of ~ 5 × was noted in the maximum power density, reaching 534 and 384 mW cm^− 2 at 800°C in H2 and 2:1 C₂H₅OH–H₂O, respectively. Stable operation over 120 h was observed in the ethanol-steam mixture without any observed carbon deposition.

Ni–LSCM–YSZ composite anode has been successfully demonstrated in a single-chamber SOFC. A maximum power density of 285 mWcm^− 2 is achieved at a gas composition of CH₄:O₂ approximate to 2:1 and a furnace temperature of 800°C. The redox cycling test demonstrates that the stability of the LSCM-based anode is better compared to the conventional Ni–YSZ anode (Zhu et al, 2010e).

14.4.2 SrTiO ₃-based anode materials

Another material which has been under extensive investigation is stron- tium titanium oxide (STO). In 1997, Slater and Irvine first proposed to use La-doped SrTiO₃ (LST) as an anode for SOFCs (Slater et al., 1997). Early research into B-site doping of LST found that niobium doping produced a redox-stable material with a metallic-type electronic conductivity of ~ 7  S cm^− 1 at 930°C (Slater et al. 1997), though, as in LST, the catalytic activity for H₂ oxidation is not good (Blennow et al, 2009). Research in the last few years has been dedicated to improving the electronic conductivity and the catalytic activity of anode materials based on SrTiO₃.

These materials have shown a high sulphur tolerance, resistance to coking and have good chemical and redox stability (Mukundan et al, 2004). The main difficulty with the use of STO is low electronic conductivity in a reducing atmosphere (Balachandran and Eror, 1981).

A-site doping of yttrium in SrTiO₃ (forming YST) has been shown to sig- nificantly improve the electronic conductivity, up to 73.7  S cm^− 1 at 800°C in forming gas with good thermal stability (Li et al, 2007). Lanthanum substitution on the strontium site (forming LST) was also found to significantly increase the electronic conductivity in a reducing atmosphere, to 16  S cm^− 1 at 1000°C. However, the catalytic activity for the oxidation of H₂ was unsatisfactory (Marina et al, 2002).

Recent reports have found significantly higher electronic conductivity for Sr₀.₉₄Ti₀.₉Nb₀₁O₃ > 120  S cm^− 1 at 1000°C in 9% H₂/N₂; however, these compounds were sintered at high temperature in a reducing atmosphere and had a high pellet density (Blennow et al., 2008). Gross et al. (2009) found similar conductivity values for dense pellets of SrTi_0.98Nb_0.02O₃ sintered under similar conditions. However, the enhanced conductivity was found to be nullified after oxidation at 1200°C.

A composite niobium-doped strontium titanate (SrTi₀.₉₉Nb₀.₀₁O₃) − yttrium-doped zirconate anode demonstrated stable conductivity of ~ 1  S cm^− 1 in humidified H₂ at 800°C over successive redox cycles. Infiltration of this material with 1  wt% Pd and 3  wt% CeO₂ produced an anode material with promising fuel cell performance, 640 mW cm^− 2 at 800°C in humidified H₂ (Gross et al, 2009). Further A-site doping was investigated by Vincent et al. (2010) with the introduction of Ba to LST. Increases in the amount of Ba were correlated with increases in fuel cell performance in dry hydrogen. However, the opposite effect was observed in dry methane. Introduction of 5000 ppm of H₂S was observed to significantly increase the fuel cell performance in both dry H₂ and dry CH₄. Additionally, exposure to hydrogen sulphide did not have any observable effect on the cell microstructure. Despite all of this, fuel cell performance was universally poor, with a maximum power density of 60 mW cm^− 2 at 900°C in 0.5% H₂S/H₂.

Introduction of A-site deficiency, with maximum deficiency forming (Y_{0 .08}Sr_0.92)_0.94TiO₃, was found to increase the ionic conductivity at the expense of the electronic conductivity (Zhao et al., 2009). The same trend is also found with increasing B-site deficiency in Y₀.₀₈Sr₀.₉₂Ti_{1 − x}O₃, with the maximum deficiency occurring for the x = 0.05 compound (Gao et al, 2008). The introduction of A-site deficiency retained greater electronic conductivity than introduction of B-site deficiency, 73.7  S cm^− 1 compared to ~ 35  S cm^− 1 at 800°C in a reducing atmosphere, while achieving similar ionic conductivity. Stability testing of both the A-site and B-site deficient compound with YSZ found that no impurity peaks were observed after firing at temperature > 1200°C for 10 h in forming gas. Previous uncertainty over the conductivity of the A-site deficient YST was ascribed to different synthesis methods and the purity of the reagents (Lu et al, 2007).

Further work into A-site and B-site deficient YST was undertaken by Ma et al. (2010), confirming that changing the sintering temperature of the samples did not affect the conductivity of the A-site deficient samples but increased that of the B-site deficient samples. Redox cycling of the samples showed a significant reduction in the conductivity of the B-site deficient samples, with a lesser effect observed for A-site deficiency. As noted in previous research (Fu and Tietz, 2008), the conductivity of A-site deficient samples maintains acceptable electronic conductivity despite the reduction upon successive redox cycles.

Similar trends in ionic and electronic conductivity were found for introduction of A-site deficiency in LST (Li et al., 2010). The electronic conductivity was found to be superior to that found for A-site-deficient YST with the expected consequence of lower ionic conductivity. An additional effect of A-site deficiency was the promotion of pellet densification and an improvement in the thermal stability of the compound. Pre-reduction of the compound was found to enhance the conductivity, albeit only for the first redox cycle. Preparation of a composite anode with A-site-deficient LST, Cu and GDC gave a maximum power density of 550 mW cm^− 2 at 750°C in humidified H₂ (Savaniu and Irvine, 2010).

Formation of composite anode materials from La-doped STO (LST) and CeO₂ exhibited increased fuel cell performance, with maximum power densities of 172.3 mW cm^− 2 in H₂ and 139.6 mW cm^− 2 in C₂H₆ at 900°C, despite a reduction in the electronic conductivity of the anode (Sun et al, 2008). Incorporation of Ni into an LST–GDC anode was found to significantly increase the fuel cell performance in comparison to the LST–GDC cermet (Yoo and Choi, 2009). The highest performance was found for Ni–50% La_0.2Sr₀.₈TiO₃–50% Gd_0.2Ce_0.8O₂ which had a maximum power density of 275 mW cm^− 2 at 800°C in humidified H₂. Unfortunately, the anodic polarisation resistance was found to significantly increase within 10 h of operation, which was attributed to nickel agglomeration due to the high level of nickel components.

Fu et al. (2007) prepared a composite Ni–YST–YSZ anode with a stable electronic conductivity of ~ 10  S cm^− 1 at 800°C in 4% H₂/Ar. Impregnation of the nickel into the porous YST–YSZ framework was found to give good dimensional stability upon redox cycling, as the volume change of the nickel was nullified due to homogenous NiO distribution throughout the pores of the framework.

Pd was also investigated for use in LST–YSZ cermets as a promoter of catalytic activity (Kim et al., 2008). Pd and CeO. impregnated LST–YSZ (75:25) composite anodes demonstrated reasonable fuel cell performance, 570 mW cm^− 2 at 800°C in humidified H₂, despite the low conductivity values (Kim et al, 2008). Further increasing the amount of CeO₂ impregnated into the scaffold did not have a significant effect on fuel cell performance, suggesting catalytic performance was not the limiting factor in increasing fuel cell performance.

Further development of strontium titanates by Ruiz-Morales et al. (2006) in 2006 led to the formation of LaoSr₈Ti₁₁Mn₀.₅Ga₀.₅O_38-δ (LSTMG). No changes in OCV or electrode performance were observed upon redox cycling at 850°C and 950°C for 48 h, with maximum power densities of ~500 and 350 mW cm^− 2 in wet H₂ and wet CH₄, respectively (Ruiz-Morales et al, 2007a). Additionally, no carbon deposition was detected after operation in wet CH₄. Substitution of Ga by Mn was found to reduce the degree of oxy- gen loss on reduction and, subsequently, reduce the electronic conductivity. The maximum electronic conductivity, 7.9  S cm.' at 900°C in 5% H./Ar, was noted for the pure gallium sample, although the TEC of the 1:1 Mn-Ga sample was closest to that of most SOFC electrolytes (Escudero et al., 2009). The best reported performance of SrTiO₃-based anode was 1066 mW cm^− 2 at 850°C when H. was used as fuel and 1.5  wt% Pd was impregnated at the anode, with La₀.₈Sr_0.2Ga₀.₈₃Mg₀₁₇O_{3 − δ} and Sm₀.₅Sr₀.₅CoO_{3 − δ} as electrolyte and cathode (Fig. 14.4) ' (Lu et al, 2009b).

14.4.3 Double perovskite-based anode materials

The conductivity and redox stability of double perovksite oxides such as Sr₂FeNbO₆ (Tao et al., 2004) and Sr₂MnNbO₆ (Tao and Irvine, 2002b) as potential anode materials for SOFCs have been investigated by Tao and Irvine since 2002. Double perovskites such as Sr.MgMoü₆ (SMM) were investigated for anode applications by Huang et al. (2006). Recent research investigated Sr₂MgMoO₆ and Sr₂MnMoO₆ and determined that both compounds have electronic conductivities of ~ 10  S cm^− 1 in both H₂ and CH₄ at 800°C, in addition to redox stability and stability to sulphurous impurities (Huang et al., 2006). Fuel cell testing established that the maximum power densities in CH. were 438 and 118 mW cm^− 2 for Sr.MgMoO₆ and Sr₂MnMoO₆, respectively (Huang et al, 2006). Huang et al. investigated substitution of Mg in SMM yielding both Sr₂NiMoO₆ and Sr₂CoMoO₆ (Huang et al., 2006). The total conductivity for Sr.CoMoO₆ was low, ~ 1  S cm^− 1 at 800°C in H₂. However, fuel cell performance in H₂. reached a maximum power density of 735 mW cm^− 2. Fuel cell performance in dry methane was significantly lower than that in wet methane, suggesting that the anode is active for internal reforming of methane, although the performance was still lower than in H₂. Power densities of 1017 and 634 mW cm^− 2 have been achieved at 800°C when H₂ or CH₄ was used as the fuel, respectively, based on a cell with Sr₂CoMoO₆ anode, La₀.₈Sr₀.₂Ga₀.₈₃Mg₀.₁₇O₂.₈₁₅/La₀.₄Ce₀₆O_{2 − δ} as electrolyte, and SrCo₀.₈Fe₀₂O_{3 − δ} as cathode (Fig. 14.5) (Zhang et al, 2011a). Despite the promising fuel cell results, further research into Sr.MgMoO₆ found degradation of the compound at 900°C in a reducing atmosphere (Bernuy-Lopez et al, 2007). Doping with La was found to improve the fuel cell performance, giving a maximum power density of 550 mW cm^− 2 at 800°C in wet CH₄. However, subsequent research discovered compound degradation upon reoxidation (Ji et al, 2007, Marrero-Lopez et al, 2009).

Research by Bernuy-Lopez et al. suggested that SMM may have limited redox stability at temperatures above 900°C (Bernuy-Lopez et al, 2007). This has been refuted by various authors who have suggested that the compound is stable and the RP-phase observed was due to the high-temperature synthesis method (Marrero-López et al, 2008). High chemical reactivity was also observed for other molybdenum-containing compounds, for example La₂Mo₂O₉ electrolyte materials (Marrero-López et al., 2010). Sintering of SMM with YSZ at 1000°C was found to form SrMoO₃, SrZrO₃ and MgO while sintering with LSGM at 1000°C produced SrMoO₄ and (Sr,La)Ga₃O₇ (Lu et al, 2009a). No significant interfacial reactivity was observed for SMM and GDC and fuel cells using a GDC interlayer showed improved performances.

Zhang et al. (2010) found that Fe substitution in SMMO increased the fuel cell performance. However, the improvement was less than was found for Ni or Co substitution. Substitution of strontium with calcium was found to increase the conductivity of the sample, although degradation of the compound was observed in an N. atmosphere at ~ 400°C. Sr.FeMoO₆ had the highest fuel cell performance, 835 mW cm^− 2 in dry H₂ at 850°C, and had 20 h of stable operation in commercial city gas, which contains impurities of H₂S, ammonia and naphthalene (Zhang et al, 2010).

Substitution of Mo in SMM by W^VI and Nb^V was attempted by Vasala et al. (2010a). Substitution with either niobium or tungsten decreased the conductivity in 5% H₂/Ar, with greater reductions noticed for tungsten than niobium. All compounds synthesised were found to be redox stable, although the range of oxygen content variation decreased with increased doping. The same authors also investigated the effect of substitution of Mg with various transition metals (Vasala et al., 2010b). Only the compound with pure magnesium content was stable in both oxidising and reducing conditions at temperatures greater than 900°C, with those with Mn or Fe stable in reducing conditions only and those with Co, Ni or Zn stable in oxidising conditions only. These data do not correlate with the previously determined stability of Sr₂NiMoO₆ (Wei et al, 2008). However, no plausible explanation for this discrepancy could be proffered by the authors.

14.4.4 Other perovskite-based anode materials

Lanthanum strontium cobalt ferrite (LSCF) is widely used as a SOFC cathode material. However, in recent years it has become of interest for anode applications due to its catalytic activity for the electrochemical oxidation of methane (Fisher et al, 2009). As LSCF does not have significantly high ionic conductivity, Haag et al. (2008) produced a composite anode of LSCF and GDC. A maximum power density of 365 mW cm^− 2 was achieved at 800°C under humidified H₂. LSCF is known to be at the limits of its stability and low conductivity under anodic conditions.

The activity of LSCF cermets with alternative fuels, such as propane and syn-gas, was investigated by Lo Faro et al. (2009a). The fuel cell performance using syn-gas increases with decreasing amounts of LSCF in LSCF–GDC cermets and increases with increasing proportions of carbon monoxide (Huang et al, 2009). When propane was used as a fuel, a maximum of 300 mW cm^− 2 was observed for the 7  wt% Ni-impregnated LSCF–GDC cermet. Gradual increases in the fuel cell performance are observed over 100 h of operation, although the anode was found to partially degrade into La₂NiO₄ and SrFe_{1 − x}Co_xO₃ (Lo Faro et al, 2009b).

Calcium-doped lanthanum chromite, a common interconnect material, was investigated as a potential symmetrical electrode material due to high electronic conductivity in both oxidising and reducing atmospheres and limited catalytic activity for both the anode and cathode reactions (Ruiz-Morales et al., 2007b). Symmetrical fuel cell performance with La_0.7Ca_{0 3}CrO₃ as both anode and cathode, with a maximum power density of 110 mW cm^− 2 in humidified H2 at 950°C, was underwhelming when compared to the performance of symmetrical LSCM electrodes.

Xu et al. utilised separate electronic conducting, ionic conducting and catalytically active components to form a cermet anode: LaCrO₃, YSZ and VO_x. respectively (Xu et al., 2007). Fuel cell performance was significantly improved with the addition of CO to the H₂ feed gas, with a maximum power density of 190 mW cm^− 2 at 850°C in syn-gas. Addition of H₂S did not affect fuel cell performance and XRD showed no changes to the anode structure after firing in syn-gas for 24 h at 900°C. Extended fuel cell operation exhib- ited reduction of the maximum power density of the cell over time, with rapid degradation after 10 h at 900°C. Carbon deposition was observed through XPS and is expected to be the cause of the performance degradation.

La_0.7Ca_0.3CrO₃ (LCC)–Ce_0.8Gd_0.2O_1.9 (GDC) composites have also been investigated as symmetrical electrodes for SOFCs on LSGM electrolyte. The electrical conductivity of the LCC-GDC composites decreases with increasing GDC content. The best electrical conductivities of 18.64  S cm^− 1 in air and 1.86  S cm^− 1 in H₂ at 850°C are achieved for an 80  wt% LCC-20  wt% GDC (LCC–GDC20) composite. The maximum power density is 573 mW cm^− 2 in dry H₂ and 333 mW cm^− 2 in humidified commercial city gas containing H2S at 900°C, respectively. These results suggest that the LCC–GDC20 composite can potentially serve as an electrode for symmetrical SOFCs operated on H₂ and commercial city gas containing H₂S (Zhang et al, 2011b). The typical fuel cell performances of SOFCs based on perovskite anode are listed in Table 14.2.

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