1.4.1
Materials requirements
In developing highly functional materials for intermediate temperature fuel cells and related applications, a number of requirements must be met. The total area specific resistivity (ASR) of the electrolyte should ideally be no more than 0.15 Ω cm2, with ionic conductivities >10–2 S cm–1 at the desired operating temperature. Low electronic conductivity and good sinterability is essential to reduce fuel cross over and short circuiting. The electrodes must have high electronic conductivity and be active to electro-oxidation or reduction whilst simultaneously be able to trans- port ions and reactant gases. Common to all components is the requirement for mechanical, chemical and thermodynamic stability over the temperature range and conditions of operation. SOFC’s experience thermal and chemical potential gradi- ents, particularly across the electrolyte, of which must therefore maintain a broad electrolytic domain. Materials should therefore be engineered with high enthalpies of formation with respect to their constituents or potential reaction products with either fuel, oxidant or trace impurities (e.g. SO2). Furthermore chemical compati- bility is needed between electrodes and electrolyte to prevent unwanted interfacial reactions that might generate blocking phases. Thermally induced fracturing dur- ing cycling can be reduced by ensuring the thermal expansion coefficients (TEC) of each component match, and mechanical stability can be promoted by tailoring stack design and selecting materials with high ceramic strength. Processing amenability is therefore essential, for example thin film or porous structures should be fabricated with ease and at low cost. Meeting these requirements begins with engineering oxide structures with the appropriate transport properties, and which can accommodate high levels of intrinsic or extrinsic disorder.
1.4.2
Electrolyte Materials
Fluorite structured oxides were the first solid oxide membrane materials to be devel- oped, and to this day constitute the electrolytes of choice for the testing of prototype fuel cells. The most well studied is that of 8 mol% yttria stabilised zirconia (8% YSZ), the properties of which remain the benchmark for continued research. The cubic fluorite (AO2) structure consists of an fcc array of large 8 coordinate A site cations with O2– in tetrahedral sites. The structure is stable when the cation-anion ratio (R+/R–) is close to 0.7, and A is around ≈1˚A which is typical of early lan- thanides such as CeO2(Ce4+=0.97˚A) [45]. Zr4+is too small to sustain the structure, and ZrO2 forms 3 polymorphs with monoclinic (25-1170°C), tetragonal (1170-2370
°C) and cubic symmetry (>2370 °C) [46]. Progressive doping with larger aliovalent cations e.g. rare earths or alkaline earths, stabilises the tetragonal (≈2-2.5%) and cubic forms (≈8-12%) to room temperature and generates extrinsic disorder, high levels of which can be accommodated [47][48][49].
Numerous dopants have been incorporated into the structure of CeO2 and ZrO2,
the conductivity depends on both dopant concentration and A site radius, however is wholly ionic (t ≈1) under oxidising conditions [47][48][49][50]. The conductivity increases as the dopant radius approaches that of the host, the highest level was recorded for 10% Sc3+ and Y3+ doped ZrO2 of 10–1 S cm–1 and 3x10–2 S cm–1 at
800 °C respectively and 20% Gd3+ doped CeO2 (CGO) of 1.5x10–1 S cm–1 at the same temperature [50][51]. The conductivity of Ce0.8Sm0.2O3–δ is 5 times higher
than that of YSZ at the same substitutional level [52]. With increasing doping level the conductivity reaches a maximum and tails off, attributable to the formation of defect vacancy associates. As Figure 1.8 shows, the magnitude of the decline reduces as the A site radius approaches that of the host, which in turn reduces the association binding energy and activation energy [53][54].
Despite the higher conductivity of CGO, the transport number decreases steadily at oxygen partial pressures below 10–5atm, eventually reaching t=0 at 10–20atm (at 800°C); YSZ however has a wide electrolytic domain [55]. Under reducing conditions Ce4+ readily reduces to Ce3+ to become an n-type electronic conductor, with small polaron hopping between the cerium sites. Although the electronic contributions are low below 700 °C, the redox flexibility of cerium causes cell volume expansion- contraction and subsequent microcracking. Experimental and computational studies have shown both YSZ and CGO to have high grain boundary resistance, and dopant segregation, and so the interfacial regions are particularly susceptible [56][57]. At- tempts to reduce this phenomenon were performed by adding a thin layer of YSZ at the anode side of CGO electrolytes, however this resulted in reactivity between the components [45].
One of the most ionically (O2–) conductive materials discovered is δ-Bi2O3, which also belongs to the fluorite family. Bi2O3 is polymorphic forming a monoclinic α- phase at room temperature which converts to the cubic δ- form at 729-824°C and is accompanied by a 3 order magnitude increase in conductivity to 1 S cm–1 and 2.3 S cm–1 at 730°C and 800 °C respectively [59][60]. On cooling Bi2O3 exhibits thermal hysteresis, and either metastable tetragonal β- or cubic γ- phases result depending on the cooling regime. δ-Bi2O3 has high intrinsic disorder, with a random distribu-
Figure 1.7: A: Comparison of the ionic conductivity of several fluorite and perovskite structured oxides. Adapted from [58].
tion of oxygen vacancies on 25% of sites, which results in displacement of half the occupied tetrahedral interstices. The sharp increase in conductivity between the α- and δ- forms results from disordering of vacancy clusters and is accompanied by a large volume expansion [61]. Many cations can be substituted for Bi3+ to extend the cubic domain to room temperature (e.g. 22-27 mol% WO3 or 25-43 mol% Y2O3)
[60][61]. However this is at a detriment to conductivity, which further declines with additions beyond that required to stabilise cubic symmetry, and results in vacancy ordering along the (111) plane below 600 °C [62]. The highest conductivity was re- ported for Bi0.8Er0.2O1.5 of 0.23 S cm–1 at 650 °C, however the substituted phases gradually age and form a vacancy ordered rhombohedral phase with significantly lower conductivity [61]. The use of Bi2O3 in devices is further hampered by elec-
tronic conductivity under low p(O2), low strength and Bi3+ volatilisation [63]. From the early work of the fluorite system, a number of empirical relations were identified that could be used as a predictor of fast ion transport [64][65][66][67]. To maximise the pre-exponential factor of the conductivity equation, the structure must accommodate a high concentration of mobile charge carriers. This is deter- mined by both the concentration of intrinsic or extrinsic defects and interactions between defects and the host lattice, which effectively trap the charge carrier at low temperature [67]. Many types of associations are possible which may be present at both low and high dopant concentrations, but may become significant leading to long or short range ordering, clustering and superstructure formation such as the case of YSZ, CGO and particularly δ-Bi2O3 [68][69]. Such associations add an extra
Figure 1.8: Activation energy of Ce1–xAxO2–δ as a function of dopant concentration.
A number of studies have shown associate formation results from both electro- static and elastic effects, and depends on both dopant concentration and type [69][66][70][71]. In fluorite systems, there has been a lot of debate as to whether the vacancy is associated with host or dopant, however a number of experimental and computational studies have suggested the mechanism varies depending on whether elastic or electrostatic terms dominate for a given composition [68][70][71][54]. For example the activation minima that accompanies substitutional concentrations cor- responding to the conductivity maxima was ascribed to the formation of dimers at low concentrations via electrostatic interactions, and higher order clusters at high dopant concentrations [71]. The activation energy minima observed as a function of dopant radius for a given substitutional level was attributed to a minima in the asso- ciation binding energy where a cross over from interactions involving first and second nearest neighbours or dopants occurs as the radius increases [70][54][72]. Therefore fast ion conduction is achieved with careful dopant selection and consideration of the concentration which maximises charge carrier density and mobility.
Both 8 mol% YSZ and 20 mol% CGO offer a good compromise, although associa- tions are still present requiring heating above 600 °C and 400 °C respectively [73]. Lattice symmetry has been advocated as being essential for maintaining energetic equivalence of the oxygen sites that are important for oxygen ion mobility, and new materials should adopt symmetry that is close to cubic [74]. The exceptionally high conductivity of δ-Bi2O3 highlights the importance of lattice softness, larger
more polarisable cations particularly those with active lone pairs deform more eas- ily reducing the migration barrier that the O2– ion has to overcome as it passes the ’bottleneck’. This lead to the notion that candidates must have open structures with weak bonding energies, and therefore free volume or specific free volume (SFV)
which can also be considered a measure of softness, should be maximised. δ-Bi2O3 has a low melting point (804 °C) and a higher SFV of 0.5 compared to 20 mol% CGO (0.38) and 20 mol% YSZ (0.31) [74]. The early work was not restricted to fluorite structures; many SOFC cathodes were developed by manipulating oxides with the perovskite structure (ABO3) which can accommodate numerous charge
and radii combinations on both the A and B site. The structure is composed of large A cations in cubo-octahedral holes and smaller corner sharing BO6 octahedra. When A is of similar radius to oxygen (1.4 ˚A) and B ≈0.58 ˚A cubic symmetry is typically adopted. As the radius of A decreases and B increases, octahedral tilt- ing reduces the symmetry to orthorhombic or rhombohedral, whilst the opposite generates hexagonal or tetragonal symmetry [75]. Distortions can be quantified by a tolerance factor, for cubic perovskites t ≈1 [76]. To date only orthorhombic LaGaO3 has been found to be suitable for electrolyte applications; the conductivity of La0.9Sr0.1Ga0.8Mg0.2O3–δ (LSGM) approaches that of CGO (0.1 S cm–1 750 °C)
whilst offering a wide ionic domain (1-10–20 atm) and low TEC [77][78][63]. LSGM also shows proton conductivity of 5 x 10–2 S cm–1 at 600°C [79]. A disadvantage of the LaGaO3 system is both the cost and volatilisation of Ga2O3 and its reactivity
with electrodes particularly NiO anodes [63].
Further doping was shown to lead to defect association, and the optimum oxygen deficiency was shown to be δ=0.2 [63][80]. It was noted that the conductivity de- creased quite linearly with tolerance factor as the A site was substituted in the order of A=Sr>Ba>Ca and A=Nd>Sm>Gd>Yb>Y [63][81]. However based on symme- try alone, cubic La0.9Sr0.1Al0.9Mg0.1O2.9 (LSAM) should offer higher conductivity,
which it does not [82]. Studies have shown a generalised positive correlation between conductivity (specifically activation energy) and SFV which should be between 30-35 ˚
A, and critical radius (rc), which is the saddle point formed by two A site cations and
one B site cation that the O2– must pass [83][84][83][85]. As such indium deriva- tives of LSGM should offer the best transport properties. These parameters are often competing however, for example SFV and rc often have an inverse relationship
to t, whilst oxygen vacancies increase the free volume but also offer higher charge carrier concentrations [85][80]. A more extensive survey of various perovskite com- positions showed that the conductivity was highest at t=0.96 which offers a good trade-off between symmetry (t) and free volume, and might explain the properties of LSGM (t=0.964) [80][74]. Other factors such as bond energy, cation polarisabil- ity and charge were shown to correlate with conductivity, and in reality a complex interplay between many factors determines the ionic transport in these structures [83][80][74][64]. As a result of these early empirical findings, a number of other cubic systems were investigated.
For example LAMOX, which belongs to the La2Mo2O9 family adopts a monoclinic superstructure at room temperature and is structurally similar to β-SnWO4. On
heating above 580 °C, LAMOX undergoes an order-disorder transition resulting in considerable volume expansion and affording cubic symmetry. This is accompa- nied by a 2 order increase in ionic conductivity (0.03 S cm–1 at 720 °C) and high oxygen diffusivity; 1.41 x 10–6 cm2 s–1 at 800 °C (La1.7Gd0.3Mo2O9) compared to La0.8Sr0.2Ga0.8Mg0.2O3–δ (4.13 x10–7 cm2 s–1) and Zr0.81Y0.19O2–δ (6.2x10–8 cm2 s–1) [58][86]. However molybdenum is sensitive to reduction; a number of dopants including bismuth have been incorporated to overcome this and depress the transi-
tion [58]. Despite these investigations YSZ remains the best candidate due to its low TEC, reasonable chemical and mechanical stability and large electrolytic domain.
1.4.3
Cathode electrode
Because the perovskite structure can tolerate extensive modifications, many ABO3
materials display semi-conducting or metallic behaviours desirable for electrode ap- plications. These properties typically rely on the electronic compensation of accep- tor dopants through the incorporation of transition metals (TM) with flexible redox couples. The charge transfer mechanism which proceeds by polaron hopping along TM-O-TM chains, depends on both the crystallographic and electronic structure, and the nature of the dopant. However, the development of MIEC cathodes is desir- able to enhance reduction kinetics and reduce polarisation losses by extending the effective triple phase region; where the gas-MIEC interface enables the simultaneous reduction and bulk diffusion of oxygen. Therefore it is desirable to enhance both the oxygen surface exchange and diffusion coefficients, and the measurement of both these parameters provides a tool for predicting high cathode performance.
Many simple perovskites have been extensively studied (LaCoO3, LaMnO3,
LaFeO3), however to obtain the best performance, these materials (and the other components) in the fuel cell must operate at temperatures in excess of 800°C [85][87]. Therefore attention has focussed on the development of cathodes with lower polari- sation resistance at intermediate temperatures. LaMnO3 perovskites have remained the state of the art cathode for many years, owing to their good intrinsic p-type con- ductivity which is enhanced on acceptor doping with alkaline earth cations. Charge compensation by Mn4+ affords conductivities in the region 2-300 S cm–1 at 900 °C for La1–xSrxMnO3–δ [88]. However its main limitation is the increase in TEC
with strontium doping and its reactivity with YSZ upon sintering to form insulating pyrochlore phases [89][90]. The latter can be minimised by adding a layer of YSZ between the cathode, and substitution of lanthanum with smaller cations Pr, Nd and Sm were shown to enhance stability [91][89]. Due to the poor ionic conductivity of LSM, YSZ composite formation is necessary [92].
Interest in MIEC cathodes initially focused on La1–xSrxCoO3 (LSC), which offers
good oxide ion conductivity and very high electronic conductivity of 1600 S cm–1 at 800 °C [93]. However it has many disadvantages that preclude device application. Although stable with CeO2 electrolytes it is much more reactive with YSZ, has a larger TEC and also undergoes both an oxygen vacancy order transition and metal insulator transition at p(O2) >1x10–3 atm [94][95]. Replacement of the A site with Gd and Sm was shown to improve the stability of LSC, and doping the B site with iron reduces thermal expansion [96]. La0.6Sr0.4Co0.8Fe0.2O3–δ (LSCF) in particular
has attracted attention owing to the materials reasonable electronic conductivity (350 S cm–1) and increase in ionic conductivity due to the stability of iron to ox- idation. Replacement of La for Nd increased the conductivity to 600 S cm–1 [93]. Both LSCF and the barium derivative Ba0.5Sr0.5Co0.8Fe0.2O3–δ (BSCF) are more suited to intermediate temperature applications owing to their high conductivity and good surface exchange properties [97]. The degree of orbital overlap is depen- dent on both the electronic structure of the cation and degree of distortion from
cubic symmetry [85]. Although comprehensive relations between charge transfer and structural features in the context of cathode perovskites are not well reported, deviation from cubic symmetry is shown to reduce the TM-O-TM angle from 180o and increases the band gap. For example the transition to metallic behaviour of LSC and the increase in Mn4+ content on strontium doping of LSM, both cause the angle to approach 180o, and are accompanied by an increase in electronic transport [85]. However unlike ionic transport in perovskites, models have not been devised to predict structures which might meet these conditions. Investigation of other per- ovskite structured cathodes has focused on triple doping the B site and derivatives of LaAlO3 which offer little added benefit [87]. The conductivity of LaNi0.6Fe0.4O3 however is 3 times that of LSM and La0.6Sr0.4FeO3–δ was shown to offer very good
diffusion and surface exchange properties [85]. Attention has mostly been focussed on optimisation and overcoming issues relating to Cr poisoning, dopant segregation and electrolyte reactivity. More recently there has been growing interest in novel layered or cation ordered non-perovskite structures (Section 1.5).
1.4.4
Anode electrode
Early anodes for high temperature fuel cells were based on porous metallic con- structions e.g. Ni, Ru, Pt, Fe, Mn, Co of which Ni had the highest electrochemical activity for H2 oxidation. Pure nickel has a high electronic conductivity (2 x 104 S
cm–1 at 1000 °C) however displays a large TEC and suffers from grain growth and aggregation which greatly reduces the triple phase region [98][99]. Addition of YSZ to form a composite ’cermet’ electrode supresses both these issues whilst extending the surface area and reaction zone to 3 dimensions which enhances the reaction kinetics [98].
To this day Ni/YSZ cermets are used as SOFC anodes, owing to their good elec- trocatalytic properties, appropriate mechanical strength and low relative cost. A number of methods have been used to prepare composites and precursor NiO-YSZ powders including, dry pressing, co-precipitation and hot moulding, however con- ventional ceramic processing of NiO-YSZ powders which are then reduced during fuel cell operation is most common [100]. The electrical conductivity of Ni/YSZ depends on a number of macro and microstructural features including the volume fraction, spatial distribution, and size of each component in addition to porosity and thickness of the electrodes [98][100]. This is further dependant on the starting materials and processing conditions.
For example the percolation threshold for a random distribution of particulate spheres in 3D space is 33% of the volume fraction, ideally the pore and Ni phases should exceed this to ensure a continuous network for electrons and gas molecules to reach electrolyte interfaces. The threshold can be reduced by manipulating the distribution of phase by altering the shape and size of the particles of each phase [98][100]. The highest conductivities are reported for fine NiO and coarse YSZ par- ticles, the former of which cluster around the later to generate a continuous pathway whilst also maintaining a good YSZ volume fraction [101]. Porosity is essential to promote triple phase boundaries and can be promoted with pore formers (e.g. car-
bon fibre), however sufficient connectivity must be maintained between each phase [98][100].
A major advantage of solid oxide fuel cells is the potential for internal reforming which greatly reduces the complexity, size and cost of external/pre-reforming con- figurations and enables the integration of current energy carriers such as natural gas. A number of catalytic routes are available for producing H2, however steam reforming followed by the water-gas shift reaction is the most common, and can be utilised directly over Ni/YSZ anodes. However, high steam-carbon ratios (S/C) in excess of 2 are required to prevent carbon deposition which results from the ac- tivity of nickel towards hydrocarbon cracking, which subsequently dilutes the fuel [102][98][99]. Nickel also suffers from sulphur poisoning which is a common contam- inant [102]. It is the desire of the research community to develop anodes which can oxidise hydrocarbon fuels directly (in the absence of steam) and which have good car- bon tolerance whilst simultaneously operating at intermediate temperatures [103]. Numerous studies have sought to lower Ni/YSZ cermet susceptibility to hydrocar- bon cracking and sulphur poisoning. The focus has been towards replacing nickel with copper due to its inactivity towards catalytic carbon formation, and YSZ with CeO2 which is an effective oxidation catalyst and offers high resistance to carbon deposition [104][99]. A number of MIEC perovskite and fluorite structured oxides have been investigated as single phase anodes or components of composite anodes [99]. Of interest is La1–xSrxCrO3, which has good electronic conductivity and sta-