Proton-conducting ceramic membranes for solid oxide fuel cells and hydrogen (H2) processing
The reduction of greenhouse gas emissions such as CO2 is a major subject of current socioeconomic, environmental and political discussion. Of a special interest is electricity production using alternative energy sources, as well as reduction in the CO2 levels emitted by existing electricity generation power plants fuelled by coal or gas. This goal can be achieved by developing innovative technologies that make use of ceramic membranes. Proton-conducting ceramic membranes are relevant for a variety of energy-related applications and hence it is necessary to develop materials with improved functional properties and durability. The present chapter will review the fundamentals and relevant application-oriented properties of different proton-conducting ceramic materials for membrane manufacturing for a variety of applications.
A reduction in CO2 emissions from power plants fired by coal or gas is a major subject in current socioeconomic, environmental and political discussion on reducing greenhouse gas emissions such as CO2. This goal can be achieved by introducing gas separation techniques making use of membrane technology that has significantly lower efficiency losses compared to conventional separation techniques.
Depending on the kind of power plant process, different membrane types (ceramic, polymer, metal) can be implemented to accomplish specific separation tasks. CO2/N2 separation is the main target in the post-combustion process. Air separation (O2/N2) is the focus of the oxyfuel process. In the pre-combustion process, an additional H2/CO2 separation is included. All separation concepts imply different process requirements but what they have in common is the need for a membrane with high permeability, selectivity and stability.
In addition to separation tasks, ceramic membranes are also used as the solid dense electrolyte in proton-conducting solid oxide fuel cells (PC-SOFCs). PC-SOFCs represent an alternative technology for clean electricity production that offers benefits in terms of fuel utilization, overall efficiency and system simplicity with reformed fossil fuels as well as hydrogen from renewables. A PC-SOFC electrolyte needs to exhibit high protonic conductivity with a very low electronic contribution while the dense ceramic material of a H2-separating membrane has to be a mixed protonic-electronic conductor.
Proton-conducting fuel cells differ from traditional solid oxide fuels in that the electrolyte membrane transports protons instead of oxygen ions. These fuel cells can operate at lower temperatures owing to the higher mobility of protons in comparison to oxygen ions while proton conductivity decreases at high temperatures (above 800 °C) owing to the lower proton (water) solubility in the oxide, that is a lower concentration of proton charge carriers. The high proton conductivity at intermediate temperatures helps to alleviate problems caused by the high operating temperatures of traditional solid oxide fuel cells.
Starting with the anode side of the cell, hydrogen is oxidized to protons and is then conducted through the electrolyte until it finally combines with oxygen to produce water at the cathode. The electrons from the anode partial reaction are routed through an external circuit and thus electricity is produced and the circuit is closed by returning the electrons to the cathode. The protons must move through the electrolyte, therefore, a membrane material with dominant protonic conductivity at elevated temperatures and in a wide range of is required as the PC-SOFC electrolyte (minor electronic conductivity contributions are tolerable). High redox stability accompanied by a high mechanical and thermal stability is also required in terms of durability and optimal operation.
Dense ceramic membranes for H2 separation can be implemented in the pre-combustion CO2 capture concept, offering an efficient way to produce high-purity hydrogen for a variety of applications, along with CO2 sequestration, which can be identified as a major benefit of great environmental relevance. The pre-combustion approach consists of several stages. Stage I is the partial oxidation of fossil fuel (methane, natural gas or gasified coal) with pure O2 resulting in a synthesis gas (CO + H2); Stage II is the water gas shift reaction (WGSR) to transform the carbon monoxide from the synthesis gas to carbon dioxide; Stage III includes CO2/H2 separation by means of a mixed protonic–electronic conducting (MPEC) membrane. Typical operating conditions in Stage III of a pre-combustion power plant are temperatures in the range of 400–900 °C, pressures between 2 and 50 bar, as well as an environment containing CO2, H2, H2S (about 5–200 ppm) and H2O. Hence, a membrane material is required with mixed protonic–electronic conductivity and pronounced chemical stability under reducing and hydrothermal conditions.
The operation principles of a PC-SOFC and a membrane for H2 separation are presented schematically in Fig. 17.1(a) and (b) respectively.
Hydrogen can be produced industrially from the electrolysis of water but at a high energy cost. High-temperature electrolysis offers higher power efficiency and hence significant cost savings over conventional low-temperature electrolysers. There have been relatively few investigations of ceramic electrolysers so far and most work has been done on low-temperature alkaline and proton exchange membrane (PEM) electrolysers which are considered to be state-of-the-art technology in electrolysis.
Elevated temperatures are required in order to optimize the cost effectiveness of the electrolysis cell. Hence ceramic materials are interesting candidates for high-temperature cells operating at high to intermediate temperatures. The following reactions take place in an electrolysis cell. At the anode water is split according to:
High-temperature proton conductors find several applications as gas sensors. They can be applied principally for the quantification of H2 (and CO) and humidity while gas sensing can be performed not only in hot gas streams but also in molten salts or metals. The advantages of the high-temperature proton conducting oxide sensors entail sensitivity, accuracy, a wide H2 concentration detection range, robust construction and possible miniaturization. Typically, in pure proton conductors, the EMF voltage depends on the hydrogen concentration gradient across the electrolyte, while the electrolyte conductivity varies as a function of the water concentration in the electrolyte gas environment (Wang and Virkar, 2004), since the hydration degree determines the concentration of charge carriers. These two phenomena, together with the kinetic effects of surface reactions when the electrolyte surface is covered with specific sensing electrode materials (electrocatalysts), allow the development of sensors for H2, CO (Yamazoe and Miura, 1988) and water for very particular applications. Several proton conducting materials hava been tested in a variety of H2 sensors, for example Yb-doped SrCeO3−δ(Kosacki and Arderson, 1998), Gd-doped BaCe3−δ (Bonanos et al., 1991), BaZr0.4Ce0.4In0.2O3−δ (Du and Nowick, 1996) and KCa2Nb3O10 (Sakthivel and Weppner, 2007). Most of these electrolyte materials were used as potentiometric sensors (monitoring the EMF response) albeit depending on the application and sensor arrangement, they can be used as conductimetric or amperometric sensors. Further details of sensor configuration, performance and sensing electrodes can be found in Korotcenkow et al. (2009). Probably, the most important application of high-temperature solid electrolyte H2 sensors is the quantification of hydrogen in molten metals, for example aluminium or copper, in the metallurgical industry (Fukatsu and Kurita, 2005). Specifically, high-quality aluminium alloy manufacture requires strict control of the solved hydrogen in the molten metal during the casting process. The best sensor for this application makes uses of a galvanic cell (Yajoma et al., 1993) based on Pt || CaZr0.9In0.1O3−δ || Pt and allows precise measurement of the hydrogen concentration in the temperature range 700–800 °C.
In summary, the applications considered above define specific demands on the candidate ceramic proton-conducting materials, for example in terms of material transport properties, stability under relevant operating conditions, and so on. In terms of their electrical properties, materials with dominant pure proton conductivity are mainly suitable for application as dense solid electrolytes for PC-SOFC, while materials with mixed proton–electronic conductivity are applicable as H2-separating membranes. Along with the required pure protonic or mixed proton–electronic conductivity, candidate materials for these applications need to exhibit a high chemical, thermal and mechanical stability under a variety of operating conditions (e.g. redox when material is applied as an electrolyte in PC-SOFC or reducing when operating as a H2-separating membrane). These strict requirements greatly limit the range of candidate materials that can operate efficiently under the relevant conditions, maintaining their functionality and durability in the long term without significant performance degradation.
In atmospheres containing water (water vapour) or hydrogen, protons are dissolved in the oxide lattice thus forming positively charged defects according to the following quasi-chemical reaction written in Kröger–Vink notation:
where is an oxygen vacancy with an effective charge (2 +), is an oxygen atom on its ordinary place (subscript O) in the crystal lattice with a neutral charge (expressed by superscript X), is a proton defect effectively charged (1 +), e′ is an electron with charge (1–), KOH · is the equilibrium constant of proton defect formation reaction, is the concentration of the proton defect, is the concentration of the oxygen vacancy, is the concentration of oxygen atoms on their ordinary place in the crystal lattice and is the partial pressure of water vapour.
According to Kreuer (2003), for large band gap oxide materials (e.g. Ce/Zr/Ti-based perovskites), proton defects are formed at moderate temperatures through a dissociative absorption of water that requires the presence of oxygen ion vacancies (Norby, 1999). Water from the gas phase dissociates to hydroxide ion and a proton. The hydroxide ion is adsorbed by the oxygen vacancies and the proton forms a covalent bond with the lattice oxygen. Thus two positive proton defects, are generated.
The formation of proton defects is accompanied by a significant weight increase (often determined via thermogravimetry). The concentration of proton defects changes with temperature and water partial pressure thus giving information about the thermodynamic stability of proton defects. The equilibrium of the hydration reaction Equation [17.1] shifts to the left with increasing temperatures since the reaction is exothermic: the more negative the hydration enthalpy, the higher the proton content and the dehydration temperature. Several factors strongly affect the stability and mobility of proton defects, for example crystal symmetry, chemical nature (acid–base sites), electronegativities of cations that interact with the lattice oxygen and repulsive forces between the proton and its surroundings (Kreuer, 1996, 1997, 1999, 2000, 2003, 2004; Ricote et al., 2009).
When a hydrogen atom is absorbed into a structure it usually loses its electron and turns into a proton. The charge densities of protons are too large to exist as separate entities in the structure and they tend to associate with electronegative atoms, such as oxygen ions, in the parent structure.
Proton diffusion mainly occurs by two mechanisms: the free proton mechanism (Grotthuss mechanism) and the vehicle-facilitated mechanism (Kreuer, 1996). The Grotthuss mechanism consists of proton jumps from a stationary oxygen atom to a neighbouring oxygen atom after it has rotated to the correct position between the two oxygen atoms.
In the vehicle mechanism, protons attach to an oxygen ion and move together with this ‘vehicle’ (via translational diffusion). The hydroxide ions thus formed diffuse through the structure, much like ordinary oxygen ions. Hydroxide ions may possibly have a lower activation energy of diffusion than oxygen ions owing to their smaller size. Hydroxide ions can diffuse through the structure using vacancy diffusion or interstitial diffusion. The electroneutrality balance is maintained by the counterdiffusion of unprotonated vehicles or most probably by oxygen vacancies.
According to Kreuer, for structural reasons, the Grotthuss-type mechanism is more likely to be found in oxides than the migration of hydroxyl ions. For perovskite structures it is considered that the proton transfer between two neighbouring lattice oxygens is the rate-limiting step. With increasing temperature a gradual transition from the Grotthuss to the vehicle type of mechanism (translational diffusion) can be expected owing to elongation and breaking of the hydrogen bonds that may suppress proton transfer and release the translational degree of freedom.
and is the mobility of protons, T is the temperature in K, ΔHmob is the enthalpy of mobility, k is the Boltzmann constant, is the proton conductivity, z is the electrical charge, F is the Faraday constant, is the concentration of protons and is the proton diffusion coefficient.
The mobility of proton defects and proton conductivity has an Arrhenius-type temperature dependence, meaning that they depend exponentially on the temperature. Proton concentration is also dependent on the temperature in terms of its thermodynamic stability when the protons are incorporated into the crystal lattice (as discussed above). The total activation energy ∆Eact for proton conduction consists of energetic terms required mainly for proton defect formation and proton migration.
In the case of dense ceramic ionic conductors operating at elevated temperatures, lattice diffusion of point defects (vacancies, interstitial defects, protons) is the dominating mechanism for the transport of charge carriers through the crystal lattice. Point defects may be formed intrinsically by varying the ratios of the basic components of the material or extrinsically by acceptor doping. In this sense, doping plays a crucial role in creating defects in the primary crystal structure. Hence, doping is a very frequently used strategy to modify and control the properties of the material by (i) distorting the crystal lattice (geometrical factor, e.g. ionic radius), (ii) influencing the chemical properties by introducing elements with a different chemical nature (electronegativity, chemical properties), or/and (iii) modifying the electrical properties by controlling the concentration and the type of electro-balancing charge carriers that are formed within the structure as a result of the doping.
In a very common case when the acceptor doping route is followed, positively charged species are generated to maintain the electroneutrality of the lattice. Positively charged species are electron holes (), oxygen. Vacancies () and proton defects ()In accordance with the mass action law, under definite conditions (, , and T) the existing defects are in dynamic equilibrium. For each set of equilibrium conditions, there are values of , , and T that determine the domination of a definite pair of defects (represented by Brouwer diagrams). Under dry oxidizing conditions, oxygen is built up in the crystal lattice, so that the concentration of the oxygen vacancies diminishes simultaneously with the generation of electron holes which are needed to neutralize the negative charge generated by the aliovalent doping element.
In water- and hydrogen-containing environments proton defects in the oxide contribute considerably to the conductivity and have to be taken into account with respect to defect chemistry and defect equilibrium.
In order to obtain considerable electron conduction levels within the material, a host lattice element is replaced by a doping element M with a variable oxidation state. As an example, the following quasi-chemical reactions describe the defect formation under oxidizing [17.11] and reducing [17.12] conditions, respectively:
where M2O3 is the oxide compound of doping element M where M has a charge (3 +), MO2 is the oxide compound of doping element M where M has a charge (4 +) and AO2 is the host oxide lattice formed by host lattice element A and oxygen.
And finally, independently of the kind of the doping scheme used, when thermodynamic and chemical equilibrium are maintained, the equality of all positive and negative species generated within the crystal lattice is also maintained.
Proton conductivity in the perovskite type of oxide ceramics was first discovered by Iwahara et al. (1981) in the early 1980s. The most extensively studied high-temperature proton-conducting ceramic materials with a perovskite structure are SrCeO3, BaCeO3, SrZrO3, as well as complex perovskites of types A2B′B″O6 and A3B′B″O9. Materials with a perovskite structure having a general formula A2 +B4 +O3 (type II–IV), where when A is Ba and B is Ce, Zr, Tb or Th they exhibit the best proton conductivities of above 10− 2 S cm− 1 (Davies et al., 1999; Fontaine et al., 2008; Iwahara, et al., 1981, 1993; Kreuer, 1996, 1999, 2003; Kreuer et al., 2001; Krug et al., 1995; Meulenberg et al., 2010; Norby, 2009; Schober et al., 1995; Schober and Bohn, 2000; Stotz and Wagner, 1966; Yamanaka et al., 2005). They have large negative hydration enthalpies and the lowest activation energies for proton conduction. As summarized by Norby (2009), when the A-site is occupied by Sr instead of Ba, the conductivity drops to between 10− 2 and 10− 3 S cm− 1 (SrZrO3, SrCeO3). Calcium-based perovskites are also proton conductors, especially CaZrO3. Within perovskite type I-V (A1 +B5 +O3) there are some examples such as KTaO3 for which a very low conductivity was reported (Lee et al., 1986).
Cerates and zirconates have been extensively studied for hydrogen separation membranes. Selected examples of doped cerates and zirconates are summarized in Table 17.1. The H2 flux data published in the literature about cerates point to Yb-doped SrCeO3 as the best state-of-the-art perovskite-based material for H2 separation applications. However, there are several disadvantages that limit the application range of these materials. Briefly, chemical stability under reducing atmospheres especially concerns the cerates. Because of their basicity they react easily with CO2 and H2S/SO2/ SO3 at moderate and with H2O at low temperatures and form alkaline earth carbonates, sulphates and hydroxides, respectively (Scholten et al., 1993; Shirsat et al., 2004). Consequently, their mechanical properties become poor owing to formation of reaction products at the grain boundaries.
Zirconate-based materials were found to be attractive for membrane applications mainly because of their higher stability of zirconates under reducing atmospheres, in particular in CO2 environments. However, existing experimental data about the hydrogen flux through zirconate-based membranes are rather limited in comparison to the data available for the cerate class of membranes. Despite their better stability in reducing environments, zirconates exhibit high grain boundary resistance that negatively influences their electrical performance. Furthermore, manufacturing of membranes with high gas tightness requires extremely high sintering temperatures (e.g. ~ 1700 °C or even higher for zirconates). In an attempt to lower grain boundary resistance and sintering temperatures, much work has been done in recent years. Examining the influence of different dopants (Naoki et al., 2008; Tao and Irvine, 2007) or manufacturing routes (Duval et al., 2007; Gibson et al., 2009; Guillaume et al., 2005; Meulenberg et al., 2006; Serra and Meulenberg, 2007; Serra et al., 2007) does not lead to a significant improvement in the electrical properties but in some cases results in better sinterability of the investigated zirconates.
Norby and Kofstad found that Y2O3 conducts protons although the conductivity values are low (Norby and Kofstad, 1984, 1986). Furthermore, proton conduction in numerous undoped and doped rare earth oxides of type Ln2O3 and LnBO3 (Ln is lanthanide element) have been investigated and it was found that they also exhibit proton conduction at elevated temperatures (Larring and Norby, 1995; Norby, 1990; Norby et al., 1992; Norby and Larring, 1997). An extensive study of the equilibrium between water vapour, protons and oxygen vacancies in rare earth oxides Nd2O3 and Sm2O3 doped on the A-site with 1 and 2 mol% Ca was conducted by Larring and Norby,1997). It was found that these rare earth oxides are mixed proton, ion and p-type electron conductors depending on the conditions. They are nearly pure proton conductors at moderate temperatures and under wet conditions.
The phosphates (e.g. zirconium phosphates and uranylphosphates) have been widely studied as proton conductors at moderate temperatures. Proton conductivity at higher temperatures was found relatively recently in materials that have no structural protons such as Li3PO4, other alkali phosphates (Mellander and Zhu, 1993) and β-Ca(PO3)2 (Greenblatt et al., 1990). LaPO4 was suggested as an interesting candidate owing to its stability and ability to dissolve and transport protons. Ca- and Sr-substituted LnPO4 ortho-phosphates, where Ln is La, Nd, and Gd were studied by Norby and Christiansen (1995). The conductivity of Ca- and Sr-substituted LaPO4 was found to be mainly of the proton type except at the highest temperatures and driest conditions, where the native conductivity dominates.
Pyrochlore-structured La2Zr2O7 was reported to show proton conductivity which was not found in the isostructural Ti-containing oxide (Labrincha et al., 1997). Shimura et al. (1996) found that Y2Ti1.8 M0.2O7−δ with M = In on Mg did not show measurable proton transport. The conductivity of undoped La2Ti2O7 was previously measured at relatively high temperatures (800–1000 °C) and with no special emphasis on possible effects of protons and proton transport (Balachandran and Eror, 1982). Haugsrud and Norby (2005) studied 2% Ca-doped La2Ti2O7 and proton conductivity was determined to be ~5 × 10− 5 S cm− 1 at 800 °C. The hydrogen flux across Ca-doped La2Ti2O7 film with a thickness of 10 μm was estimated on the basis of measurement of partial conductivities. It was estimated that such a membrane would yield a hydrogen flux in the order of 0.1 mlN cm− 2 min− 1 at 750 °C, where subscript N stands for normalized flux, this value being lower by a factor of ~5 compared to that of 5% Yb-doped SrCeO3 under the same conditions. However, the protonic conductivity of rare earth oxides, phosphates and pyrochlores is far too low for these materials to be implemented and efficiently operated in the envisaged applications.
Dense H2-conducting ceramic materials achieve much lower values of H2 flux compared to dense metallic membranes (based on Pd alloys) and microporous silica membranes, which display much better performance. However, these membranes face serious degradation problems when integrated in processes of coal gasification and/or steam reforming. Moreover, these two material classes (dense metallic and ceramic microporous materials) are not the subject of this paper and they will not be considered here.
Rare earth ortho-niobates (LnNbO4) and tantalates (LnTaO4) are materials with a fergusonite structure (monoclinic symmetry) that changes to a scheelite structure (tetragonal symmetry). The temperature of phase transition is different for these two classes (see Section 17.5.3). Proton conductivity was found in several acceptor-doped rare earth ortho-niobates and ortho-tantalates (LnNbO4 and LnTaO4) at elevated temperatures and in a humid atmosphere (Haugsrud and Norby, 2006). Tantalates exhibit a proton conductivity approximately one order of magnitude lower than that of niobates (Norby and Haugsrud, 2007). The total conductivity of undoped LaNbO4 was found to be about 4 − 5 × 10 − 5 S cm − 1 at 1000 °C (Mokkelbost et al., 2008) compared to the A-site-doped (Haugsrud and Norby, 2006; Mokkelbost et al., 2008) material La0.99Ca0.01NbO4, which exhibits conductivity of less than 2 × 10− 3 S cm− 1 in a wet O2 atmosphere and conductivity of about 7 × 10− 4 S cm− 1 in a wet H2 environment, both at 1000 °C. At 800 °C the total conductivity of La0.99Ca0.01NbO4 in wet O2 is ~ 7.5 − 8 × 10− 4 S · cm− 1 and less than 6 × 10− 4 S cm− 1 in wet H2 (Haugsrud and Norby, 2006), while the conductivity of undoped LaNbO4 at the same temperature is ~ 8 × 10− 6 S × cm− 1.
A variety of studies on rare earth ortho-tantalates with a general formula LnTaO4 have shown that these materials exhibit proton conductivity at lower temperatures, but this conductivity is generally an order of magnitude lower than that of niobates (Norby and Haugsrud, 2007). At 1000 °C La0.99Ca0.01TaO4, for example, exhibits a total conductivity of about 4 × 10− 4 S cm− 1 and a protonic conductivity of ~ 2 × 10− 4 S cm− 1. At the same temperature, the values of the total and protonic conductivities for Nd0.99Ca0.01TaO4 are σtot ~ 2 × 10− 4 S cm− 1 and σH < 2 × 10− 4 S cm− 1; for Gd0.99Ca0.01TaO4 are σtot ~ 8 × 10− 5 S cm− 1 and σH < 5 × 10− 5 S cm− 1, and for Er0.99Ca0.01TaO4 are σtot ~ σH = 4 × 10− 5 S cm− 1, respectively. With decreasing ionic radius in the row La → Nd → Gd → Er, the activation energy for proton conduction increases and the same tendency is true both for ortho-niobates and ortho-tantalates.
Rare earth tungstates (Ln6WO12) are materials with a defective fluorite structure, although the crystal symmetry is not completely clarified. A correlation was found between the type of the lanthanide element and the crystal lattice symmetry that shows a tendency toward reducing the crystal symmetry from cubic through tetragonal and finally to rhombohedral when Ln changes in the order La → Nd → Er.
Shimura et al. (2001) studied La5.8WO11.7 in electrochemical hydrogen pumping experiments and proton conduction for this material was demonstrated. The proton transport number was determined to be between 0.7 and 0.9 in a H2-containing atmosphere. A proton conductivity of 5 × 10− 3 S cm− 1 was measured in wet H2 at 900 °C. Haugsrud (Haugsrud, 2007) studied defects and transport properties in Ln6WO12 where Ln is La, Nd, Gd and Er. A general feature of these materials is that they exhibit protonic conductivity in wet atmospheres, below ~ 900 °C. Under oxidizing and reducing conditions and at elevated temperatures, Ln6WO12 becomes predominantly an electronic semiconductor. Maximal conductivity was observed for undoped La6WO12 compared to La6WO12 doped with 0.05 mol% Ca on the A-site.
In several publications, the feasibility of tungstates as membranes for H2 separation was studied by flux measurements or estimated theoretically (Escolástico et al., 2009; Haugsrud and Kjølseth, 2008; Norby and Haugsrud, 2007; Shimura et al., 2001). Haugsrud and Kjølselh (2008) compared the ambipolar conductivity of undoped La6WO12 to that of 5%-Yb-doped SrCeO3 (state-of-the-art material) and came to the conclusion that the undoped La6WO12 is a good candidate for membrane applications. A permeability data prediction was communicated by Norby et al. (Norby and Haugsrud, 2007) for La6WO12, SrCeO3 and Er6WO12. The flux was calculated for 10-μm-thick layers with a feedside pressure of 10 atm H2 and assuming that the flux was ruled by the Wagner equation, that is via a bulk diffusion mechanism. The predicted hydrogen permeation for La6WO12 at 800 °C has a value of about 2.0 mlN cm− 2 min− 1, which is higher compared to SrCeO3 and Er6WO12 yielding a hydrogen flux at the same temperature of about 1.0 mln cm− 2 min− 1 and less than 0.1 mln cm− 2 min− 1, respectively.
Table 17.1 presents a summary of selected examples of ceramic materials developed and tested or theoretically evaluated for H2-separating membrane applications.
The cubic perovskite structure is observed in many compounds featuring a composition ABX3, whereby A and B are cations with different oxidation numbers and X is an anion. The A-cation occupies the centre of the unit cell, while the B cation and the X anions are arranged at the corners and the edges of the unit cell, respectively. Figure 17.2 shows the unit cell of CaTiO3 as an example of the perovskite structure. Owing to steric constraints, caused by different combinations of ion radii, the unit cell of most perovskites is slightly distorted. A perfect cubic unit cell is only present within SrTiO3.
Proton-conducting compounds with a perovskite structure are, for example, the acceptor-doped lanthanum rare earths LaErO3 (Larring and Norby, 1994) and LaScO3 (Fujii et al., 1998). In so-called complex perovskites, significant non-stoichiometries are used in order to enhance the proton conductivity. A prominent example with a high cation non-stoichiometry is Ba3Ca1.18Nb1.82O8.73 (BCN18) (Norby, 1999). Hydration of proton-conducting perovskites with high oxygen deficiency allows for increased proton contents. Ba2InSnO5.5 and Ba2In2O5, for example, may take up 0.5 protons (Schober, 1998) and one proton (Fischer et al., 1999) per unit cell, respectively.
A number of titanates featuring the composition A2Ti2O7 with A = Y, and the elements in the row from Sm to Yb as well as some zirconates, for example La2Zr2O7 and Gd2Zr2O7, crystallize in the pyrochlore structure. The classical pyrochlore composition is (Na,Ca)2(Nb,Ti)2(O,F)6(O,F) [= A2B2X6Y] and is sometimes referred to in the literature as anion-deficient fluorite derivative (Jona et al., 1955). Basic building components are arrays of corner-connected tetrahedra formed by the A atoms while the Y atoms occupy the centres of the tetrahedra. Each A atom is coordinated by six X atoms arranged in a puckered hexagon and a pair of Y atoms normal to the mean plane of the hexagon, resulting in hexagonal–bipyramid coordination (Nyman and Andersson, 1978). Figure 17.3 presents the unit cell of Gd2Ti2O7 as an illustration of the pyrochlore structure.
As described in Section 17.4.5, rare earth ortho-niobates (LnNbO4) [= ABO4] and rare earth ortho-tantalates (LnTaO4) possess proton conductivity. In these material classes a phase transition takes place from a monoclinic crystal structure (fergusonite type) at low temperatures to a tetragonal crystal structure (scheelite type) at high temperatures. In the case of the niobates, the transition temperature ranges from about 500–800 °C and increases with decreasing radii of the rare earth cations (Gingerich and Bair, 1964). The transition temperature of the tantalates is significantly higher and ranges from 1300–1450 °C (Stubican, 1964). The scheelite structure can be described as two intercalated diamond lattices, one for A cations and the other for B cations. The A cations are coordinated by eight oxygen anions and the B cations are tetrahedrally coordinated by the oxygen anions. Figure 17.4 shows the unit cell of the mineral scheelite (CaWO4). The fergusonite structure can be considered as a distorted scheelite structure. Each A cation is surrounded by eight oxygen anions and each B cation is surrounded by four oxygen sites and two additional near oxygen sites. As reported by Errandonea et al. (2004), the scheelite-to-fergusonite transition can be described as a shear transformation in which the scheelite structure is partially conserved while certain sheets are slightly shifted.
Rare earth tungstates with a stoichiometry close to Ln6WO12 (including Y6WO12) are interesting material candidates in terms of proton conductivity (e.g. Haugsrud, 2007). The crystal structures of most materials of this class, however, have not yet been completely solved. Detailed structure reports are only given in the case of Y6WO12 and Ho6WO12 (Diot et al., 2000). Diot et al. report that the rhombohedral structure of these compounds can be described by an arrangement of seven edge-sharing cubes, where the central cube is occupied by W cations.
The crystal symmetries of the lanthanide tungstates were published by McCarthy et al. (1972) in 1972. The crystal symmetry depends on the rare earth element and is cubic in the case of La, Ce and Pr, pseudo-tetragonal from Nd to Gd, and is rhombohedral from Tb to Lu and for Y. The structure of the rare earth tungstates is related to the fluorite structure and is therefore frequently referred to in the literature as an ordered defective fluorite structure or as a disordered pyrochlore structure (e.g. Diot et al., 2001). Magrasó et al. (2009) reported that the structure of La6WO12 is closely related to the cubic fluorite structure and is characterized by an ordering of the La and the W atoms, in other words a superstructure is present. This finding is also confirmed by transmission electron microscopy (TEM). Figure 17.5 shows a high-resolution TEM micrograph of La6WO12 in the  orientation. Atom columns along this direction consisting of La atoms and W atoms appear as white spots on a dark background. Owing to the contrast-forming conditions applied, oxygen sites are not visible within this image. Because of their higher electron-scattering power, W atoms appear with higher intensity. The superstructure of the cations is clearly visible as an ordering of the brighter and darker spots.
In a contrast to the structures considered above, the phosphates and rare earth oxides are not classified by their structure but by their elements. This makes an attempt to summarize and describe the existing great variety of different structures amongst these material classes a rather difficult task even when common structural features are applied. Therefore, presenting a structural overview of these groups of materials in this general structural section of the chapter was not found to be relevant nor of benefit to the reader.
Perovskitic proton-conducting materials are principally based on the type AIIBIVO3. As a result, these materials contain large proportions of alkaline-earth elements, basically Sr and Ba. However, as previously mentioned, the high basicity of these elements results in the high stability of the corresponding Ba-Sr carbonates and sulphates (with respect to the mixed oxide compound) and, therefore, these proton conductors tend to react with acid gases (CO2 and S-containing gases) under the envisaged operating conditions and decompose irreversibly. The stability of the perovskite compound depends on the specific lattice composition and the kind of transition metal present in the B-atom position is particularly crucial. The order for the three main metals used as B-atoms is (in decreasing order of stability): titanates, zirconates and cerates (Kreuer et al., 2001). In general, it seems that there is an opposing trend between mixed oxide stability and proton conductivity, that is the higher the stability, the lower the proton conductivity. Doped SrCeO3 and BaCeO3 are thermodynamically stabilized slightly with regard to decomposition into the single oxides while they react rapidly with CO2 even at low concentrations. This fact prevents their use in realistic applications entailing the conversion of hydrocarbons or syngas. Nevertheless, the choice of the acceptor dopant would allow stability to be increased and a balanced stability to be achieved between perovskite oxide and protonic defects principally caused by modification of local symmetry.
A realistic stability test was carried out using powdered membrane materials based on several doped BaZrO3 and BaCeO3. The test period was three days and the conditions were 800 °C and 700 °C (independent tests) respectively under wet CO2-rich reducing conditions, that is CO2 + CH4 gas (1/9 molar ratio) at atmospheric pressure (Serra and Escolástico, unpublished results). All doped zirconates remained stable during the treatment as revealed by X-ray diffraction (XRD) diffraction while all doped cerates reacted dramatically to form carbonates and CeO2-based oxides.
Another stability test comprised the treatment of a similar set of samples under wet CO2-rich reducing and corrosive conditions, that is in gas flow 115 ppm H2S, 4.43% CO2, 2.12% CO and 92.09% H2 at 30 bar for three days under continuous flow of 50 ml min− 1. Again, the zirconates successfully passed the test while the cerate decomposed to form carbonates and sulphates.
Ln6WO12 membrane materials are believed to possess high chemical stability to carbonation owing to the absence of highly basic cations in their structure. The most H2 permeable material reported up to now consists uniquely of undoped La6WO12. Similar tests to those reported above confirmed the stability to carbonation and in the presence of significant amounts of H2S, as concluded from XRD investigations. Nevertheless, a long-term stability test is needed to assess the applicability of these materials. Moreover, at high temperatures, Ln6WO12 shows high reactivity with other ceramic compounds such as YSZ, CeO2 and NiO. Figure 17.6 illustrates the high stability of Ln6WO12 (where Ln is Nd) in a wet H2 atmosphere at elevated temperatures. A similar situation was observed for the compounds LaNbO4 and LaTaO4, which remained stable after treatment in CO2. However, it was observed that some specific multidoped LaNbO4 compounds decomposed under wet CO2-rich reducing and corrosive conditions.
Inorganic proton conducting membranes can be implemented in a variety of applications and therefore the ceramic materials have to meet the corresponding requirements. During the past 20 years significant progress in research and development (R&D) has been made, mostly motivated by the benefits that these membranes offer in a number of applications. Improved proton conductivity and stability under relevant operating conditions are of a particular importance. Proton conductivity is rather sensitive to temperature owing to the thermodynamic stability of proton defects in terms of proton content that contributes to the total conductivity of the material. Moreover, the stability and mobility of proton defects is unique for each kind of material. It strongly depends on the crystal structure, chemical nature and concentration of doping elements. However, as seen in the literature, the stability of protons is not an indication of either the chemical stability of the material or high proton conductivity. In this context, it should be mentioned that proton defects have the highest stability in cerates followed by the zirconates but in terms of chemical stability zirconates are more stable to carbonation than cerates. In order to illuminate the basics of these phenomena and complex dependencies for different classes of materials, intensive research is strongly required. In any case, proton–electron or dominant proton conductivity, the stability/mobility of proton defects and chemical stability must be balanced by a trade-off for obtaining the optimal properties of the membrane.
Strategies for the development of novel, application robust membranes rely on progress in materials science-based improvements as well as technology-driven optimization of the process and application conditions. R&D activities are consequently directed towards mechanically robust porous structures and functional layers in the nanometre range and the development of mixed-conducting oxides by means of theoretical approaches like atom modelling and experimental screening, based both on prior knowledge and modelling outcome. Decisive results based on a parallel investigation of (i) new materials, design and processing of components and equipment, (ii) integration into power plants and the related process engineering and (iii) energy systems analysis will pave the way to efficient and affordable membrane technologies including carbon capture and storage and hydrogen generation from carbon-based fuels by means of new gasification processes for fossil and biomass energy carriers.
The authors acknowledge the Initiative and Networking Fund of the Helmholtz Association, contract HA-104 (‘MEM-BRAIN’), German Federal Ministry of Education and Research (BMBF) through the Northern European Innovative Energy Research Project, contract 03SF0330 (‘N-INNER’) and Spanish Ministry for Science and Innovation (Project ENE2008-06302).
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