On the development of proton conducting materials for technological applications

Solid State Ionics 97 (1997) 1–15

On the development of proton conducting materials for technological applications K.D. Kreuer* ¨ Festkorperforschung ¨ Max-Planck-Institut f ur , Heisenbergstr. 1, D-70569 Stuttgart, Germany

Abstract Some aspects of the simultaneous optimisation of material properties of proton conductors which are relevant for their use in electrochemical cells such as fuel cells, electrochemical reactors and sensors (high proton conductivity, chemical, electrochemical and morphological stability) are discussed. Suggestions are made for the further development of proton conducting perovskite type oxides, proton conducting polymer membranes and medium temperature proton conducting materials. Keywords: Oxide; Polyaromatic membrane; Nafion; Imidazole; SOFC; PEM-fuel cell; Conduction mechanism; Stability; Perovskite Materials: Ba 2 YSnO 5.5 ; Ba 3 CaNbO 9 ; Polyetherketone; Imidazole; Pyrazole; Perovskite

1. Introduction Although high proton conductivity has been reported for a large number of solid compounds and materials [1] and many of them have been suggested for applications as electrolytes in electrochemical cells, their use in technological devices is still quite limited. Apart from the fact that proton conductivity in the solid state remains significantly below the upper limit for proton conductivity in liquids, major problems arise from the numerous additional material requirements, other than proton conductivity, which have to be met for a particular application. Therefore, the simultaneous optimisation of all relevant material properties appears to be an interest*Fax: 149 711 689 1722; e-mail: [email protected] 0167-2738 / 97 / $17.00  1997 Elsevier Science B.V. All rights reserved PII S0167-2738( 97 )00082-9

ing route in the development of new proton conducting materials for specific applications. Such a strategy is promising especially when the mutual dependences of the relevant properties can be expressed in terms of the same structural, thermodynamical and dynamical parameters. Along this line, this paper presents a few issues in the selection or development of: (i) high temperature proton conducting perovskite type oxides, (ii) low temperature proton conducting hydrated polymer membranes, (iii) intermediate temperature proton conducting materials relying on proton transfer between nitrogen acting as proton donor as well as proton acceptor. The first two have attracted increasing interest in recent years especially because of their potential use as separator materials in fuel cells, electrochemical

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reactors and sensors. Whereas the existence of proton conductivity in these families of compounds has been known for a long time, the presented approach also leads to materials based on distinctly different chemistries. This may even open up new fields of application as demonstrated for the third case.

2. Perovskite type oxides The systematic investigation of proton conducting perovskite type oxides (ABO 3 ) started with the work of Takahashi and Iwahara in 1980 [2]. It was already this paper which reported proton conductivity in indium doped zirconates (In:SrZrO 3 ), the only class of oxides, which has later been developed for a technological application, i.e. as a separator material for high temperature potentiometric hydrogen sensors [3]. In contrast to this application, the use of oxides in fuel cells, electrochemical reactors or amperometric sensors requires much higher proton conductivities with high transference number and it appears that this is hard to achieve while maintaining good chemical, electrochemical and mechanical stability. In the following, some important parameters, which determine the above mentioned properties, are identified and suggestions are made for the further development of such materials.

2.1. Charge carrier concentration Protonic defects in oxides correspond to hydroxide ions on oxide ion sites (OH ?O ) which are positively charged defects competing with other positively charged defects, i.e. electronic holes (h ? ) and oxygen ion vacancies (V ??O ) charge compensating negatively charged defects, mostly acceptor dopants such as R B (where R is frequently a rare earth cation). The successive formation and annihilation of these positively charged defects may be written as: H 2 O 1 V O?? 1 O O3 5 2OH O? , K 5 (K1 K2 )21 ,

(1)

2OH ?O 1 1 / 2 O 2 5 H 2 O 1 2h ? 1 2O O3 , K1 ,

(2)

?? 2h ? 1 O 3 K2 , O 5V O 1 1 / 2 O 2 ,

(3)

where the formation reaction of protonic defects [ Eq. (1)] is a linear combination of Eqs. (2) and (3). Because of the net creation of gas molecules in Eqs. (2) and (3) they are expected to have a positive reaction entropy, i.e. they shift to the product side with increasing temperature. With increasing temperature one therefore expects the appearance of positively charged defects in the order OH ?O →h ? →V ??O . At which temperatures this happens depends on the dopant concentration, the external parameters pH 2 O and pO 2 and the equilibrium constants of all involved reactions. They additionally comprise the intrinsic formation of electronic defects: h ? 1 e9 5 nil, K3 .

(4)

The equilibrium conditions for Eqs. (2)–(4) and the electroneutrality condition: 2[OH ?O ] 1 2[h ? ] 1 [V O?? ] 2 2[R 9B ] 5 0

(5)

then provide the four independent equations defining the concentration of all four defects (OH ?O , V ??O , h ? , e9) as a function of K1 (T ), K2 (T ), K3 (T ) and pH 2 O and pO 2 where the acceptor dopant concentration [R 9B ] is taken as fixed. For each set of equilibrium constants K1 2K3 there is a range of partial pressures (high pH 2 O , low pO 2 ) where protonic defects are directly replaced by oxygen ion vacancies as majority charge carriers with increasing temperature and a partial pressure range where electronic holes become apparent in between. For a high concentration of protonic defects with respect to that of oxygen ion vacancies and electronic holes, the reaction constants K1 , K2 and K3 must be low and the product K1 K2 also low so that K becomes high compared to K1 and K2 . Let us therefore first discuss the parameters determining the equilibrium constant for the formation of protonic defects directly from oxygen ion vacancies by the dissociative absorption of water K. This is a pure acid / base reaction not involving electronic defects, which will be addressed at the end of this section. Different contributions to the Gibbs free energy of reaction become apparent by formally separating Eq. (1) into the following partial reactions [4]: 22 (H 2 O) g 5 2H 1 g 1 Og ,

(6)

K.D. Kreuer / Solid State Ionics 97 (1997) 1 – 15 ?? 3 O 22 g 1 VO 5 OO ,

(7)

3 ? 2H 1 g 1 2O O 5 2OH O .

(8)

As opposed to the first partial Eq. (6) the two latter Eqs. (7) and (8) depend on the kind of oxide. In Eq. (7) the formation enthalpy for oxygen ion vacancies is gained back. For binary oxides this approximately scales with the molar density as a measure for the lattice energy [5] reduced by the energy associated with the lattice relaxation around oxygen ion vacancies [4]. For more complex oxides (e.g. perovskites of the type ABO 3 ) the thermodynamics of the lattice and therefore also for the formation of oxygen ion vacancies is very sensitive to the compatibility of the different species with the particular structure. In a rough approximation simple empirical parameters, such as the Goldschmidt tolerance factor for perovskites (e.g. [6]), may account for this. Defect relaxation energies depend not only on the particular compound but also on the oxygen ion vacancy concentration. For BaCeO 3 based perovskites, e.g., this decreases from approximately 0.3 eV for low defect concentrations to approximately 0.2 eV for high defect concentration [4]. Eq. (7) would therefore be favoured for oxides with thermodynamically stable lattices and a high concentration of oxygen ion vacancies, e.g. as a result of a high acceptor concentration. Eq. (8) describes the formation of two OH bonds. Its thermodynamics should therefore mainly reflect the basicity of the lattice oxygen, i.e. Eq. (8) should be favoured for basic oxides forming strong OH bonds. The changes of the charge distribution associated with both Eqs. (7) and (8) are expected to be similar for different oxides and should therefore not enter significantly into the relative thermodynamics of hydration for oxides of similar dielectric behaviour. Unfortunately, there is not a lot of data available for K(T ) so far. Within the accuracy of the data for some perovskite type oxides shown in Fig. 1 there is a qualitative consistency with the above considerations. In the order Sr→Ba and Nb→Zr→Ce, Er for the occupation of A and B site of perovskite type oxides, the standard hydration enthalpies become more negative in accordance with the increasing basicity of the corresponding oxides. Unfortunately,

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Fig. 1. Equilibrium constant K for the hydration reaction of different perovskites. The standard hydration enthalpies obtained from the slopes are indicated, and the standard hydration entropy can be read from the intersection with the ordinate (data are taken from [4,5,7,8]).

the few data do not yet allow verification of the effect of lattice and defect relaxation energies on K(T ). It is interesting to note that the standard hydration entropy roughly scales with the hydration enthalpy and eventually reaches almost the standard entropy of water for oxides with the most negative enthalpies of hydration (Fig. 1). In these, almost the entire standard entropy of the water in the gas phase is cancelled by the hydration Eq. (1). Besides the equilibrium constant K(T ), the limiting concentration of protonic defects [OH ?O ]8 is the other relevant parameter determining the concentration of protonic defects. This is generally found to be lower than the effective acceptor concentration which has been suggested to be the result of a decreased activity coefficient for oxygen ion vacancies charge compensating for the acceptor doping, so that [OH ?O ]852f [V ??O ]8. Fortunately, the activity coefficient f is increasing with increasing dopant concentration [4] which has been explained by lattice relaxation effects around oxygen ion vacancies, which decrease with increasing concentration of oxygen ion vacancies. This also implies a dependence of f from [OH ?O ] which is indeed indicated by the hydration isotherms [4]. Although this second order effect suggests the absence of a well defined

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solubility limit, the apparent solubility limit will be used in the following qualitative considerations. The reduction of the water solubility limit also varies significantly for different oxides. Whereas the activity coefficient for a 3% Y-doped BaCeO 3 is still 0.2 [4], it is as low as 0.04 for a 3% Y-doped SrZrO 3 [9]. The very few data seem to indicate that the reduction of the oxygen ion vacancy activity coefficient is most severe for the most densely packed oxides (e.g. KTaO 3 , CaZrO 3 , SrZrO 3 ) as opposed to loosely packed oxides such as BaCeO 3 and Ba 3 CaNb 2 O 9 . But also the degree of covalency in the interatomic interactions may be a relevant parameter as suggested by the very low water solubility of rather loosely packed Y:BaZrO 3 single crystals (recent results from the authors’ laboratory). The actual concentration of protonic defects in different acceptor doped oxides is shown in Fig. 2 as a function of temperature for a given water partial pressure. At low temperatures the concentrations simply equal the saturation limit [OH ?O ]8 and are therefore independent on temperature. The temperature regime, in which the protonic charge carrier concentration descends via dehydration Eq. (1) is mainly determined by the enthalpies of this reaction, which are also indicated (see also Fig. 1). Whereas the presence of oxygen ion vacancies is not critical in some applications (e.g. fuel cells), conductivity contributions from electronic charge

carriers must be avoided. The extension of the oxygen partial pressure range, in which the concentration of electronic defects remains below a certain limit, is given by the reaction constant K3 , i.e. the electrochemical band gap, which is generally somewhat lower than the optical band gap [11]. This is a feature of the electronic structure of the oxide and in most cases determined by the energy levels of cation (e.g. M 13 / M 14 ) and oxygen (O 2 / O 22) redox couples, where the former is generally the species associated with the electron, and the latter corresponding to the electronic hole. Details of this situation have been discussed by Goodenough from a chemical point of view [12,13]. Apart from the influence of the crystallographic structure, small band gaps are expected for oxides containing small transition metal cations of the third period (e.g. Ti, V). But even for titanates this is still well above 2 eV, which provides a rather wide ionic domain at sufficiently low temperatures. The absolute position of the ionic domain with respect to the oxygen partial pressure, is determined by the relative stability of the other positively charged defects (OH ?O , V ??O ) competing with electronic holes (h ? ) as described by Eqs. (2) and (3) (K1 , K2 ). For BaCeO 3 the ionic domain extends over more than twenty orders of magnitude of the oxygen partial pressure for temperatures lower than 8008C as required for applications in fuel cells [14].

Fig. 2. Concentration of protonic defects as a function of temperature for different acceptor doped perovskites (data from [4,7,8,10]).

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2.2. Stability For the thermodynamic stability of the oxide with respect to a certain reaction relative stabilities with respect to the reaction products have to be considered. As already pointed out above, a high basicity of the oxide is advantageous for the formation of protonic charge carriers. On the other hand, basic oxides are expected to react easily with acidic or even amphoteric gases like SO 3 , CO 2 or H 2 O to form sulfates, carbonates or hydroxides. Therefore, the chemical stability will be discussed on the background of this conflicting situation which is relevant for many applications. Considering a simple perovskite of the type ABO 3 , the reaction with CO 2 may be written as: ABO 3 1CO 2 5AO1BO 2 1CO 2 5ACO 3 1BO 2 . For reactions of this kind, the formation enthalpy of the perovskite from the binary oxides, which mainly reflects the compatibility of the cations with the perovskite structure, and the stability of the carbonate with respect to AO help to understand its thermodynamics. The effects of both parameters almost cancel for the decomposition of perovskites with low perovskite tolerance factors, i.e. low ratios of the ionic radii rA /r B . For Sr and Ba cerates the thermodynamic data for the reactions with CO 2 are almost identical [15]. Whereas the formation of BaCO 3 from BaO is more advantageous than that of SrCO 3 from the less basic SrO (see Fig. 3a and b) the increased stability of BaCeO 3 compared to SrCeO 3 is obviously compensating for this. For perovskite type oxides with alkaline earth metals on the A-site one can therefore consider the stability with respect to the above reaction to be mainly determined by the choice of the B-cation. In accordance with an increasing perovskite tolerance factor one therefore observes an increasing stability in the order cerates→zirconates→titanates (Fig. 3). As pointed out above, the lattice energy and the hydration enthalpy are related. It is therefore interesting to compare the phase equilibria with the defect equilibria of Eq. (1). Whereas the oxide basicity favours the formation of protonic defects as well as the decomposition in acidic gases, the stability of the oxide is anticipated to increase the formation of protonic charge carriers but to suppress the decomposition reactions. This is, in fact, supported by

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the few experimental data which are available. As can be seen from Fig. 3, the equilibrium constant K(T ) for SrZrO 3 is higher with respect to the water partial pressure of decomposition into the hydroxide compared to the corresponding parameters for the thermodynamically less stable cerates. Nevertheless, the formation of protonic defects as well as the decomposition into hydroxides appear at lower water partial pressures for cerates compared to zirconates. The thermodynamic stability of the oxide may also have an impact on the mechanical stability of the microstructure of corresponding ceramics. This is especially critical for perovskites with big B cations, i.e. with low tolerance factors and low thermodynamic stabilisation with respect to the binary oxides. As can be seen from Fig. 3a and b the decomposition partial pressures for the alkaline earth cerates are only a little higher than for the corresponding alkaline earth oxides which is due to the small Gibbs free energies of formation from the binary oxides. There is even a recent paper which claims that BaCeO 3 is instable with respect to its decomposition into BaO and CeO 2 [17] which is still within the accuracy of the other thermodynamic data. In a recent paper [18] it has been shown that Ba 21 may even occupy the B-site in BaCeO 3 , which leads to a further destabilization of the structure. It is probable that such less favorable structures may first tend to change in their grain boundary region thus weakening the microstructure as observed for BaCeO 3 -based ceramics [19]. But even for relatively stable perovskites (BaTiO 3 , SrTiO 3 , PbTiO 3 and KNbO 3 ) structural instabilities leading to the forma´ and Ruddlesden– tion of an intergrowth of Magneli Popper phases at the surface of the perovskite have been identified by different experimental techniques [20]. The latter has been suggested to be induced by changes in the oxygen partial pressure and seems to go along with the formation of peroxides at the surface of the perovskites.

2.3. Mobility of protonic charge carriers In perovskite type oxides the diffusion coefficient of protons is generally orders of magnitude higher than for oxygen indicating that proton diffusivity must involve proton transfer between oxygens and hydroxide ion reorientation. Although the latter is an

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Fig. 3. Phase equilibria and the defect equilibrium for the hydration reaction of cerates and zirconates (data are taken from [4,7,15,16]). The CO 2 partial pressure of air and the possible application ranges in air (shaded) are indicated (see text).

activated process, quantum MD simulations [21,22] as well as QNS experiments [23] suggest that this is a very rapid process and particularly not rate limiting proton mobility. Considering the proton transfer to be critical, it is surprising that the highest proton diffusivities are observed for perovskites with high lattice constants corresponding to large structural oxygen separations Q o . In the case of BaCeO 3 , Q o 5312 pm at room temperature corresponding to a proton transfer barrier of about 1.8 eV as revealed by a quantum chemical calculation for a fixed oxygen separation Q [22].

For such a large oxygen separation one does not expect a significantly large hydrogen bond interaction providing a path for proton transfer. In pure hydroxides, only small hydrogen bond interactions indicated by small red shifts of the OH stretching frequency (nOH ) are observed. The small tendency of basic OH groups towards acting as a proton donor even allows a blue shift of nOH with respect to a free OH 2 (n OH 53556 cm 21 ) depending on the metal– oxygen interaction [24]. The IR spectra of hydroxide ions in some perovskite type oxides (Fig. 4 [25– 29]), however show distinctly different features.

K.D. Kreuer / Solid State Ionics 97 (1997) 1 – 15

Fig. 4. IR-spectra in the OH stretching region for different perovskites containing protonic defects OH O . The frequency of a free OH 2 is indicated for comparison.

Especially perovskites with large structural oxygen separations show a tremendous red shift of the OH stretching frequency thus indicating a softening of the fundamental vibration already, especially for those oxides, which also show the highest proton mobilities. One reason for these strong hydrogen bonds in oxides compared to hydroxides is definitely the oxide ion being a stronger proton acceptor than the hydroxide. The increase of this effect with increasing Q o , however, can only be explained by the increasing thermal vibrations of the oxygen. The dynamics of the next nearest environment of the proton is thought to significantly assist proton diffusivity [30], and it is the excitation of the O–B–O bending mode which has been suggested to be involved in the transition state of proton transfer in ABO 3 perovskites. This greatly determines the fluctuations of the oxygen separation coordinate Q and therefore also the respective hydrogen bond interaction, which increases with decreasing donor / acceptor separation. The oxygen vibrations induce transient hydrogen bonds which may even lead to a

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relative stabilisation of configurations in the vicinity of the transition state. Therefore, the softness of the oxygen separation coordinate Q seems to be more important for a high proton diffusivity than the structural oxygen separation Q o . This situation can only be described quantitatively by considering the full dynamics of the oxide. This is, e.g., accessible by quantum MD simulations which have been carried out for barium cerate, zirconate and titanate [21,22]. Fig. 5 shows the Gibbs free energy changes of the systems as a function of the oxygen separation DG(Q). The transition state parameters are obtained from the minima in the sums of DG(Q) and the proton transfer barrier as a function of Q. Within the accuracy of such simulations, the few data indicate no clear relation between the transition state parameters and the packing of the perovskite. To what extend neighbouring oxygens approach for a given thermal excitation may rather be described by the packing defining the structural oxygen separation Q o and the stiffness of the B–O bond mainly determined by its covalency. The latter appears to be very low in BaCeO 3 thus allowing extended oxygen vibrations and an energetically low proton transfer transition state despite the large structural oxygen separation (Q o 5312 pm). The situation in the zirconate is less advantageous because of the rather rigid oxygen separation coordinate. The lowest transition state is indeed observed for BaTiO 3 which combines a rather small structural oxygen separation (Q o 5286 pm)

Fig. 5. Gibbs free energy change DG(Q) of Ba-cerate, zirconate and titanate as a function of the oxygen separation co-ordinate Q obtained by quantum MD-simulations [22]. The calculated transition state parameters for proton transfer are shown in the insert.

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with a medium rigid oxygen separation coordinate. Rapid hydroxide reorientation is observed in all cases anticipating that high proton diffusivities may also be observed in closely packed oxides such as BaTiO 3 . But more data from experiments (e.g. reliable transport data from single crystals, local dynamics from diffraction and scattering experiments) and simulations are required for a more detailed and general understanding. It should be mentioned that association of protonic defects with the acceptor dopant has been used for the interpretation of QNS [31] and IR data [32]. Recent PFG-NMR experiments on BaCeO 3 based crystals, however, show that the proton diffusion coefficient is almost independent of the proton concentration and on the kind and concentration of dopant indicating only minor defect interactions at least for high proton concentrations [4]. Only for very low concentrations of protonic defects the increase of the activation enthalpy of the proton diffusivity suggests some contribution from the association of protonic defects with acceptor dopants.

2.4. Perovskite type oxides for electrochemical applications One important requirement for any application is the stability of the material under the respective operation conditions. Most applications comprise the presence of CO 2 , either due to the presence of air or as an oxidation product of organic fuels. The CO 2 partial pressure of air defines the lowest operation temperature in this environment as indicated in Fig. 3. For cerates, zirconates and titanates this is around 700, 350 and 2008C respectively. On the other hand, the formation of protonic charge carriers at a given pH 2 O is characterized by an upper temperature limit, where dehydration occurs (see also Fig. 2). Assuming a technologically feasible maximum pH 2 O of 100 hPa and a minimum concentration of protonic defects [OH O ].0.8[OH O ]8 (this corresponds to KpH 2 O .1) this yields the application ranges indicated as the shaded fields in Fig. 3a–c. For cerates this is actually a very small region just above 7008C where for Sr-zirconate a possible operation window

between approximately 350 and 4508C is anticipated. From the discussion in Section 2.3 it is expected that this range is increasing with increasing thermodynamic stability of the perovskite and shifting to lower temperatures with increasing acidity of the B cation. The latter has some influence on the OH bond strength in the perovskite but does not enter into the thermodynamics of decomposition as long as the B cation is not involved in the formation of carbonates. Frequently both parameters are expected to change, like in the case of Nb 15 on the B position, which is expected to introduce more acidity into the perovskite lattice but also a higher thermodynamic stabilisation due to a more advantageous perovskite tolerance factor. While the equilibrium constant K(T ) essentially defines the upper temperature limit with respect to dehydration, a high solubility limit is important for a high concentration of protonic defects. This requires a high acceptor concentration and a loosely packed structure in order to avoid a severe reduction of the oxygen ion vacancy activity coefficient. High proton mobilities are compatible with loosely packed structures providing the absence of too much covalency especially in the B–O bond which tends to suppress the required oxygen dynamics. Considering only simple alkaline earth perovskites the following rules for the occupation of A and B-site may help to find the appropriate perovskite for a given application. With respect to thermodynamic stability, heat of hydration, water solubility limit and mobility of protonic defects the occupation of the A-site does not require much of a compromise. Except for the stability with acidic gases, which is almost independent of the choice of the A-cation, all relevant properties are superior for an A-site occupation by the big barium compared to other alkaline earth ions. The choice of the B-cation, however, requires some compromising. It should be of medium size, amphoteric nature and should form no significant covalent bonding with its oxygen ligands. High packing densities as a result of small B-cations reduce the water solubility limit, whereas the poor compatibility of big B-cations with the perovskite structure reduces the thermodynamic stability including the stability in acidic gases. For too acidic

K.D. Kreuer / Solid State Ionics 97 (1997) 1 – 15

B-cations the hydration enthalpy is not negative enough to retain the protonic defects up to the operational temperature, whereas too basic B-cations may cause decomposition reactions in acidic gases. The occupation of the B-site with ions of different acid / base properties is also of great interest. This is expected to further increase the thermodynamic stability while the basicity is expected to be some average of that of the binary oxides. The acceptor dopant concentration should be as high as possible. This not only increases the concentration of oxygen ion vacancies [V ??O ] accessible for the hydration reaction, it also minimizes the lattice relaxation around oxygen ion vacancies thus leading to sufficiently high activity coefficients f([V ??O ]) of oxygen ion vacancies and therefore to high concentrations of protonic defects [OH ?O ]~( f([V ??O ])[V O?? ])1 / 2 . Even the association of protonic defects with acceptor dopants is expected to be less severe for high dopant concentrations because of the more even charge distribution. For applications which require high proton mobility but not a high concentration of protonic defects somewhat smaller A and B-cations may be chosen. This drastically reduces the water solubility but high proton mobilities may be obtained as observed for SrZrO 3 [26] and anticipated for BaTiO 3 [22]. Ca doped Ba 2 CaNb 2 O 9 , which has been suggested by Nowick as a proton conductor [27], seems to be interesting in these respects. Preliminary results indicate very good thermodynamic stability and proton conductivities as high as those in BaCeO 3 based materials. The very low hydration enthalpy (0.6 eV [8]) however leads to a loss of protonic charge carriers at temperatures not much above 2008C (Fig. 2) where the thermally activated proton mobility is not yet sufficiently high for, e.g., fuel cell applications (Fig. 6). Less acidic cations on the B position which provide a similar tolerance factor may be more advantageous as indicated by the hydration and conductivity behaviour of Ba 2 YSnO 5.5 (Figs. 2 and 6). Complex perovskites such as Ca:Ba 2 CaWO 6 may even combine a similar basicity of the lattice oxygen with an even higher thermodynamic stability due to an extra energy gain as a result of internal acid–base reactions. But also perovskites with big trivalent cations on A and B site such as Ba:LaYO 3

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Fig. 6. Proton conductivity of different acceptor doped perovskites (the doping levels are identical to those in Fig. 2). References may be taken from [1]; the data on the niobate [8] and the stannate [10] have not been published yet.

[2,5] should be considered. The avoidance of any transition metal cations has the additional advantage of a more extended ionic domain (see Section 2.1) and a reduced tendency to form covalent bonds.

3. Low temperature proton conducting polymers Stimulated by the legislative pollution control in most industrialised nations, the development of polymer electrolyte membrane fuel cells (PEM-FC) has attracted an increasing interest. Generally perfluorinated ion exchange membranes in their protonic form such as Nafion are being used as membrane materials. These combine the required chemical, electrochemical and mechanical stability with high proton conductivity. In fact, such membranes have been used in hydrogen PEM fuel cells for a long time, e.g. in space crafts and submarines, but the technology is far too expensive for application in mass products such as cars. One essential reason for this is the cost of the proton conducting membranes. This situation stimulates the development of new

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proton conducting polymers. In contrast to oxides, where applicable materials are hardly available, Nafion meets most requirements except for a tolerable price. Therefore, an interesting approach is to first study Nafion, in order to understand its superior combination of properties, and then to materialize similar situations by a less expensive chemistry.

3.1. Nafion Nafion is the trade mark of a perfluorosulfonic polymer with the following structure:

Fig. 7. Water diffusion and conductivity diffusion coefficient (DH 2 O , Ds ) for Nafion as a function of temperature and water content (recent data from the author’s laboratory).

It combines the extreme hydrophobicity of the perfluorinated polymer backbone with the extreme hydrophilicity of the terminal sulfonic acid function (–SO 3 H). Already in the dry state, this leads to a spontaneous hydrophilic / hydrophobic nano separation [33]. In the presence of water only the hydrophilic part of the micro structure is hydrated. The water of hydration then acts as a plastiziser mobilising the polymer backbone which leads to a further phase separation. Eventually, a stationary micro structure is formed which absorbs and desorbs water almost reversibly at moderate temperatures. One important feature of such micro structures is that their hydrophobic part provides relatively good mechanical stability even in the presence of water, while the hydrated hydrophilic domains provide the very high proton conductivity. The latter depends very much on the presence of water. As can be seen from Fig. 7 the activation enthalpy for water diffusion in such membranes is almost identical to that in pure water for water contents higher than required for approximately primary hydration of the sulfonic acid functions (n5[H 2 O] / [SO 3 H].3) while the absolute value of the diffusion coefficient is decreasing with decreasing water content [34]. These features have been interpreted with water dynamics on a molecular scale

(nm) similar to that of liquid water and geometric restrictions for water diffusion within the heterogeneous microstructure increasing with decreasing water content. The polymer material acts as a nano porous inert ‘sponge’ for the water of hydration, which in fact shows very little interaction with the polymer except for the first three water molecules per sulfonic acid group required for its primary hydration as indicated by the water absorption isotherm (Fig. 8). Fig. 7 also shows the proton conductivity diffusion coefficient Ds which roughly follows the water diffusion coefficient DH 2 O , i.e. the conductivity is essentially carried by the diffusion of hydrated protons such as H 3 O 1 and H 5 O 1 2 which originate from the complete dissociation of the sulfonic acid functions.

3.2. New proton conducting membrane materials With respect to the above discussed transport properties the situation in Nafion appears to be quite simple, and indeed there are other sulfonated membrane materials such as sulfonated styrene grafted FEP [35], sulfonated ORMOCERs [36], partially fluorinated polystyrenes [37] and organically modified layered phosphonates [38] which show similar transport coefficients. But the lack of either chemical

K.D. Kreuer / Solid State Ionics 97 (1997) 1 – 15

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Fig. 9. Water diffusion and conductivity diffusion coefficient (DH 2 O , Ds ) for a sulfonated polyaromatic membrane [41]. Fig. 8. Water absorption isotherm for Nafion (recent data from the author’s laboratory).

or mechanical stability still prevents them from substituting perfluorosulfonic membranes in fuel cell applications. Failure analyses demonstrate the chemical instability of non-fluorinated aliphatic sections of such materials while not only the perfluorinated but also the aromatic constituents appear to be durable under fuel cell conditions. Therefore, membranes based on purely aromatic high performance polymers such as Ryton and PEK [39] offer an interesting option for the development of new chemically durable proton conducting polymers [40,41]. Some of them allow direct electrophilic sulfonation and casting from organic solutions, which is by far a less expensive process than compared to the fabrication of perfluorosulfonic membranes. However, transport and mechanical properties are distinctly different from those of Nafion. Fig. 9 shows Ds and DH 2 O for a sulfonated polyaromatic membrane [41]. Only for very high water contents both diffusion coefficients are almost identical as observed for Nafion. With decreasing water content, however, Ds decreases more rapidly than DH 2 O corresponding to A 5 Ds /DH 2 O dropping significantly below unity as shown in Fig. 10, which also includes the corresponding data for Nafion and aqueous solutions of hydrochloric acid. The data

Fig. 10. The ratio A 5 Ds /DH 2 O for aqueous solutions of hydrochloric acid, hydrated Nafion and a sulfonated polyaromatic membrane [41].

indicate an increasing association of the protonic charge carriers with the anionic counter charge (–SO 2 3 ) immobilized on the polymer backbone. Although the separation between the sulfonic acid functions is well below the Debye length of water (approx. 800 nm at room temperature) the effect of electrostatic potential differences due to space charges surrounding the sulfonic anions can not be neglected even for small separations of the sulfonic anions. A perfectly homogeneous distribution of sulfonic acid functions correspond to a separation of 1.2 nm and a maximum electrostatic potential differ-

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ence of approximately 70 meV between neighbouring sulfonic functions. This energy may be considered an association energy which is expected to show up as a contribution of the activation energy of proton diffusivity. This, however, can hardly be observed in Nafion, in which the condensation of sulfonic functions in the hydrophilic phase obviously decreases their mutual separation to such an extent that ionic association effects practically disappear. In aromatic polymers the observed association effects suggest a less effective phase separation. This also has a negative impact on the morphological stability as can be seen from comparing the swelling behaviour of a sulfonated aromatic polymer to that of Nafion (Fig. 11). Whereas for the latter irreversible swelling starts only for temperatures higher 1308C, which corresponds to the glass transition temperature of Nafion, the aromatic membrane is morphologically stable in water only up to 808C. It is worth noting that the extra water incorporation associated with the swelling leads to a further increase of the proton conductivity. Thermodynamically speaking, swelling increases the amount of water as a second phase, and the increase in conductivity is the result of changes of a two phase effect [42]. As can be seen from the hydration

isotherm of a polyaromatic membrane (Fig. 12) and the room temperature proton conductivity as a function of the water content for this membrane compared to that of Nafion (Fig. 13), high proton conductivity in the polyaromatic membrane relies

Fig. 11. Water uptake of Nafion and a sulfonated polyaromatic membrane in water as a function of temperature [41].

Fig. 13. Room temperature proton conductivity of NAFION and a sulfonated polyaromatic membrane as a function of water content. The conductivity contributions depending on the amount of water as a second phase are indicated.

Fig. 12. Hydration isotherm of a sulfonated polyaromatic membrane. The increase of the amount of water as a second phase as a result of swelling is indicated.

K.D. Kreuer / Solid State Ionics 97 (1997) 1 – 15

much more on the amount of water as a second phase than is the case for Nafion. This is also due to the less effective phase separation resulting in a poorer connectivity of the hydrophilic phase for a given water content. The comparison of the properties of the two polymers shows that nano separation has exclusively positive effects on the considered properties, i.e. proton and water diffusion and morphological stability. Therefore, the control of the nano structure of polyaromatic and other membranes provide an interesting route for simultaneously improving most membrane properties relevant for fuel cell applications. In particular, this approach does not require any compromising with respect to conductivity and stability.

4. Intermediate temperature proton conductors Whereas some proton conducting oxides and hydrated polymers are close to being technologically applied at high (.5008C) or low temperatures (,908C) respectively, only few materials have been reported to show high proton conductivity at intermediate temperatures (100–4008C). This gap arises from the fact that all compounds, for which proton conductivity has been reported so far, belong to a limited number of families of compounds with respect to the species ‘solvating’ the proton [1]. In the case of polymers, this is either water as in the case of Nafion or pure oxo-acids as in the case of polybenzimidazole?H 3 PO 4 [43], where these moieties participate in the formation of protonic charge carriers (e.g. H 3 O 1 , H 4 PO 41 , H 2 PO 42 ) their dynamics being essential for the diffusivity of such defects. The potential for a further development of the available materials is, however, limited because of the inherent properties of the species the proton conductivity is linked to. In particular, they impose constraints on the operation temperature and the chemical and electrochemical stability with other materials. Therefore, it is not only very interesting to adjust the polymeric host to the requirements of a particular application as described in the preceding section, but also to optimize the intimate environment of the proton. In order to support high proton diffusivity this must allow for the formation of

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protonic defects and provide strongly fluctuating proton donor and acceptor functions in an otherwise non-polar environment [1]. Heterocycles such as imidazole, which already have been shown to exhibit moderate proton conductivity as a pure material [44], are promising in this respect. Their basic nitrogen sites act as strong proton acceptors with respect to Brønsted acids such as sulfonic acid groups thus forming protonic charge carriers (C 3 H 3 NH 2 )1 . The rather compact molecules are advantageous for extended local dynamics, and their protonated and unprotonated nitrogen functions may act as donors and acceptors in proton transfer reactions while the ring itself is rather non-polar. The melting points of such molecules are higher than that of e.g. water or phosphoric acid, which makes them interesting candidates for supporting proton conductivity at medium temperatures. Their high basicities provide another interesting feature with respect to the chemical compatibility with other compounds such as metal hydrates or hydrogen bronzes being present in certain applications. Fig. 14 shows the conductivity of a sulfonated polyaromatic membrane (see Section 3.2) with imidazole intercalated in different concentrations n5

Fig. 14. Conductivity of imidazole intercalated into a sulfonated polyaromatic membrane [45]. The conductivity of the hydrated polymer is given for comparison.

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K.D. Kreuer / Solid State Ionics 97 (1997) 1 – 15

[C 3 H 3 NH] / [–SO 3 H] [45]. For the highest imidazole content a conductivity of 2310 22 V21 cm 21 is observed at 2008C which is comparable to the conductivity in the hydrated polymer at lower temperatures (Fig. 14). A detailed comparison of the transport data for the hydrated and the imidazole containing polymer shows striking similarities, where the transport coefficients of the imidazole containing systems are just shifted to higher temperatures with respect to the hydrated system by about 80 K. Of course, this only shows that high proton conductivity can be materialized by so far less considered families of compounds. The further development of such systems for particular applications, however, requires the simultaneous optimization of all relevant properties as described above for high and low temperature proton conducting materials.

Acknowledgments The author thanks J. Clauß and G. Frank (Hoechst AG, Frankfurt) for providing the polymers, A. Fuchs ¨ ¨ for technical assistance and W. Munch, H.J. Schluter and B. Gibson for reading the proofs. Part of the work is being supported by the BMBF under the contract number 0329567

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