Ion Conducting Membranes for Fuel Cells and other Electrochemical

Perspective pubs.acs.org/cm

Ion Conducting Membranes for Fuel Cells and other Electrochemical Devices Klaus-Dieter Kreuer* Max-Planck-Institut für Festkörperforschung, Heisenbergstrasse 1, D-70569 Stuttgart, Germany ABSTRACT: Transport and stability issues of proton and hydroxide ion conducting separator membranes for fuel cells are critically discussed from a fundamental point of view. Considerations of structure and dynamics on the molecular scale to the device level equally imply polymer-chemical and electrochemical aspects which are closely related for this class of materials. The importance of ion/solvent, residual ion/ion, and solvent/polymer interactions for the formation and mobility of ionic charge carriers and selective ionic transport and even as driving forces for nanoscale ordering is emphasized, and it is shown that, apart from simple electrostatics, specific chemical interactions must be considered. On the basis of this understanding, suggestions are being made for the modification of existing and the development of new membrane types, not only for fuel cells but also for other electrochemical energy conversion and storage devices such as redoxflow and alkaline ion batteries. KEYWORDS: ion conducting membrane, fuel cell, redox-flow battery, Li ion battery, proton, hydroxide, diffusion, conductivity, nanomorphology, hydration, visco-elastic constants, phosphate, polyelectrolyte, ionomer, block-copolymer, Nafion, Aquivion

1. INTRODUCTION Most electrochemical conversion and storage devices, such as f uel cells, redox-f low, and alkaline ion batteries, rely on the amazing properties of ion conducting polymer membranes. These devices can function only if the used membranes efficiently separate the electrochemically active masses (electrodes) and mediate the electrochemical reactions taking place at the anode and cathode by conducting specific ions. These ions may be protons or hydroxide ions in the case of PEM-fuel cells, Li+ in lithium ion batteries, or non-electrochemically active anions or cations in the case of redox-f low batteries. Apart from these key properties, many other requirements, especially concerning stability, render the development of such membranes a formidable task! This article briefly summarizes our current understanding of membrane structure, reactivity, and dynamics on different length scales with special emphasis on some surprising recent insights; these have not only extended the understanding of well-established membrane materials such as Nafion and polybenzimidazole-phosphoric acid adducts but also consolidate the basis for the modification of existing and the development of new membrane types for the above-mentioned applications. When done efficiently, this is a multidisciplinary process comprising ab initio calculations, complex organic/ inorganic synthesis, and comprehensive physicochemical characterization. The great beauty of such a combined approach is exemplified by retracing the development of proton conducting multiblock copolymers for PEM f uel cell applications and Li+ conducting polyelectrolytes for application in Liion batteries. © 2013 American Chemical Society

The present article develops a few perspectives from a specific point of view, in that both the analysis of the state of the art and the presented research perspectives are biased by the author’s own work. Because of the limited space and the increasing abundance of the relevant literature, references are restricted to a few key papers, reviews, and our own work. For different and broader views, the reader may refer to the huge number of existing reviews in the field. A recent publication of the American Chemical Society1 can serve as a suitable starting point for accessing the comprehensive literature.

2. TYPES OF ION CONDUCTING POLYMERS AND THEIR APPLICATIONS IN ELECTROCHEMICAL ENERGY CONVERSION AND STORAGE It goes without saying that common polymers such as polyethylene (PE) or Teflon (PTFE) do not conduct ions per se. Generally speaking, ionic conductivity occurs only in the presence of moieties dissociating into ionic species which, to some extent, diffuse in an uncorrelated way allowing the irreversible separation of ionic charges, which is a requirement of any ion conduction process. Special Issue: Celebrating Twenty-Five Years of Chemistry of Materials Received: August 14, 2013 Revised: October 23, 2013 Published: November 19, 2013 361

dx.doi.org/10.1021/cm402742u | Chem. Mater. 2014, 26, 361−380

Chemistry of Materials

Perspective

Figure 1. Ion conducting polymeric separators in eletrochemical conversion and storage devices: (a) fuel cells, (b) redox-flow batteries, and (c) alkaline metal ion (Li+, Na+) batteries. Note that the direction of charge flow corresponds to the discharging mode.

leads to a concentration of ionic groups within single domains of their phase- separated morphology. Because of the high local concentration of sulfonic acid groups, these domains are polyelectrolytes,2 rather than ionomers, and show high proton conductivity even at low levels of hydration. This, together with acceptable mechanical properties emerging from the unsulfonated domain, renders these and related classes of ion conducting polymers an interesting alternative to PFSA membranes. Apart from the major focus on sulfonic acid functionalized polymers, there is an increasing interest in anion-exchange polymers because of the promise of using them in fuel cells with non-noble metal catalysts within their electrode structures. Most anion exchange membranes are hydrocarbons, and virtually all contain some sort of quaternary amine as basic functional groups. They fall into the category of ionomers albeit with relatively high concentrations of anion exchanging groups, making cross-linking necessary in order to prevent the membranes from dissolving in water. Recent progress in the field notwithstanding, stability issues and the fact that any membrane of this type readily converts into its bicarbonate (HCO3−) form make the application of anion exchange polymer membranes in PEM fuel cells scarce. It is worth mentioning that for a special case of PEM fuel cells, the direct methanol fuel cell (DMFC), PFSA membranes are still used because of their robustness and the availability of compatible electrode structures, although hydrocarbon membranes generally exhibit lower water and methanol permeation. Hydrocarbons with very polar backbones such as highly sulfonated polyphenylene sulfones even show very high selectivity for water compared to methanol absorption.3 Traditionally, PFSA membranes are also being used as separators in redox-flow batteries (Figure 1b), but more recently hydrocarbon based anion exchange membranes are considered as well. Especially when the redox active species are cations (like in vanadium flow batteries), the high permselectivity of many hydrocarbon based anion exchange membranes is a particular advantage, and it is interesting to note that many membranes which are not stable in their hydroxide (OH−) form show acceptable stability in their, e.g., Cl−, Br−, or SO42− form under acidic conditions, which are typical for many types of redox-flow batteries. All membrane types mentioned so far require some hydration to conduct ions, and the diffusions of hydration water and ions are generally related. Approaches aiming at “dry” proton conducting polymers containing no low-molecular-

In the case of the best known ion conducting polymer, Nafion, sulfonic acid groups (−SO3H) are part of the macromolecular structure; i.e., they are covalently immobilized. Since these groups are super acids, they strongly interact with water leading to hydration (solvation) and virtually complete dissociation, even at low water activity (relative humidity). The hydration water not only stabilizes the separation of cations (here protons) and anions (here sulfonic anions), it also enables the mobility of the hydrated protons. While initially developed as a Na+ conducting separator membrane for chloralkali electrolysis, Nafion in its acid form (the sulfonic acid functional group is sometimes termed protogenic group) later became the benchmark membrane for PEM-fuel cell applications as well (Figure 1a). According to the IUPAC nomenclature,2 Nafion belongs to the family of cation-exchange polymers because, in contact with aqueous solutions, especially monovalent cations can easily exchange, albeit with some preference, for certain ions. With respect to the molecular structure, Nafion is an ionomer, a polymer with a small but significant proportion of the constitutional units having ionic or ionizable groups.2 In fact, Nafion is the proto-type of a class of perfluorosulfonic acids (PFSA) which comprises a variety of different side chain architectures. In particular, short side chain (SSC) PFSAs seem to have the potential to outperform Nafion when it comes to increasing the operation temperature of PEMfuel cells. Although the modification of PFSA membranes, e.g., by dispersing nanoparticles in the membrane, became extremely popular, there is no accepted understanding of the reported effects (e.g., conductivity enhancement). What seems to be clear at this point is that such composites show higher conductivity at low levels of hydration only if the particles add to the total ion exchange capacity (IEC) of the membrane. Membranes hosting particles with a high degree of sulfonic acid functionalization seem to have systematically higher conductivities. There is also a wealth of sulfonic acid functionalized hydrocarbon membranes with conductivities similar to these of PFSAs at high degrees of hydration, but low conductivities at low hydration levels and unsatisfactory mechanical properties are unfortunately characteristic features making them inferior to the more costly PFSA membranes for many applications. This applies to hydrocarbons with a random distribution of sulfonic acid groups which are still ionomers, as opposed to di-, tri-, and multiblock copolymers with a similar overall concentration of ionic groups. In the latter case, the protogenic groups are located on one kind of block of the molecular structures which 362



Perspective

interaction with the polymeric matrix, the remaining proton conductivity is still high enough for fuel cell applications. The dependence of ionic conductivity on the presence of low molecular weight solvent molecules is also a critical issue for alkaline ion conducting electrolytes, e.g., Li ion conductors used in lithium ion batteries (Figure 1c). Here, standard electrolytes are solutions of lithium salts (e.g., LiCF3SO3, LiPF6) in aprotic polar solvents such as ethylene or propylene carbonate (EC, PC) and dimethyl carbonate (DMC). Although the ionic conductivity has contributions from both cations and anions, the total conductivity is generally more than 1 order of magnitude lower than that of proton conducting membranes used in PEM-fuel cells. At the moment this is not a severe problem, since the current densities of current batteries are limited by poor electrode kinetics in most cases, and small separator thicknesses may compensate for low specific conductivities. There are also polymers, such as polyethyleneoxide, which dissolve Li-salts to form solid polymer electrolytes,2 but as in the case of proton conducting polymers, the ionic conductivities remain another order of magnitude below typical conductivities of liquid systems. Therefore, most separator materials in Li-batteries contain some low molecular weight solvents, even when this is trapped in gel-like structures within a stable polymer network. Nevertheless, fully polymeric, nonflammable separators are highly desirable because their use may solve the severe safety issues of alkaline ion batteries.

weight solvent (e.g., water) essentially make use of moieties such as heterocycles and phosphonic acids which combine proton donor and acceptor functions with the possibility of covalent immobilization as terminal functions of fully polymeric structures.4−7 In this way, it was possible to completely decouple the long-range diffusion of protons from this of the polymer,4 but because of the inherent conductivity limits (usually more than 1 order of magnitude lower than this of hydrated systems, Figure 2) and interferences with the oxygen

Figure 2. Proton conductivity of some typical phosphorus5−7,12 acid (green) and heterocycle functionalized membranes (orange)4 as compared to the conductivity of pure phosphoric acid10 and hydrated sulfonic acid based systems (blue).25,40,53 Note that the conductivities of phosphonic and sulfonic acid systems were recorded at a water pressure of pH2O = 1 atm (105 Pa).

3. MEMBRANES FOR PEM-FUEL CELLS Polymer electrolyte membrane (PEM) fuel cells mostly use hydrogen or hydrogen-rich reformates as fuel, and oxygen is usually supplied as a humidified air stream within a temperature range constrained by the properties of the membrane material. In the case of Nafion, this is limited to about T = 90 °C because of the increasing hydration requirement with increasing temperature, the decay of the morphological stability in this temperature range, and the membrane decomposition through the attack of peroxo and hydroxo radicals. These radicals form at the surface of the platinum electrocalalyst in the presence of hydrogen and oxygen. Especially at OCV (open circuit potential) conditions, highly dispersed platinum of the cathode structure oxidizes, dissolves into the acidic membrane, and migrates toward the anode side until precipitation occurs within the membrane when hydrogen crossing over from the anode side is encountered.13 Among others, the main research objectives for this class of ion conducting polymers therefore are reduced humidity dependence of the proton conductivity, reduced acidity, lower gas permeation, and improved chemical and mechanical stability. When hydroxide ion (OH−) conducting anion exchange membranes are used, lifetime and performance are also limited by the membrane properties. Such fuel cells can currently operate up to a temperature of only T ∼ 60 °C. As a strong nucleophile, OH− ions start to attack the chemical bonding of the anion exchanging group (quarternary ammonia or some other amine) of the polymeric structure, while reaction with acidic CO2 from the air converts OH− into bicarbonate (HCO3−), which shows little electrochemical activity (see Section 3.2). In the case of HT-PEM fuel cells operating at T ∼ 160 °C, the commonly used PBI−phosphoric acid membranes have demonstrated stable performance for 104 hours, but phosphoric acid leaching upon start/stop cycles and inferior mechanical properties at the required high phosphoric acid “doping” levels

reduction reaction at platinum surfaces,8 such ion conducting polymers have not yet found their way into electrochemical applications. Although this class of materials is not discussed in the present article, one should keep an eye on fully polymeric electrolytes for several reasons. The fact that these are not or only moderately acidic makes them chemically compatible with less costly non-noble metal catalysts, and the use of very thin membranes may overcome the problems associated with their inherently low specific conductivity. While proton conductivities of heterocyclic (especially imidazole, pyrazole, and triazole functionalized) systems are almost insensitive toward changes of the relative humidity (RH), it must be mentioned that, in the case of phosphonic acid functionalized systems, some water activity is required to prevent the systems from condensation (Figure 2) which is strongly detrimental to proton conductivity.5 As expected from the related acidity increase, fluorination of phosphonic acid functionalized polymers leads to a suppression of condensation and an increase of hygroscopicity and proton conductivity.9 Interestingly, the problem of condensation is less severe for pure phosphoric acid, a highly viscous liquid with the highest intrinsic proton conductivity of any compound.10,11 While phosphoric acid fuel cells (PAFC) take advantage of the unique properties of phosphoric acid absorbed in some porous matrixes, adducts of polybenzimidazole with phosphoric acid are used in high-T PEM fuel cells operating around T = 160 °C.12 Under these conditions hydrogen rich reformates can be used as fuel without further purification. Although phosphoric acid loses much of its conductivity through the acid/base 363



Perspective

fraction Φwater as low as 0.2 (20%) corresponding to a relative humidity (RH) of about 80%). But our own work17 raised severe doubts about this model: a constant number of cylinders requires significant structural reorganization to adjust changes of the water content which is in contrast to extremely fast equilibration once water has entered the membrane.18,19 It is also the severe separation of different and accumulation of equal charges (e.g., protonic charge carriers accumulating within the interior of the cylinders) which is energetically unfavorable when the charges are not completely screened by the water of hydration. We have therefore suggested the water structures to be locally f lat and narrow which allows the protonic charge carriers to electrostatically interact with several sulfonic groups.17 Generally speaking, water “f ilms” may act as positively charged “glue”, keeping together the oppositely charged polymer structures. In the case of Nafion, this is not only evidenced by the linear scaling of the structural correlation length (obtained from the position of the ionomer peak) with the polymer volume fraction but also by reasonable fits of the regime of SAXS patterns representing correlations in the 1−5 nm range using water volume fractions, determined experimentally with a high precision. While the parallel cylinder model was most likely biased by a large uncertainty of the experimentally determined water content, locally flat morphologies are not only consistent with the evolution of Nafion SAXS patterns recorded over a wide range of water contents (up to a water volume fraction Φwater = 1 − Φpolymer ∼ 0.5), but they also seem to be a common feature of most dissociated ionomers and polyelectrolytes (Figure 3). Details of this

are still serious downsides relevant for their use in fuel cells. Other disadvantages are the required high platinum loading of the oxygen electrode and the very high activation potential of the oxygen reduction reaction (ORR). Relatively low current densities actually let the fuel cell potential drop into the range around U = 0.7 V which is about 60% of the OCV at the operation temperature, and the related efficiency loss may not be acceptable when compared to other technologies. 3.1. Systems Containing Sulfonic Acid. As indicated in the Introduction, the behavior of ion conducting polymers in fuel cells is affected by transport and stability issues on different length and time scales. Since our current understanding of this complex situation is far from complete, I will focus on some recent advances in the field in an attempt to approach a more consistent comprehension of the complex relationships. 3.1.1. PFSA Membranes and Their Modifications. Perfluorosulfonic acid (PFSA) membranes generally consist of a polytetrafluorethylene (PTFE) backbone with perfluorinated side chains of different lengths attached to the backbone through ether linkages and terminated by sulfonic acid (−SO3H) groups (Scheme 1). The unique properties of Scheme 1. Molecular Structures of (a) Nafion, (b) Aquivion, and (c) the 3M-Ionomer

PFSA membranes are most likely the immediate consequences of the PTFE backbone to pack in an ordered way (similar to the crystallization of pure PTFE (Teflon)), the high persistence length of the backbone, and the fact that the polymeric structure combines the extreme hydrophobicity of the backbone with the extreme hydrophilicity of the superacidic −SO3H group (note that the latter is related to the electron withdrawal property of the adjacent CF2 group). Even at low RH, a significant number of water molecules are absorbed at the −SO3H groups (the amount of water is usually expressed as hydration number λ = [H2O]/[−SO3H]), and a hydrophobic/ hydrophilic phase separated morphology develops spontaneously. The absorbed water tends to strongly bind together forming a continuous aqueous domain in which not only proton and water transport but also transport of dissolved gases (especially oxygen) take place. The accessible polymer conformations constrain the separation to the nanoscale as evidenced by the appearance of an “ionomer peak” in smallangle X-ray scattering (SAXS) patterns, as first observed by Gierke et al.14 Nanomorphology. Surprisingly, details of the nanomorphology of PFSAs were under debate for a long time until a recent work by Schmidt-Rohr and Chen who obtained seemingly the first quantitative picture of Nafion’s morphology from simulating a previously published small-angle X-ray scattering (SAXS) pattern.15 The so-called “parallel cylinder model” is now widely accepted and apparently supported by a recent NMR study on elongated membranes.16 The key feature of this microstructure is inverted micelle cylinders with large diameters even at low water contents (e.g., 2.4 nm for a water volume

Figure 3. Linear scaling of structural correlation length d with the polymer volume fraction Φpolymer as obtained from small-angle X-ray scattering (SAXS) of Nafion17 (inset), diverse sulfonated poly-etherketones (S-PEK), and a fully sulfonated poly-phenylene-sulfone (S220) which indicate locally flat morphologies as illustrated for a highly swollen situation (reprinted with permission; copyright 2013 WileyVCH).

nanomorphology may be different (in the case of Nafion, there may be water films between polymer ribbons, as suggested by the Grenoble group15), but in any case, the water films are relatively narrow (0.3−2 nm) which has important implications for transport, the visco-elastic behavior at high T and low humidification, and the evolution of the nanomorphology with T and RH. In long side chain PFSAs with high IEC, the hydrophobic/ hydrophilic separation is very robust; i.e., it is still observed at 364



Perspective

very high T (up to T = 160 °C) provided RH is not too low.20 The observation that this morphological feature tends to decay at low RH (the ionomer peak is almost lost at T = 160 °C and RH = 10%) is consistent with electrostatics driving the formation of slightly ordered flat morphologies. It is expected that, with decreasing RH and increasing T, which dramatically reduces the dielectric constant of water and therefore its solvating properties, dissociation of the ionomer, preceding electrostatic cross-linking, is reduced. A short side chain (SSC) Aquivion membrane with the same ion exchange capacity as Nafion actually loses its structural correlations at a lower temperature around T = 100 °C for RH = 10%,20 which points toward other competing interactions involved in the formation of the nanomorphology, as will be explained below. Another typical structural feature of all PFSA membranes is a certain degree of crystallinity arising from PTFE backbone ordering. Of course, this “crystallinity” is more pronounced for SSC PFSAs in which the properties are more dominated by the backbone and continuously decreases with increasing IEC, i.e., with increasing side chain concentration. The internal structure of the ordered parts shows up as a correlation peak in the WAXS (wide-angle X-ray scattering) regime,21 and the structural correlation between these parts is indicated by the so-called “matrix knee” which follows the ionomer peak toward lower q corresponding to correlation lengths around 15 nm. This feature appears to be quite robust compared to the very low q part of the SAXS patterns. It is clear that even on these large scales (>30 nm) there is still some structure, but the corresponding features are strongly affected by the sample history (including swelling, deswelling, aging, stretching, and pressing). A q−1 slope, as reported by the Grenoble group,22 indicates the existence of elongated objects, but the observation of slopes ranging from 0 to ∼3.517 (depending on the sample pretreatment and aging) strongly suggests that the PFSA structure on this length scale (∼30−100 nm) is subject to severe changes. These changes also seem to include the onset of the so-called “ultrasmall angle upturn”, a distinct increase of the scattered intensity at very low q.15,22 Frankly speaking, there is still no satisfactory understanding of the PFSA structure neither on the molecular (subnanometer) scale nor on very large scales. There is only reasonable evidence for a few features on the nanometer scale (see above). Fortunately, these features are those most relevant for understanding transport properties and, to some extent, also the visco-elastic behavior of PFSA membranes. Visco-Eleastic Properties. Apart from common stress/strain measurements, the visco-elastic properties of PFSA membranes are generally examined by complex dynamical mechanical analysis (DMA). There are actually two major trends in the evolution of the storage modulus (real part of the elastic constant): For relative humidity values in the range RH ∼ 10−100%, water is acting as a plasticizer; i.e., the storage modulus monotonically decreases with increasing hydration level, and at any relative humidity, there is a severe decay of the storage modulus with temperature up to T ∼ 140 °C (Figure 4). Surprisingly, the latter is not visible in the SAXS regime where T dependent scattering data are available (d = 1 − 30 nm). For a given RH, the SAXS patterns are virtually independent of T. Only at significantly higher T and low RH does the intensity of the ionomer peak decrease,20 and this perfectly correlates with another feature in the visco-elastic behavior. In this regime (high T, low RH), water actually stif fens the membrane (Figure 4), and the decay

Figure 4. Complex visco-elastic constants expressed as storage modulus E′ and loss tan δ for Nafion as a function of temperature T and given relative humidities RH.27 Note that the maxima of tan δ occur at higher T than the decays of the storage modulus E′ indicating the presence of different interactions controlling the visco-elastic properties (see text). The inset shows a comparison with the storage modulus of two short side chain PFSA membranes (Dow) recorded at a constant water pressure pH2O = 16 hPa25 (reprinted with permission; copyright 2013 Elsevier).

of the corresponding interaction shows up as a maximum of the elastic loss (tan δ in Figure 4) shifting toward higher T with increasing RH. The interaction, fading away in this T, RH regime, is obviously identical to the interaction driving the formation of the locally flat nanomorphology (see above) which decays in the same regime. This presumably electrostatic cross-linking only accounts for a minor part of the total storage modulus at lower T, and the decay of the modulus already visible at room temperature is most likely associated with the decay of another interaction. This cannot be the physical crosslinking through the crystalline parts of PFSA membranes, because the corresponding structural correlation (matrix peak in the SAXS) remains unchanged far beyond T = 100 °C,20 and also DSC (differential scanning calorimetry) shows an onset of “f usion” of the ordered parts only above T = 200 °C.23,24 There are actually two features observed in the DSC trace, and both have been interpreted as endothermal fusion. While this is undisputed for the distinct peak in the T range of PTFE melting (>300 °C) only visible for low IEC PFSAs, the second feature just above T = 200 °C looks more like indicating a glass transition. If this holds true, the corresponding glass transition temperature for the hydrated ionic form is expected to be lower than for the PFSA precursors (containing SO2F as nonionic precursory of the sulfonic acid group) used for the DSC measurements. The decay of the elastic modulus with temperature observed experimentally (Figure 4) is therefore most likely associated with a glass transition of the amorphous part. The fact that this shifts to higher T by ∼30 K for SSC PFSA compared to Nafion (inset Figure 4) suggests disintegration of the hydrophobic backbone interaction in this regime. In short: apart from electrostatic interactions driving the nanomorphology and leading to some ionic cross-linking, hydrophobic backbone interaction (probably in the amorphous phase) seems to govern the overall visco-elastic properties. 365



Perspective

Hydration Behavior. The hydration of PFSA membranes affects virtually all properties: transport coefficients of ions and water and even the permeation of dissolved gases generally increase with the level of hydration,25 and the selectivity for exchanging cations or anions (perm-selectivity) is lost while the elastic properties usually deteriorate with increasing hydration (see above). These dependences therefore critically control their performance not only in PEM fuel cells but also in redoxflow batteries, as will be discussed later. Apparently, the main force driving water into PFSA membranes is the extreme hydrophilicity of the super-acidic sulfonic acid group (−SO3H). Hydration is exothermal, and the heat of hydration decreases with increasing water uptake; in the case of Nafion the hydration enthalpy has been measured directly26 by calorimetry showing that about six water molecules per sulfonic acid group are absorbed exothermally. Accordingly, the water uptake in this regime is expected to decrease with increasing T, which recently has been proven experimentally (Figure 5a).27 From this decrease as a function of RH, the heat of hydration as a function of λ has been obtained and found to be close to that of sulfuric acid. The fact that hydration isotherms for most sulfonic acid functionalized ionomers and polyelectrolytes fall into a very narrow range for low RH, corresponding to water uptakes of λ < 6, clearly supports the assumption that hydration in this regime is governed by the hydrophilicity of the sulfonic acid group. The observed small variations may not only result from small differences of the acidity but also from hydration entropies, which depend on the degree of water dispersion. In fact, Nafion shows the highest water uptake in this regime (low RH), which correlates with the fact that Nafion undergoes a very pronounced hydrophobic/hydrophilic separation leading to disordered bulk-like water structures even at low water content. Apart from the exothermally absorbed water, which is involved in the solvation and separation of acidic protons and their conjugated base (−SO3−), there is much more water taken up by PFSA membranes at high RH (up to λ ∼ 20, Figure 5). The corresponding driving force has been suggested to be the entropy increase associated with the dilution of the protonic charge carriers within the hydrated hydrophilic domain.27 This is nothing but osmosis, and considering the high concentration of hydrated protons within the hydrophilic domain, the osmotic pressure may reach values limited by the corresponding counter-pressure built up within the polymer matrix. With the storage modulus shown in Figure 4, internal pressures up to 12 MPa (120 bar) have been calculated. While the related equilibrium water content is very close to the experimental one at room temperature, the severe decay of the elastic modulus with T leads one to expect a decreasing internal pressure and accordingly a severe increase of hydration (swelling) with T especially at high water activity (RH). Indeed, there is some increase in hydration observed, but this increase remains below what is expected from the significant decrease of the bulk modulus. This is true in the vapor phase, i.e., at RH below 100%; once the dew point is reached, and water condenses on the membrane surface and water uptake abruptly increases (Figure 5b). Observing two different water contents at the same water activity apparently violates the Gibbs phase rule which is why this phenomenon is known as “Schroeder’s paradox”. But this may be resolved by introducing the surface tension as a free parameter which has been done in various implicit and explicit ways.28−30 Recently we have suggested the formation of an “extended layered surface skin”, which is very

Figure 5. Hydration behavior of Nafion: (a) water uptake in terms of hydration number λ = [H2O]/[−SO3H] as a function of temperature T and relative humidity RH; the hydration number in water is given for comparison. (b) The data at T = 90 °C are shown as a hydration isotherm and proposed morphologies just below the dew point of water and in contact with water are schematically illustrated (see text) (reprinted with permission; copyright 2013 Elsevier).27

tough parallel and little permeable to water normal to the membrane surface. The surfaces of PFSA membranes are wellknown to be hydrophobic in the vapor phase, and the formation of such an extended surface skin may start at the hydrophobic membrane surface under the same driving forces which govern the formation of Nafion’s weakly ordered layered nanomorphology (see above). As a simple consequence of the boundary condition (flat hydrophobic surface), the periodic water “films” may arrange parallel to the membrane surface in the near surface region. There is experimental evidence for the formation of such a highly anisotropic surface layer from gazing incident SAXS,31 and a thickness of ∼60 nm is estimated from a study of very thin Nafion layers.32 This “skin” may then sustain a higher inner pressure than expected from the bulk modulus, and changes of the anisotropic “skin” structure in contact with liquid water are suggested as a means to release most of the internal pressure (insets Figure 5b). This may lead to a decrease of the chemical potential of water and therefore to an 366



Perspective

around λ = 6. Of course, the macroscopic water diffusion coefficient at this water content (∼ 10−6 cm2 s−1) is still more than an order of magnitude smaller than that of bulk water or acidic aqueous solutions,34 but this is essentially due to the fact that diffusion is geometrically confined to the volume fraction occupied by the hydrated hydrophilic domain. In PFSA membranes, especially in the long side chain variety Nafion, this percolation effect is relatively small, indicating high connectivity in the aqueous domain. For λ > 6, the water diffusion coefficient approaches the bulk water diffusion coefficient (2.2 × 10−5 cm2 s−1) with increasing water volume fraction in an almost linear way (in the double logarithmic plot (Figure 7) the corresponding slope is close to 1.2). As

uptake of additional water, taking the system back to local thermodynamic equilibrium. It is also worth mentioning that the reorganization of the membrane surface dramatically changes the water exchange rate with the environment; in contact with water, PFSA membranes equilibrate about 2−3 orders of magnitude faster than in humid vapor.27 Transport Properties (Water Transport and Proton Conductivity). Since the transport properties of Nafion have been discussed extensively in another review,25 here the major features are summarized with a focus on a few recent insights which have altered our general understanding of transport in PFSA membranes. Most of the transport takes place within the hydrated hydrophilic domain of PFSA membranes, and the major parameters to be considered are the morphology, the ion exchange capacity, and the hydration behavior, which is related to the viscoelastic properties (see above). The local dynamics of water exothermally absorbed at low RH is retarded to some extent because this water is essentially involved in ion solvation. But, interestingly, the sulfonic acid anion (−SO3−) and the hydrated proton start to separate far before the ionic solvation shells are completed. For the model system methyl sulfonic acid (CH3SO3H)−water (H2O), 85% of the acid is found to be dissociated at a water content as low as λ = 2.33 At this hydration level, the ion concentration of the model system passes through a maximum, and the striking observation that the diffusion coefficients of all species of the system pass through a minimum at the same water content has been explained by an ordering of the incompletely screened ions in order to minimize the electrostatic energy. Any molecular and ionic transport then requires some extra thermally activated disordering, and the overall transport rate is reduced. Translating this observation to the transport within hydrated hydrophilic domains of PFSA membranes, some local proton conductivity is expected even at such low water levels, but with higher activation enthalpies than those of water diffusion and proton mobility in pure water. Activation enthalpies for local diffusion and conduction processes are not available, but the anticipated behavior is also visible in the T dependence of the macroscopic water diffusion and proton conductivity (Figure 6): the activation enthalpy of water diffusion decreases with increasing water content approaching the value of pure water at

Figure 7. Proton and water transport coefficients as a function of water volume fraction Φwater: DH2O (water tracer diffusion), Dσ (proton conductivity diffusion coefficient), DFick (diffusion in water concentration gradient), and Dp (pressure driven permeation diffusion coefficient).25

expected, the conductivity diffusion coefficient Dσ (proton mobility) closely follows the water diffusion coefficient especially at low water content (Figure 7); i.e., as in acidic aqueous solutions, proton conductivity is brought about by a vehicle mechanism, the cooperative diffusion of protonated and unprotonated molecules (here water).35,36 With increasing water content, there is an additional conductivity contribution from structure dif f usion the prevailing proton conduction mechanism in pure water and dilute acids.37−39 Since this mechanism comprises intermolecular proton transfer, the terms “proton hopping” and “Grotthuss mechanism” are also used. But use of the term “structure dif f usion” may be more expedient because the protonic charge carrier just follows the hydrogen bond pattern (structure) of the solvent (water) around the excess proton diffusing by hydrogen bond breaking and forming processes.39 In other words, this is a solvent driven process involving the dynamics of many molecules displacing the position of the protonic defect by just one molecular separation. In particular, there is no indication for the formation of extended “Grotthuss chains” as still suggested in some textbooks. It also should be mentioned that solvent and acidic protons interchange their identity on a sub-picosecond time scale; i.e., there are no physically fast protons (for aqueous HCl, a maximum effects of 4% has been confirmed experimentally34).

Figure 6. Activation enthalpy Ea of water diffusion and proton conductivity of Nafion as a function of hydration number λ:53 with increasing λ, the corresponding values for pure water are approached. 367



Perspective

It is the very nature of structure dif f usion that it is very sensitive to symmetry reduction.39 In acidic solutions, this is caused by the effect of neighboring ions biasing the hydrogen bonds within the solvation structure of the hydrated acidic proton; within the separated morphology of an ionomer, such as PFSAs, also confinement effects reduce symmetry and therefore the rate of structure diffusion. In PFSA membranes, confinement effects are relatively small, i.e., the decrease of structure dif f usion with decreasing water content resembles that of acidic aqueous solutions (e.g., ref 34). When it comes to understanding water transport and the formation of water concentration profiles within PFSA membranes in running fuel cells, the effect of internal pressure gradients (see section on hydration) on molecular transport must be considered as well. In PFSA membranes, the internal pressure gradient is a driving force for hydrodynamic water flow, which is a fundamentally different mechanism than the diffusional transport processes discussed above. Water flows in a collective way when the water structures are sufficiently wide, which is the case in PFSA membranes at high levels of hydration. This leads to the remarkable situation that collective pressure driven water transport increases substantially while Fickian water diffusion decreases close to the dew point of water (Figure 7) because of the vanishing chemical driving force (chemical water potential gradient). The interested reader may find the corresponding data in ref 25 and in ref 40 as a way to implement these two different water transport processes (diffusional and hydrodynamic) into the transport equations. Interestingly, the dramatic onset of hydrodynamic water transport is also clearly visible in the so-called electroosmotic water drag, defined as the number of water molecules per proton dragged through the membrane as a consequence of a protonic current. Of course, this is a key feature for any electrochemical application involving high current densities, which is true not only for fuel cells but also for redox-flow batteries. In the case of Nafion, the electroosmotic drag coefficient at room temperature is in the range of 2 below the dew point of water, but at higher swelling (Φwater > 30−40%), e.g., at higher T or in the presence of methanol,41 the drag coefficient K increases dramatically (Figure 8). This transition actually corresponds to a width of the flat water films around 1.5 nm17 which seems to correspond to what is called “slip” in fluid dynamics.42 In today’s PEM-fuel cells, operating close to the dew point of water, hydrodynamic water transport is of paramount importance for describing water distribution within PFSA membranes. Apart from the boundary conditions (gas humidification, membrane surface hydrophilicity), this is essentially determined by the electroosmotic water drag from the anode to the cathode and the pressure driven water permeation (Figure 7) in the opposite direction.40 Within the family of PFSA membranes, the hydrophobic/ hydrophilic separation of the SSC varieties is slightly less pronounced which leads to slightly lower hydrodynamic water transport for a given water volume fraction,40 but practical SSC membranes usually have higher IECs than Nafion with higher water uptake at a given RH, undoing this advantage. Since direct methanol f uel cells (DMFC) is still of interest in PEM fuel cell technology, it must be mentioned that the abovediscussed properties severely change in the presence of alcohols such as methanol. With their hydrophobic alkyl residue and their polar OH group, alcohols tend to behave like surfactants; i.e., they mediate interaction of the PTFE backbone with the

Figure 8. Electroosmotic drag coefficient K as a function of water content at room temperature for Nafion, the polyelectrolytes S-220 and S-360 (fully and half-sulfonated polyphenylene sulfone), and the multiblock-copolymer SU14−FS15 (see Figure 9)59 (reprinted with permission; copyright 2012 Wiley-VCH). Note that Nafion has a wellseparated hydrophobic/hydrophilic nanomorphology whereas the hydration water in the two polyelectrolytes is highly dispersed (see text).

hydration water leading to more swelling.41 Severe increase of hydrodynamic solvent (water and methanol) and gas transport is an immediate consequence,25 which is frequently called “crossover”. The tracer diffusion coefficient of methanol is only a factor of 2 lower compared to the water diffusion coefficient25 and electroosmotic drag coefficients are virtually identical,41 but the presence of any alcohol dramatically reduces dissociation of the sulfonic acid group (−SO3H) and therefore also proton conductivity.25,41 Stability. Frequent swelling and deswelling, severe drying close to sealings, fuel cell stack mounting pressures of more than 2 MPa, and the extreme chemical conditions in PEM fuel cells require the proton conducting polymer membrane to be mechanically and chemically robust. As opposed to many other membrane types, PFSA membranes are indeed mechanically very robust, although their modulus is comparatively low. This has to do with their extreme fracture toughness (elongation to break) characteristic for ionomers with PTFE backbone. The severe softening at high T and RH (Figure 4), however, clearly limits the operation window of PFSA membranes, and the higher glass transition temperature of SSC PFSAs (see above and inset of Figure 4) is therefore a distinct advantage. When it comes to chemical durability, SSC PFSAs seem to have a slight advantage over the long side chain variety Nafion.43 This especially applies to the degradation of the side chains, which starts from the C−S bond and is significantly faster than backbone degradation.44 Prior to this 19F-NMR investigation of membranes subjected to fuel cell test protocols, the so-called unzipping reaction starting from COOH groups terminating the backbone had been identified;45,46 fortunately, this reaction can be suppressed by ensuring that the main chains are terminated by −CF3 groups. The occurrence of the observed reactions is actually the consequence of the presence of peroxo- and/or hydroxo-radicals which form on the surface of the electrocatalysts in the presence of both oxygen (O2) and hydrogen (H2).47 Therefore, gas permeability has to be 368



Perspective

statistically sulfonated polyarylenes than for PFSAs.25,41 A closer inspection of the diffusional transport coefficients including their T-dependence shows that both transport on the nanoscale and long-range diffusion are retarded in sulfonated polyarylenes compared to PFSA membranes. This clearly suggests that the aqueous nanostructures are not only narrower but also less connected with a higher tortuosity on larger scales. As expected from this nanomorphology, the onset of excessive electroosmotic water drag is at distinctly higher water volume fractions. The interested reader may find a comprehensive set of transport data in ref 25. Statistically sulfonated polyarylenes are ionomers just as PFSAs, and that is why their visco-elastic properties are still determined by the nature of their polymer backbones. In contrast to low IEC PFSAs (see above), these are generally amorphous and swell more in water which makes them very soft under wet conditions with a low fracture toughness (small elongation to break), while they are brittle in the dry state. Blending with a small concentration of compatible unsulfonated high molecular weight polymer actually leads to a distinct improvement of the visco-elastic properties without changing the transport properties too much.49 The mode of sulfonation has actually severe implications on the chemical durability which is critical with respect to acid− base and redox reactions. For electrophilic sulfonation, the rule of the thumb “easy on−easy off” applies to the hydrolytic desulfonation as described and explained in ref 52. Not only can electron rich polyarylenes be easily sulfonated and desulfonated, but they are also very susceptible against oxidative attack, i.e., through H2O2 and even more so by oxidizing radicals. The fact that several electrophilically sulfonated polyarylenes survived in operating fuel cells for more than 1000 h has probably to do with the limited operation temperature (T < 80 °C) and the fact that the highly reduced gas crossover compared to PFSA membranes efficiently prevents radical formation. Sulfonated Polyarylenes with Block-Structure. The above qualitative analysis naturally guides the way toward the current approaches for overcoming the main disadvantages of hydrocarbon membranes, namely, the steep conductivity decrease with decreasing level of hydration, the susceptibility toward hydrolytic and oxidative attack, and the poor mechanical stability, especially brittleness in the dry state. The most effective way to increase the conductivity at low RH is to increase the IEC because this measure increases the charge carrier concentration and the diffusion coefficient of water, greatly influencing the ionic mobility. The latter has mostly to do with the fact that the number of water molecules absorbed per ionic group is approximately identical for all ionomers and polyelectrolytes in this regime (see above); i.e., the water uptake is approximately proportional to the IEC. The percolation within the aqueous domain, however, increases in a highly nonlinear fashion. As a result, any increase in IEC dramatically increases the conductivity at low RH where water percolation is still low. This has been demonstrated by the extremely high conductivity of fully sulfonated polyphenylenesulfone53,54 (Figure 2) and poly-para-phenylene.55 According to the IUPAC nomenclature,2 these are polyelectrolytes with saltlike properties: they are extremely brittle in the dry state and soluble in water. The high concentration of ionic groups makes any phase separation redundant; i.e., the hydration water is highly dispersed within a morphology which is well organized

included into the stability considerations. At high hydration levels, the water in PFSA membranes dissolves comparatively high concentrations of gases which are transported with the hydration water. This actually explains why PFSA degradation is highest under OCV (open circuit potential) conditions when no protonic current and therefore no electroosmotic water flux from the anode to the cathode prevents oxygen from permeating from the cathode to the anode side. Just by considering the chemical structure of PFSA membranes, one may also anticipate breaking of ether linkages within the side chains through acidic attack, but to the best of my knowledge there is no convincing experimental indication for this. Long-term hydrothermal aging experiments at T = 80 °C and 100% RH show a continuous decrease of the IEC,48 but this can be completely reversed through the usual standardization procedure (e.g., 2 h in 1 M of a strong acid at T = 80 °C). The authors actually found spectroscopic indication for condensation reactions between sulfonic acid groups, but it could well be that slow structural relaxation driven by residual electrostatic interaction (see above) may also have prevented the membrane from rapidly exchanging ions in the titration process used to determine the IEC. 3.1.2. Hydrocarbon versus PFSA Membranes. Since the early 1990s, there has been continuously high interest in hydrocarbon membranes as potential alternatives for the well established PFSA membranes. Originally research and engineering activities were driven by the high costs of PFSA membranes and the putative environmental problems associated with the use of perfluorinated polymers in mass products. Later, it became clear that hydrocarbon membranes have distinctly different properties;49 i.e., they may be even more suitable than PFSAs for certain applications. Statistically Sulfonated Polyarylenes. Because of their superior stability, initially polyarylenes have been generally chosen as starting polymers for polymer analogous sulfonation. Direct electrophilic sulfonation, e.g., in fuming sulfuric acid, was very popular, but soon also nucleophilic substitution routes were used for the sulfonation of electron poor base polymers. The interested reader may refer to the review by Rikukawa and Sanui,50 a chapter in the Fuel Cell Handbook,41 and the excellent recent review by Hickner et al.51 For these preparation routes, sulfonation is essentially random along the polymer chain as it is in PFSAs, but the properties of the cast membranes differ from these of PFSA membranes in a characteristic way.49 For similar volume densities of sulfonic acid groups (e.g., a sulfonated polyarylene-ether-ketone with an IEC of 1.4 mequiv g−1 corresponds to Nafion with an IEC of 0.9 mequiv g−1) and a given water volume fraction, SAXS patterns of sulfonated polyarylenes and PFSA membranes are similar in the regime of the ionomer peak, again supporting the hypothesis that some residual electrostatic interaction between the hydrated ions drives the nanomorphology (see above). The influence of the backbone properties still shows up as a slight but systematic increase of the scattering intensity within the Porod regime and a small shift of the ionomer peak toward higher q, along with some broadening.49 This points toward a less pronounced hydrophobic/hydrophilic separation which naturally explains the significantly lower permeation coefficients for water and dissolved gases. The effect on the diffusional transport (water diffusion and proton conductivity) is insignificant at high hydration levels, but with decreasing hydration the decay of water diffusion and proton conductivity is more severe for 369



Perspective

on a low nanometer scale, as evidenced by distinct X-ray diffraction patterns.53 Nevertheless, they can be used as a constituent of more complex polymeric structures. If the polyelectrolyte part separates to form a continuous domain, the conductivity can be expected to be higher than for similar materials with a random distribution of sulfonic acid group simply because of the highly nonlinear increase of the local conductivity with local IEC (see above). An elegant way to form such a material is through so-called block-co-polymers, an approach pioneered by several groups.56−60 Here, typically segments of a highly sulfonated polyarylene with a defined molecular weight of a few thousand g·mol−1 are combined with unsulfonated segments of a similar defined molecular weight through a coupling group so as to form di-, tri-, or even multi-block copolymers. Of course, such supramolecular structures may phase separate in a defined way when cast from solutions of polar aprotic solvents preferentially as thin membranes. This phase separation is then typically on a scale of a few tens of nanometers and should not be confused with the separation observed in ionomers on smaller scales between water and polymer. In fact, the latter is observed as the internal structure of the hydrophilic domain with typically much lower correlation lengths than observed for ionomers (inset Figure 9).53

Figure 10. Proton conductivity of a multiblock-copolymer (SU14− FS15) containing S-220 segments forming a proton conducting domain at a water pressure of pH2O = 105 Pa (see also Figure 9).59 The conductivity of pure S-220 and Nafion is shown for comparison.

carriers is subjected to higher activation enthalpies,59 i.e., such membranes are particularly suitable for high-T operation. Needless to say, polyelectrolytes do not form hydrophobic surface skins, and this also seems to hold for the corresponding multiblock copolymers. Compared to PFSA membranes, they therefore take up more water at high relative humidity and the exchange of water with the vapor phase is generally fast (see section on hydration behavior of PFSA membranes). No compromise can be made with respect to hydrolytic and oxidative stability of the sulfonated segments because, in all PEM fuel cells, they are in intimate contact with water, oxidizing gases, and radicals. A very useful way toward stable sulfonated segments is through the use of presulfonated monomers in the polymerization reaction, an approach which has been extensively promoted by the McGrath group.51 This allows one to sulfonate electron rich phenyl rings in the usual way before reducing the electron density, and therefore the reactivity, by introducing electron acceptor groups such as −SO2− by oxidation of sulfide groups.52,54,59,61 3.1.3. Exploring the Limits of Sulfonic Acid Functionalized Systems. The operation of PEM fuel cells at high temperature and low humidification is a critical issue, and therefore understanding the physicochemical conductivity limits of sulfonic acid based systems under these conditions is of particular interest. Since the values of the hydration number λ of sulfonic acid functionalized systems fall into a very narrow range for low water activities (RH < 65% corresponding to λ ∼ 6), the transport coefficients as a function of λ are informative parameters. In this regime, the hydration water is involved in ion solvation, retarding both water diffusion and proton conductivity (see above) . This also becomes visible as an increase of the corresponding activation enthalpies with decreasing hydration number λ (see also Figure 6). Comparing activation enthalpies of systems with higher water dispersion (e.g., SSC PFSAs40 and sulfonated polyphenylene-sulfones52) with the corresponding values for Nafion clearly suggests that, apart from chemical interaction (solvation), water confinement effects must be considered as well. For a given hydration number, the activation enthalpies for water and proton transport, which are indicative of the local dynamics, exhibit the lowest values for Nafionthe ionomer with the most

Figure 9. Morphology of a multiblock-copolymer (SU14−FS15 consisting of blocks of fully sulfonated phenylene sulfones and blocks of phenylene ether sulfones with an overall IEC of 1.6 mequiv/g) as observed by transmission electron microscopy.59 The inset shows the corresponding SAXS pattern compared to this of S-220 (see text) (reprinted with permission; copryight 2012 Wiley-VCH).

If the phase separation is well developed, as the one shown in Figure 9, the properties of the polyelectrolyte part are locally preserved. Such membranes may then combine in a unique way some of the high conductivity of the hydrated domain and the low hydrodynamic water transport with the mechanical properties of the water free preferentially elastic domain. Of course, this is only the case if the morphology is bicontinuous. Then acceptable compromises between transport and mechanical properties may still allow for proton conductivities higher than that of Nafion (Figure 10). As in the case of pure polyelectrolytes, the transport of water and protonic charge 370



Perspective

hydrocarbon polymers is of paramount importance, since this provides an additional means to increase the local density of ionic groups.63 At what density of ionic groups emerging polyelectrolyte effects start to reduce dissociation has yet to be clarified, but in any case, the ionic groups must be concentrated in a well percolating volume increment (domain). Beyond fulfilling the conductivity requirements, membranes must be flexible so as to adapt to changes of the water content. In particular, the formation of a free volume upon drying has to be avoided, since this dramatically reduces both the local and long-range transport coefficients. On the other hand the maximum swelling of such high IEC materials has to be limited through a corresponding toughness of the polymer matrix. If the maximum swelling were constrained to λ = 5, the hydration level and proton conductivity may even become independent of RH for RH > 50%. For such materials, the detrimental effects of swelling and deswelling on the proton conductivity are expected to be small while the conductivity at high RH remains reasonably high. It goes without saying that optimum membranes must be morphologically and chemically stable. As part of membrane electrode assemblies, they experience an external uniaxial pressure in the range 2−5 MPa within typical PEM fuel cell stacks, while fluctuations of the internal swelling pressure in space and time lead to additional strain. The response of the membranes to the related forces is not only controlled by their storage modulus but also by their fracture toughness (elongation to break). Although PFSA membranes are quite soft at T = 90 °C (see Figure 4), their enormous elongation to break (∼150%) prevents mechanical failure. On the other hand, most hydrocarbon membranes have a significantly higher storage modulus under these conditions, provided hydration is not too high. But for low hydration levels, elongation to break is very small, i.e., the membranes are brittle; high hydration levels lead to softening with no significant increase of elongation to break. For this reason the most frequent failure mode of hydrocarbon membranes is the formation of cracks and pinholes under the conditions of a running fuel cell. The increase of the membrane fracture toughness over a large RH range must therefore be one of the targets of future membrane developments. The issue of chemical stability appears to be also quite complex. Since there is nothing like “thermodynamic stability” of organic membranes, one may rather use the term “durability” referring to the reactivity of a membrane material under specific conditions. As pointed out above, the most important types of reaction are the attack by oxo and hydroxo radicals and desulfonation through hydrolytic cleavage of the C−S bond. Because the intrinsic susceptibility of this bond does not vary significantly with the chemical environment, steric shielding seems to be the most efficient way to prevent this bond from being attacked. This is probably the reason why short side chain PFSA membranes are chemically more robust than the long side chain ionomer Nafion.43,44 Interestingly, the formation of radicals depends on a membrane property, namely, the gas crossover, which is significantly higher for PFSA compared to hydrocarbon membranes. Therefore, hydrocarbon membranes have the potential to perform stably in PEM fuel cells, even though virtually all of them fail Fenton’s test. In particular, membranes with locally high IECs desirable for obtaining high proton conductivity (see above) generally have a high dispersion of water and therefore a low dissolution and transport of gases.

developed phase separation. On the other hand, water dispersion depends on the nature of the backbone and the IEC; i.e., it is higher for PFSA than for hydrocarbon membranes49 and generally increases with increasing IEC.40,53 A high IEC actually leads to a higher concentration of protonic charge carriers and a high water uptake in terms of water volume fraction, which is associated with a more efficient percolation of the hydrated hydrophilic domain. These are actually reasons for the very high proton conductivity of high IEC SSC PFSAs such as DOW 858 (see Figure 2 and ref 40) or the 3M membrane62 and explains why high IEC hydrocarbons such as fully sulfonated-polyphenylenesulfone (S-220) show proton conductivities comparable to that of Nafion at room temperature.53 At this temperature, the retardation of transport on the local scale is compensated by the better percolation on larger scales and the higher charge carrier concentration. Because of the higher activation enthalpy, high IEC hydrocarbon systems may show increasingly higher proton conductivity than Nafion with increasing T. Since conductivity increase with IEC is highly nonlinear, the accumulation of ionic groups into a distinct well percolating phase is a suitable strategy to maximize proton conductivity. The physicochemical limit can then be estimated by considering the water diffusion coefficient and proton conductivity for aqueous solutions of small sulfonated molecules such as methyl-sulfonic acid (MSA).33 Here, percolation and confinement effects are virtually nonexistant, while the local dynamics is expected to resemble the situation in well phase separated membranes. The comparison of the conductivity of the MSA−H2O system with that of Nafion reveals a conductivity exceeding the conductivity of Nafion at λ = 3 (RH ∼ 30%) by a factor of ∼40 (Figure 11). For real

Figure 11. Ionic conductivity of methyl sulfonic acid as a function of hydration number λ.33 Note that the conductivity at low hydration number is significantly higher than that of Nafion (reprinted with permission; copyright 2010 Elsevier).

membranes with nonconducting hydrophobic volume increments and some residual confinement effects, a possible conductivity increase by a factor of ∼5 appears to be plausible. Practically, high conductivities may be achieved by using multiblock architectures (see above), by reinforcing high IEC PFSAs62 or hydrocarbon membranes, e.g., through appropriate blending, or by dispersing nonsoluble high IEC nanoparticles. In this context, preparation routes which allow the placement of more than one sulfonic acid group on the phenyl rings of 371



Perspective

When it comes to the desulfonation reaction, the general observation is: “easy on easy off”; i.e., polymers which are easily sulfonated are easily desulfonated at high water activities. Of course, this has to do with the free energy of the common transition state of both reaction pathways, and in the case of polyphenylenes this can be increased by reducing the electron density of the sulfonated phenyl ring.52 The use of sulfone (−SO2−) as a strong acceptor group in ortho position to the sulfonic acid group has not only been proven to be particularly efficient in increasing the hydrolytic durability52 but additionally reduces the susceptibility against radical attack.64 3.2. Anion Exchange Membranes (AEM). For various reasons, anion exchange membranes are not simply mirror images of above-discussed cation exchange membranes. The most common anion exchanging groupsquaternized amines such as trimethyl ammonium (TMA), methyl-imidazolium, penta-methyl-guanidinium, and diazabicyclo[2,2,2]octane (DABCO) (see Scheme 2)are not as strong bases as the

Figure 12. Water diffusion and conductivity diffusion coefficients (DH2O, Dσ) as a function of water volume fraction Φwater for a noncross-linked poly(arylene sulfone) functionalized with TMA groups (FAA-3 supplied by Fumatech, see text).66

Scheme 2. Molecular Structures of Anion Exchanging Groups: (a) Tri-Methyl Ammonium (TMA), (b) Methyl Imidazolium, (c) Penta-Methyl Guanidinium, and (d) Diazabicyclo[2,2,2]octane (DABCO)

coefficient; i.e., counterion (OH−) condensation at the conjugated acid (TMA+) excludes an increasing part of the anions from being transported within the aqueous domain. It is important to note that, in this regime, the water uptake in terms of λ at a given RH is lower than for sulfonic acid functionalized systems which further reduces the conductivity. Despite these downsides, reported conductivities up to 6.8 × 10−2 S cm−1 at high RH are absolutely interesting for fuel cell applications65 for which a maximum power density of 823 mW cm−2 has been reported.67 The reason why anion exchange membranes are not yet established in PEM fuel cell technology is their intrinsic chemical instability and the above-mentioned susceptibility to CO2 contamination. More than 20 years ago, Bauer et al. already determined the reaction rates of nucleophilic substitution and Hofmann elimination of the amine group through OH− ions.68 In the first case, the attack is at the carbon next to the amine nitrogen, while in the second case the reaction is initiated by attack of a hydrogen in the β position to the leaving group (e.g., NR4+) (Scheme 3). In both cases, amine is released leaving behind an alcohol group in the first case or a CC double bond and a water molecule as a second leaving group in the case of Hofmann elimination. It also should be mentioned that even the commonly used polyarylene backbones may decompose by reacting with hydroxide ions.69 As

sulfonic acid group is an acid (as constituent of PFSAs the latter is even a superacid). Furthermore, the hydroxide ion (OH−), as the anion of choice for AEM fuel cell applications, not only tends to react with acidic gases, it also behaves as a strong nucleophile in general. The OH− conductivities reported in the early literature are significantly lower than the very high proton conductivities of sulfonic acid functionalized systems, but there is increasing evidence that this is the consequence of CO2 contamination leading to the conversion of OH− into carbonate (CO32−) and eventually bicarbonate species (HCO3−). With a careful exclusion of carbon dioxide, OH− conductivities close to the proton conductivities of PFSA membranes have recently been reported,65 and a quantitative determination of water diffusion and hydroxide conductivity of un-cross-linked polyarylene functionalized with quarternary ammonium groups (supplied by Funtech under the trade name FAA-3) reveals clear similarities with transport in proton exchange membranes (Figure 12).66 At intermediate water contents (Φwater ∼ 20−30%), the OH− conductivity diffusion coefficient Dσ follows the H2O diffusion coefficient; i.e., the ionic groups are highly dissociated and hydrated OH− diffuses at a similar rate as the water molecules. With increasing water content, the occurrence of OH− structure diffusion is evidenced by the fact that Dσ is exceeding the water diffusion coefficient approaching the value for pure aqueous solutions (5 × 10−5 cm2 s−1) at very high water contents. As expected from the lower rate of structure diffusion in pure water, the hydroxide mobility in this regime is about a factor of 2 lower than the mobility of protonic charge carriers in acidic membranes. At very low water contents (Φwater < 20 vol %, λ < 10), the moderate basicity of the anion exchanging group shows up as a severe decrease of Dσ compared to the water diffusion

Scheme 3. Splitting Off Amine Functional Groups through (a) Nucleophilic Substitution and (b) Hofmann Elimination68

372



Perspective

and that there is still some power generation in the presence of low levels of CO2, but at the moment anion exchange membranes are still far from being utilizable in PEM fuel cell technology (for up-to-date reviews see refs 78−80). At this stage, however, they have an immediate potential for redox-flow battery applications, as will be discussed later. 3.3. Adducts of Polybenzimidazole and Phosphoric Acid. A fundamentally different type of membranes with proton conductivities less dependent on external humidification takes advantage of the properties of phosphoric acid (H3PO4). As opposed to water, phosphoric acid is an intrinsic proton conductor as a consequence of its high degree of selfdissociation and a very high mobility of protonic charge carriers.81,82 In the nominally dry state, self-dissociation leads to the formation of a variety of charged species:

opposed to the desulfonation of cation exchange membranes which can be suppressed by acceptor substitution (see above), the opposite is true for anion exchange membranes, in which the presence of electron withdrawing groups (e.g., −SO2−, benzyl, −CN) increases the rate of reactions with nucleophiles. Approaches to stabilize the amine via charge delocalization (e.g., by using imidazolium, guanidinium) may not have been successful as they trade increased stability through delocalization with reduced stability from dramatically reduced steric shielding due to their planar geometry. 66 The rapid decomposition of polybenzimidazole (which has a larger delocalization range then imidazolium and guanidinium) in alkaline media70 strongly suggests that delocalization alone cannot sufficiently stabilize a quaternary amine.66 Currently, the most promising approach for stabilizing quaternary amines seems to be through the use of spacers coiling around the ammonium group, thereby achieving a high degree of steric shielding in a comparatively simple way.71 The spacer may either separate the amine from the polymer backbone or may be tethered to the amine with a methyl as the other terminal group. Currently, claims about base stable quaternary amines are abundant,72 but so far, the results could not be verified by independent studies.73−75 From this literature, current anion exchange membranes in their OH− form appear to be intrinsically durable only for temperatures not exceeding T ∼ 60 °C. Knowing that the polymeric environment may have some effect on the stability of the anion exchanging group, a preselection through testing small molecules yet appears to be an efficient strategy to search for suitable candidates. On the basis of the systematic work of Schwesinger et al.76 this has led to the identification of various phosphonium cations as a putatively base stable anion exchanging group.69 Apart from these chemical instability issues, the susceptibility to CO2 contamination seems to be a no-go for fuel cell applications in air. In this environment, any anion exchange membrane converts from its OH− into the HCO3− form within minutes66 leading to a dramatic decrease of the conductivity. This is not only the consequence of the approximately five times lower mobility of HCO3− compared to the rate of OH− structure diffusion in aqueous solutions; anion exchange membranes in the HCO3− form also take up less water at a given relative humidity.66 Interestingly, the oxygen reduction reaction taking place at the cathode produces OH− only in the absence of CO2. Otherwise, HCO3− or even CO32− may form, and the transport of these ions from the cathode to the anode may mediate an electrochemical reaction even in the presence of CO2. But this reaction is associated with a large anodic overpotential, which shows up as a severe voltage drop even at low current densities.77 At high current densities and CO2 levels < 1000 ppm, as is the case for operation in air (400 ppm), the production of OH− in the cathode reaction and the transport of OH− from the cathode to the anode leads to a displacement of carbonate species in the membrane, which is well-known as the so-called “self-purging effect”.77 But at such high current densities (>500 mA cm−2), the fuel cell efficiency is already down to about 25% which is too low for most types of applications. It should clearly be stated that the detrimental precipitation of solid carbonates characteristic for KOH based fuel cells is not an issue in fuel cells making use of anion exchange membranes

5H3PO4 ⇌ 2H4PO4 + + H 2PO4 − + H3O+ + H 2P2O7 2 −

The fact that their concentration (about 10 mol %81) and the mobility of protonic charge carriers are very high is a direct consequence of the phosphoric acid hydrogen bond network topology. Since this is key to the unique properties of phosphoric acid based proton conducting materials in general a brief explanation of these relations is given here: With three out of four oxygens being protonated, the phosphoric acid molecule (H3PO4) has a severe imbalance of the number of potential proton donors and acceptors. When phosphoric acid molecules condense to form a solid or a liquid the dominant intermolecular interaction is hydrogen bonding, indeed. Individual hydrogen bonds of the type O−H···O are actually very short ( 1.5 which correspond to O−H groups which are involved in hydrogen bonding as acceptor but only little as proton donor (Figure 13). This “frustrated” high energy state is important (i) as a precursor for the intrinsic condensation reaction, (ii) for the extreme hygroscopicity of phosphoric acid, and (iii) for the rapid proton conduction mechanism in the nominally dry state. This state is not far from being an implicit water molecule (Figure 13); i.e., completely donating the proton within the hydrogen bond to the “frustrated” O−H, which cannot release its covalently bonded proton, leads to the formation of a water molecule (H−O−H) and breaking of the O−P bond. Of course, this reaction goes along with the formation of a bond between this phosphorus and an unprotonated oxygen of another phosphate (formation of a diphosphate) and the transfer of an additional proton (preferentially a “frustrated” proton from another molecule) to the water molecule leading to the formation of a hydronium ion (H3O +). As a consequence, condensation products (e.g., diphosphates) and aqueous species always coexist within phosphoric acid, which has some water partial pressure even in the nominally dry state. At the typical operation temperature of a HT-PEM fuel cell (T ∼ 170 °C), this partial pressure corresponds to RH ∼ 1%.84 For 373



Perspective

interrupted and reversal of the polarization process inhibited. The charges remain separated and the protonic charge carriers diffuse in an almost uncorrelated way before two different defects start to attract and neutralize.82 This process leads to the proton conductivity of nominally dry phosphoric acid (H3PO4) which is very sensitive to any perturbation (see below). The only additive, which increases the proton conductivity, is water, which is an intrinsic constituent of phosphoric acid even in the nominally dry state (see above). The addition of small amounts of water does not change the principal conduction mechanism, but it appears to increase the rate of the elementary reactions involved.85 In the following, proton conducting adducts of phosphoric acid and polybenzimidazole (PBI) are discussed in the light of above-described mechanisms. Initially, PBI has been chosen because of its high concentration of basic nitrogen sites associated with the benzimidazole moiety interacting with acids either through complete proton transfer or just by hydrogen bonding.12 How much phosphoric acid is taken up by PBI very much depends on the kind of PBI and the formation process of the adduct. In fact, a huge number of different PBIs and other basic polymers have been reported to form adducts with phosphoric acid (for recent reviews, see refs 86−89). Because of its availability, initially, meta-PBI (Scheme 4a) was

Figure 13. Degree of oxygen protonation of a single H3PO4 molecule and phosphoric acid in the condensed hydrogen bonded liquid state as obtained from ab initio MD simulation.85 Two populations are clearly visible corresponding to characteristic bond patterns (insets). Note that the frustrated proton can be considered to be part of a water precursor (see text).

Scheme 4. Molecular Structures of Diverse PolyBenzimidazoles: (a) PBI, (b) AB-PBI, (c) iso-AB-PBI, and (d) PBI-OO

any higher humidity, phosphoric acid takes up more water; for any lower humidity, the concentration of condensation products increases. The driving force for the uptake of additional water, i.e., the reason for phosphoric acid’s extreme hygroscopicity, is probably the high hydrogen bond network “frustration” as well. Absorption of water molecules leads to a reduction of this “frustration” through proton transfer from the “frustrated” network to the absorbed water molecules, which actually blend into the hydrogen bond network by forming strong hydrogen bonds with the phosphate species.85 Apart from the condensation reaction and the formation of aqueous species, simple proton transfer between adjacent phosphoric acid molecules also leads to the formation of charged species (essentially H2PO4− and H4PO4+), which are most important for phosphoric acid’s high proton conductivity. This proton transfer is driven by the very weak electrostatic proton/proton coupling within the hydrogen bonded network with an otherwise almost random proton “rattling” dynamics.85 Once intermolecular proton transfer takes place, the hydrogen bond structure around the overprotonated phosphate (H4PO4+) slightly contracts, thus preparing the next proton transfer step. With some delay time, this leads to the formation of a polarized hydrogen bond chain, which may comprise up to about five phosphate moieties with a ∼20% probability.83 In contrast to the situation in water, this is energetically cheap because the proton displacements in the very short hydrogen bonds and the corresponding dipolar moments are small, and the dielectric response of the “solvent” environment, stabilizing the polarization, is very fast because of its protonic nature (Zundel polarizability). Of course, a chain of polarized hydrogen bonds (which may be termed Grotthuss chain83) loses its polarization by a reversal of the proton transfer events. In phosphoric acid, however, frustrated protons “attack” the negative partial charges appearing along the dipolar chain through forming hydrogen bonds; hence, the chain is

commonly used before other types of PBI such as AB-PBI (Scheme 4b) were considered as well. The higher density of basic imidazole groups of the latter leads to a higher uptake of phosphoric acid for given conditions,87 which is reported to be particularly true for iso-AB-PBI, in which the benzimidazole moieties are oriented in alternating head to head and tail to tail configurations (Scheme 4c).90 Replacing the phenyl ring in conventional PBI by the moderate base pyridine not only improves the capability to take up phosphoric acid, it also increases the solubility of para-PBI in polar aprotic organic solvents,88 which is relevant for the film forming process. Since para-PBI shows a significantly higher tensile strength for a given phosphoric acid content, pyridine containing para-PBI combines high proton conductivity and mechanical robustness.87 Polybenzimidazoles containing ether linkages (Scheme 4d) are also attracting increasing interest, not because of high phosphoric acid uptakes but for its high flexibility even at low phosphoric acid contents. 374



Perspective

In short, adducts of polybenzimidazoles and phosphoric acid are gels under most conditions. For obtaining high proton conductivity, the materials need to contain a high concentration of phosphoric acid, which naturally tends to soften the gel structure. Since this is most likely the consequence of the strong interaction between the benzimidazole unit and phosphoric acid (internal salt formation92), alternative approaches may also consider other polymeric matrices different from the commonly used polybenzimidazoles. From the above considerations, backbone flexibility and kind and density of basic sites are expected to be sensitive parameters in narrowing down interesting membrane compositions which may show high proton conductivity at lower phosphoric acid contents. Another important parameter is the nanomorphology: apart from the percolation of the phosphoric acid phase, progressive phase separation is expected to reduce the number density of interacting sites which already has been shown to increase proton conductivity.93 It is definitely true that PBI−phosphoric acid membranes are the only ones that have been used successfully in PEM fuel cells at high temperature (T ∼ 160 °C) without external humidification, but the severe voltage drop of the corresponding fuel cells at rather small currents reduces the fuel efficiency leading to the production of large amounts of waste heat which cannot always be used in satisfactory ways. This is generally thought to be the consequence of the adsorption of phosphate species onto the platinum electrocatalyst at the equilibrium potential of oxygen reduction. Another challenge for further developments, therefore, is to bind phosphoric acid in such a way that it still supports high proton conductivity without adsorbing on platinum at cathodic potentials. The fact that CsH2PO4 based fuel cells show lower over-potentials for the oxygen reduction reactions94 provides some hope that there is space for further improvements.

Compared to the dependence of phosphoric acid uptake on the type of polybenzimidazole, the preparation process seems to have an even stronger influence on how much phosphoric acid is kept within the diverse PBI structures. For a long time, PBI membranes cast from solutions of, e.g., DMAc were just immersed in aqueous solutions of phosphoric acid, but later, a process allowing for a more homogeneous mixing of PBI and phosphoric acid was developed by Benicewicz et al.91 This process relies on the fact that both PBI and their precursor monomers are soluble in pyro-phosphoric acid (PPA). The polymerization process of various kinds of PBI is carried out at high temperature under dry conditions in PPA, before casting the homogeneous solution of PBI in PPA onto a glass plate at ambient. In the course of cooling, PPA hydrolyzes, converting into ortho-phosphoric acid, which is a nonsolvent for PBI. As a consequence, PBI and phosphoric acid separate during a sol/gel transition which leads to the formation of relatively stable films. Thermodynamically speaking, it is the non-negligible water activity of “ortho-phosphoric acid” which makes it a nonsolvent for PBI, and the gelification of the solution, therefore, depends on the relative humidity which is sometimes controlled at a high level during the casting process.86 The final products are characterized by high phosphoric acid contents which are well kept within the gel structure, and adducts with a phosphoric acid content of close to 90 vol % still have a surprising mechanical strength at ambient. Since the proton conductivity very much depends on the phosphoric acid content (see below), this is considered a distinct advantage, but on the other hand, the sol/gel transition may be at least partially reversed when the membrane is heated to the operation temperature of a HT-PEM fuel cell (T ∼ 160 °C). The softening going along with the corresponding structural rearrangements is a wellknown problem of membranes produced by the PPA process. Unfortunately, ionic and covalent cross-linking strategies88 appeared to be a trade off between improving tensile strength and avoiding the appearance of brittleness. A general requirement for obtaining membranes with reasonable mechanical stability is a very high molecular weight of PBI while crystallinity is practically no issue. Because of the plasticizing effect of phosphoric acid, the weak order (crystallinity) of pure PBI is completely lost in PBI−phosphoric acid adducts. In any case, such materials are gels at not too low water activity; i.e., liquid phosphoric acid is just entrapped within a network of PBI with one phosphoric acid per benzimidazole unit strongly interacting through hydrogen bonding or pure Coulomb forces.92 Therefore, leaching of phosphoric acid in the presence of water is always an issue of this type of membrane material. Nevertheless, the apparent conductivity of PBI−phosphoric acid gels is significantly lower than the conductivity of bulk phosphoric acid. In order to achieve high proton conductivity similar to this of fully hydrated Nafion, the phosphoric acid content must be as high as ∼80 vol %.88 For this, the conductivity is already about a factor of 5 lower than for pure phosphoric acid. Apart from the small reduction of percolation, the major effect explaining this dramatic conductivity decrease could well be the decrease of hydrogen bond network “frustration” associated with the subtraction of protons through polybenzimidazole.92 As pointed out above, this is expected to reduce not only phosphoric acid’s proton conductivity in the nominally dry state but also its hygroscopicity, another factor affecting proton conductivity (see above).

4. MEMBRANES FOR REDOX-FLOW BATTERIES Redox-flow batteries are related to both PEM fuel cells and batteries (Figure 1). As other type of batteries, redox-flow batteries store electrical energy almost reversibly. Their electrochemically active masses are not stationary like in conventional batteries, they rather are solutions of redox couples pumped through the electrode compartments just as humidified fuel and air is channeled along the gas diffusion electrodes of PEM fuel cells. The cell designs (also stack designs) therefore resemble this of fuel cells, but the separator membrane requirements are quite different. The membrane has to separate the so-called anolyte and catholyte solutions (Figure 1b) which contain the electrochemically active redox couples and some supporting electrolyte (for the different types of redox couples, see, e.g., ref 95). When the flow battery is charged or discharged, an equivalent amount of ionic charge has to cross the membrane, while the ions involved in the redox process have to be efficiently separated. Apart from the obvious stability requirements (the membrane has to be durable with the oxidizing and reducing solutions) the most important issues therefore are selectivity and transport. Unfortunately, there is virtually no systematic membrane development for this very demanding application. But available membrane types have been tested in situ analyzing power output, Coulomb efficiency, and the compositional changes of anolyte and catholyte during operation. Generally speaking, PFSA membranes such as Nafion allow for the highest power 375



Perspective

densities, especially when “forced convection” of the active electrolytes is applied,96 but low Coulomb efficiency and changes of the ion concentration in the electrolyte point toward a severe crossover of electrochemically active species and water. Diverse modifications, such as the dispersion of inorganic nanoparticles (e.g., SiO2, TiO2), blending with other polymers (e.g., PVDF), or coating the membrane surface with a more selective (e.g., polyethylenimine) layer, are claimed to improve the membrane’s barrier function.97 More recently, the specific advantages associated with the use of cation and anion exchanging polyarylene membranes have clearly been recognized. Compared to PFSA membranes, hydration water is more dispersed in these membranes,49 and this is thought to be the reason for the generally higher perm-selectivity. In the case of vanadium redox flow batteries (VRFB), anion exchange membranes have the advantage of excluding electrochemically active VO2+ from entering the membrane, which also reduces the reaction rate with this highly oxidizing species.97 For this application, even moderately basic anion exchange membranes can be used, because their functional groups (commonly nonquarternized simple amines) are fully protonated under the acidic conditions of VRFBs. Unfortunately, polyarylenes not only suffer from limited oxidation stability, but their insufficient ionic conductivity under the conditions of redox flow batteries is another severe downside. Since the issue of “barrier function for the electrochemically active species versus high ionic conductivity” is still unresolved, the only system developed beyond the laboratory scale is the VRFB using vanadium redox couples for both electrolytes (VO2+/VO2+ and V3+/V2+). In this case, any crossover reduces the Coulomb efficiency and increases the internal heat production but does not lead to irreversible chemical contaminations. A framework which may allow for a more systematic development of membranes for particular redox-flow systems is briefly presented in the following. An obvious approach, followed by several groups, relies on Donnan exclusion98 referring to the effect that ions of the same sign as the fixed ions (e.g., −SO3− in the case of PFSA membranes) are prevented from entering the membrane. Therefore, anion exchange membranes may be preferred for redox-flow batteries making use of cations as redox-couples (e.g., VRFB), but in many cases cation exchange membranes are chosen because of power and stability considerations. Fundamentally, the point at which Donnan exclusion (perm-selectivity) starts to decay is expected to depend on hydration level, kind of fixed and mobile ions, nanomorphology, and chemical nature of the backbone. At low hydration levels, at which all water molecules are involved in exothermal ion solvation (see Section 3.1.1), counterions may exchange for other ions of the same sign, but this is not possible for the fixed ions of opposite sign (Figure 14a). Even in the regime where entropy is the main driving force for membrane hydration (see Section 3.1.1), the uptake of ions of the same sign as the fixed ions is unfavorable albeit not completely excluded (Figure 14b). How Donnan exclusion decays with the degree of hydration is expected to depend not only on the solvation energies of all ions involved but also on the nanomorphology and solvent interaction with the backbone. For the same hydration number λ (here, water molecules per ion), the dielectric constant of a sulfonated polyether ketone was found to be significantly lower than that of a PFSA membrane99 clearly indicating that water dispersion and backbone interaction enters into the solvation capability of

Figure 14. Schematic illustration of ion distribution in an anion exchange membrane: (a) in the Donnan exclusion regime, where only anions can exchange, and (b) beyond Donnan exclusion, where also co-ions can enter the membrane.

hydration water in ion exchange membranes. Very useful information is therefore quantitative data on the decay of Donnan exclusion with increasing hydration number λ which may help to identify the maximum hydration for the targeted ion selectivity. Of course, the latter depends on the ion mobility, which, in aqueous environment, essentially depends on ion size and charge. This is probably the reason why cation exchange membranes still show reasonable barrier function in VRFB in which the electrochemically active ions are all cations. But before addressing the issue of selectivity between different cations or anions, it should clearly be noted that the hydration level of the membrane during operation must be controlled. This can be done by controlling the external conditions, in particular the concentration of the supporting electrolyte, which allows adjusting the osmotic pressure difference. Provided there is no internal membrane pressure (see Section 3.1.1), the ion concentration within the membrane will roughly match that of the electrolytes. The other strategy is reinforcing the membrane so as to limit the hydration level. It may also be feasible to benefit from the selective uptake and mobility for ions of the same sign. Apart from simple electrostatic considerations, here, also specific interaction such as hydrogen bonding or even some covalency may enter solvation and ion pairing energies. The experimental determination of Donnan equilibria and electronic structure calculations of solvation and ion pairing energies are therefore helpful. The latter is of particular interest for understanding counterion condensation (condensation of fixed and “mobile” ion) excluding ions from being transported through the membrane. Apart from ionic transport, the transport of solvent (in most cases water) is another critical issue. This may be the 376



Perspective

propylene and ethylene carbonate, dissolve high concentrations of Li-salts, but these salts do not fully dissociate into single solvated ions. The presence of contact ion pairs and even triple ions100,101 and the appearance of correlated ionic diffusion102,103 are characteristic for such systems. The fact that it is the oxygen of the small solvent molecules coordinating to the Li ions led several groups to dissolve salts into polymeric ethers, in particular poly(ethylene oxide) (PEO), and study their ionic conductivity more than 40 years ago.104,105 The corresponding solvation chemistry is characterized by strong Li+ coordination through the ether oxygen while the anions are only weakly interacting with PEO.106 As a result, such electrolytes are mainly anion conductors with Li+ transference numbers even smaller than for liquid electrolytes.107 It must, however, be noted that in the case of liquid electrolytes there is another pathway for Li-ion transfer: the counter diffusion of contact ion pairs and anions.108 Low transference numbers and absolute conductivities more than 1 order of magnitude lower than these of liquid systems (Figure 16) has led to the development of gel-

consequence of osmosis (driven by the chemical water potential difference of anolyte and catholyte) and electroosmotic water drag (the coupled transport of ions and water; see Figure 8). As shown above (see Section 3.1.1), both transport modes dramatically increase beyond a certain level of hydration, and it is not clear yet whether this level coincides with the decay of Donnan exclusion. By all means, it is definitely useful to keep swelling below this limit. High selectivity must be combined with high ionic transport of some ion different from these involved in the electrochemical charging and discharging processes. Since this definitely increases with increasing hydration, like in the case of ion conducting membranes for fuel cells, the challenge is to obtain high ionic transport at a hydration level still acceptable for the required selectivity. This situation very much resembles this of high temperature, low humidity operation of fuel cell membranes, and an interesting approach therefore is to maximize the local concentration of appropriately chosen ionic moieties. For the strongly acidic cation exchanging group −SO 3 − high proton conductivities have already been demonstrated (see Section 3.1.2), and basic anion exchanging amine groups may allow for high OH− conductivity at high levels of hydration. Exchanging OH− versus anions with lower hydration enthalpy (e.g., Cl−), as commonly present in electrolytes of redox-flow batteries, leads to a severe drop in conductivity even at high water contents. In the case of the Cl− form of an anion exchange membrane, the conductivity diffusion coefficient systematically remains a factor of 2 below the water diffusion coefficient (Figure 15). This points toward

Figure 16. Comparison of the total conductivities of typical Li+ conductors: mixed conducting solution of a Li-salt in aprotic polar solvent,118 solid polymer electrolyte (salt in polymer),109 single Li-ion conducting polyelectrolyte solvated with an aprotic polar solvent,111 and a fully polymeric single Li-ion conductor.115

type electrolytes, in which liquid electrolytes are simply entrapped within the pores of a polymeric matrix (e.g., PVDF). Of course, these are not true polymer membranes (like is the case for PBI−phosphoric acid gels; see Section 3.3), but they can mechanically better separate the active masses while providing conductivities close to these of liquid electrolytes. A comprehensive comparison of gel-type and solid polymer electrolytes is given in ref 109. Both types of electrolyte conduct anions and cations, but in the case of liquid solutions, the ions have a stable solvation shell, while for salt in polymer solutions, e.g., Li+, it must be transported without the polymeric solvent. The strong chemical interaction between polymer and Li+ then requires a dynamical coupling between segmental motion and Li+ diffusion as directly evidenced by NMR.110 What is still missing for battery technology is a highly conducting fully polymeric single Li-ion conductor containing no flammable low molecular weight solvent. Recently, single ion conductors with very high conductivity (>10−3 S cm−1) have been obtained by ion exchanging the highly proton conducting polyelectrolyte S-220 (see Figure 10) with Li+

−

Figure 15. Water and Cl conductivity diffusion coefficients for a noncross-linked poly(arylene sulfone) functionalized with TMA groups (FAA-3 supplied by Fumatech).66 The suggested formation of ionic cross-links is illustrated (see text).

some stable associates (cross-links) which only break up at very high water content (inset Figure 15),66 which is another example for the relevance of residual ionic interaction.

5. ELECTROLYTES FOR ALKALINE ION BATTERIES Although safety issues ask for fully polymeric separator materials in alkaline ion batteries, fully polymeric alkaline ion conducting membranes hardly found their way into high drain battery applications yet. Since it is beyond the scope of this article to address the many specific aspects associated with this particular application, only a few general issues of the formation and mobility of ionic charge carriers are briefly discussed. Here, residual ionic interaction is even more of an issue. Liquid polar aprotic solvents, such as commonly used 377



Perspective

the investigation of structure and dynamics on a range of length and time scales.

before solvating with dimethyl-sulfoxide (DMSO) (Figure 16).111 Interestingly, Li+ diffusion is about five times slower than the diffusion of the solvent, which points to either strong counterion association (incomplete dissociation) or a retardation of Li+ mobility as a consequence of its heavy solvation by DMSO.112 Since the anions are completely immobilized, no concentration polarization effects 113 (formation of salt concentration gradients) can occur in such electrolytes, and the choice of DMSO as a solvent makes them interesting candidates for Li/oxygen batteries.114 This is actually surprising, because the high donor number (DN ∼ 30) renders DMSO’s oxidation limit to be very low. The downside of DMSO’s flammability is resolved for a very recent triblock copolymer consisting of two PEO blocks and a segment bearing Lisulfamide as ionic group.115 This is a true polymeric Li+ single ion conductor, but conductivity values around 10−5 S cm−1 (Figure 16) are still too low for battery applications. An even more efficient decoupling of the Li+ dynamics from this of the polymeric solvent and the immobilized anion, therefore, appears to be an interesting challenge for future research.

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] Notes

The authors declare no competing financial interest. Biography After receiving his Diploma in Mineralogy at the University of Cologne, Klaus-Dieter Kreuer did a PhD in the department of Chemistry at the University of Stuttgart. As a fellow of the “Studienstiftung des Deutschen Volkes” he benefited from a research stay at the California Institute of Technology (group of R. Vaughan) and a Max-Planck award allowed him to join the Massachusetts Institute of Technology as a visiting scientist. Later, Klaus-Dieter Kreuer built an R&D group within a Swiss-German company (Endress & Hauser) before joining the Max-Planck-Institute for Solid State Research, where he assisted J. Maier in building his new department. Since 1990 Klaus-Dieter Kreuer is lecturing at the University of Stuttgart from which he received his Habilitation degree.

■

6. FINAL REMARKS According to the large disparities of their levels of development, the perspectives for future research appear to be very diverse for the different types of ion conducting membranes and electrochemical applications, albeit a few general aspects are also clearly identified. Different strategies toward high temperature, low humidity operation in fuel cells, the combination of high conductivity and selectivity in redox flow batteries, and polymeric single ion conductors for battery applications are adumbrated above. In all cases, high mobility of ionic charge carriers is an issue which, for most liquid electrolytes, is closely related to the solvent diffusion coefficient and the viscosity through the Nernst− Einstein and Stokes−Einstein relationships. One of the challenges for obtaining high ionic mobility in polymeric environment is the effective decoupling of ionic diffusion from the dynamics of the polymeric structure. While this is possible for protonic charge carriers in some hydrogen bonded liquids such as phosphoric acid83 (see Section 3.3) and specially designed polymers containing heterocycles4 or phosphonic acid5 as protogenic groups, similar mechanisms have not yet been described for other ions. The separation of ionic charge carriers, i.e., the dissociation process, is closely related to ion solvation, which needs to be better understood beyond simple electrostatic consideration (Debye−Hückel approach). A lot can be learned from the electrochemistry of solutions and polyelectrolytes,116 for which simple empirical concepts such as “hard, softacid, base” or Manning counterion condensation are aiming at a qualitative description of ionic interactions in particular systems. But also spectroscopic tools, especially NMR,117 and ab initio electronic structure calculations may surely contribute to a better understanding of specific chemical interactions in processes such as charge carrier formation and mobility. Since in many cases the ionic groups are part of the polymeric structure, such interactions affect the polymer conformations as well. Because of the close relation between ionic transport and nanomorphology, insights into how these polymer conformations constrain the formation of ordered nanomorphologies must be another key area of future research. This is where electrochemistry and polymer chemistry meet with experimental, simulation, and theoretical tools applied to

ACKNOWLEDGMENTS The many critical discussions with J. Melchior, A. Wohlfarth, and M. Marino, technical assistance through A. Fuchs, and proof reading by A. Kuhn and J. Jackson (all MPI-FKF) are gratefully acknowledged.

■ ■

DEDICATION Dedicated to my colleague and friend Per Jacobsson. REFERENCES

(1) Polymers for Energy Storage and Delivery: Polyelectrolytes for Batteries and Fuel Cells; ACS Symposium Series 1096; Page, K. A., Soles, C. L., Runt, J., Eds.; American Chemical Society: Washington, DC, 2012. (2) Hess, M.; Jones, R. G.; Kahovec, J.; Kitayama, T.; Kratochvil, P.; Kubisa, P.; Mormann, W.; Stepto, R. F. T.; Tabak, D.; Vohlidal, J.; et al. Pure Appl. Chem. 2006, 78 (11), 2067−2074. (3) Saarinen, V.; Kreuer, K. D.; Schuster, M.; Merkle, R.; Maier, J. Solid State Ionics 2007, 178, 533. (4) Herz, H. G.; Kreuer, K. D.; Maier, J.; Scharfenberger, G.; Schuster, M. F. H.; Meyer, W. H. Electrochim. Acta 2003, 48, 2165. (5) Steininger, H.; Schuster, M.; Kreuer, K. D.; Maier, J. Solid State Ionics 2006, 177, 2457. (6) Rager, T.; Schuster, M.; Steininger, H.; Kreuer, K. D. Adv. Mater. 2007, 19, 3317. (7) Steininger, H.; Schuster, M.; Kreuer, K. D.; Kaltbeitzel, A.; Bingöl, B.; Meyer, W. H.; Schauff, S.; Brunklaus, G.; Maier, J.; Spiess, H. W. Phys. Chem. Chem. Phys. 2007, 9, 1764. (8) Schuster, M.; Rager, T.; Noda, A.; Kreuer, K. D.; Maier, J. Fuel Cells 2005, 5, 355. (9) Atanasov, V.; Kerres, J. Macromolecules 2011, 44, 6416. (10) Dippel, T.; Kreuer, K. D.; Lassègues, J. C.; Rodriguez, D. Solid State Ionics 1993, 61, 41. (11) Kreuer, K. D. Chem. Mater. 1996, 8, 610. (12) Wainright, J. S.; Wang, J. T.; Weng, D.; Savinell, R. F.; Litt, M. J. Electrochem. Soc. 1995, 142, L121. (13) Peron, J.; Nedellec, Y.; Jones, D. J.; Rozière, J. J. Power Sources 2008, 185, 1209. (14) Gierke, T. D.; Munn, G. E.; Wilson, F. C. J. Polym. Sci., Part B: Polym. Phys. 1981, 19, 1687. (15) Rubatat, L.; Rollet, A. L.; Gebel, G.; Diat, O. Macromolecules 2002, 35, 4050. 378



Perspective

(16) Li, J.; Park, J. K.; Moore, R. B.; Madsen, L. A. Nat. Mater. 2011, 10, 507. (17) Kreuer, K. D.; Portale, G. Adv. Funct. Mat. 2013, DOI: 10.1002/ adfm.201300376. (18) Kusoglu, A.; Modestino, M. A.; Hexemer, A.; Segalman, R. A.; Weber, A. Z. ACS Macro Lett. 2012, 1, 33. (19) Lyonnard, S.; Berrod, Q.; Gebel, G.; Guillermo, A. Presented at E-MRS Spring Meeting, Symposium F, Strasbourg, 2012. (20) Kreuer, K. D. et al. In preparation. (21) Gebel, G.; Diat, O. Fuel Cells 2005, 5, 261. (22) Rubatat, L.; Gebel, G.; Diat, O. Macromolecules 2004, 37, 7772. (23) Tant, M. R.; Darst, K. P.; Lee,K. D.; Martin, C. W. In Multiphase Polymers: Blends and Ionomers; ACS Symposium Series 395; Utracki, L. A., Weiss, R. A., Eds.; American Chemical Society: Washington, DC, 1989; pp 370−400. (24) Ghielmi, A.; Vaccarono, P.; Troglia, C.; Arcella, V. J. Power Sources 2005, 145, 108. (25) Kreuer, K. D.; Paddison, S. J.; Spohr, E.; Schuster, M. Chem. Rev. 2004, 104, 4637. (26) Escoubes, M.; Pineri, M. In Perfluorinated Ionomer Membranes; Eisenberg, A., Yeager, H. L., , Eds.; ACS Symposium Series 180; American Chemical Society: Washington, DC, 1982; pp 9−24. (27) Kreuer, K. D. Solid State Ionics 2013, 252, 93−101. (28) Choi, P. H.; Jalani, N. H.; Datta, R. J. Electrochem. Soc. 2005, 152, E84. (29) Choi, P. H.; Datta, R. J. Electrochem. Soc. 2003, 150, E601. (30) Freger, V. J. Phys. Chem. B 2009, 113, 24. (31) Bass, M.; Berman, A.; Singh, A.; Konovalov, O.; Freger, V. Macromolecules 2011, 44, 2893. (32) Eastman, S. A.; Kim, S.; Page, K. A.; Rowe, B. W.; Kang, S.; Soles, C. L.; Yager, K. G. Macromolecules 2012, 45, 7920. (33) Telfah, A.; G. Majer, G.; Kreuer, K. D.; Schuster, M.; Maier, J. Solid State Ionics 2010, 181, 461. (34) Dippel, T.; Kreuer, K. D. Solid State Ionics 1991, 46, 3. (35) Kreuer, K. D.; Weppner, W.; Rabenau, A. Solid State Ionics 1981, 3/4, 353. (36) Kreuer, K. D.; Rabenau, A.; Weppner, W. Angew. Chem., Int. Ed. Engl. 1982, 21, 208. (37) Tuckerman, M. E.; Laasonen, K.; Sprik, M.; Parrinello, M. J. Chem. Phys. 1995, 103, 150. (38) Agmon, N. Chem. Phys. Lett. 1995, 244, 456. (39) Kreuer, K. D. Solid State Ionics 2000, 136−137, 149. (40) Kreuer, K. D.; Schuster, M.; Obliers, B.; Diat, O.; Traub, U.; Fuchs, A.; Klock, U.; Paddison, S. J.; Maier, J. J. Power Sources 2008, 178, 499. (41) Kreuer, K. D. Hydrocarbon membranes. In Handbook of Fuel Cells: Fundamentals, Technology and Applications. Vielstich, W., Lamm, A., Gasteiger, H., Eds.; John Wiley & Sons Ltd.: Chichester, U.K., 2003; Vol. 3, Part 3, pp 420−435. (42) Almeras, Y.; Barrat, J. L.; Bocquet, L. J. Phys. IV 2000, 10, 27. (43) Merlo, L.; Ghielmi, A.; Cirillo, L.; Cirillo, L.; Gebert, M.; Arcella, V. J. Power Sources 2007, 171, 140. (44) Ghassemzadeh, L.; Marrony, M.; Barrera, R.; Kreuer, K. D.; Maier, J.; Müller, K. J. Power Sources 2009, 186, 334. (45) Curtin, D. E.; Lousenberg, R. D.; Henry, T. J.; Tangeman, P. C.; Tisack, M. E. J. Power Sources 2004, 131, 41. (46) LaConti, A.; Hamdan, M.; McDonald, R. In Handbook of Fuel Cells: Fundamentals, Technology, and Applications; Vielstich, W., Lamm, A., Gasteiger, H., Eds.; John Wiley and Sons Ltd.: Chichester, U.K., 2003; pp 647−662. (47) Ghassemzadeh, L.; Kreuer, K. D.; Maier, J.; Müller, K. J. Phys. Chem. C 2010, 34, 14635. (48) Collete, F.; Thominette, F. M.; Mendil-Jakani, H.; Gebel, G. J. Membr. Sci. 2013, 435, 242. (49) Kreuer, K. D. J. Membr. Sci. 2001, 185, 29. (50) Rikukawa, M.; Sanui, K. Prog. Polym. Sci. 2000, 25, 1463. (51) Hickner, M.; Ghassemi, H.; Kim, Y. S.; Einsla, B. R.; McGrath, J. E. Chem. Rev. 2004, 104, 4587−4612.

(52) Schuster, M.; Kreuer, K. D.; Andersen, H. T.; Maier, J. Macromolecules 2007, 40, 598. (53) de Araujo, C. C.; Kreuer, K. D.; Schuster, M.; Portale, G.; Mendil-Jakani, H.; Gebel, G.; Maier, J. Phys. Chem. Chem. Phys. 2009, 11, 3305. (54) Schuster, M.; de Araujo, C. C.; Atanasov, V.; Andersen, H. T.; Kreuer, K. D.; Maier, J. Macromolecules 2009, 42, 3129. (55) Si, K.; Dong, Dong, D.; Wycisk, R.; Litt, M. J. Mater. Chem. 2012, 22, 20907. (56) Shi, Z. Q.; Holdcroft, S. Macromolecules 2005, 38, 4193. (57) Ghassemi, H.; McGrath, J. E.; Zawodzinski, T. A. Polymer 2006, 47, 4132. (58) Bae, B.; Yoda, T..; Miyatake, K.; Uchida, H.; Watanabe, M. Angew. Chem., Int. Ed. 2010, 49, 317. (59) Titvinidze, G.; Kreuer, K. D.; Schuster, M.; de Araujo, C. C.; Melchior, J.; Meyer, W. H. Adv. Funct. Mater. 2012, 22, 4456. (60) Takamuku, S.; Jannasch, P. Macromol. Rapid Commun. 2011, 32, 474. (61) Atanasov, V.; Buerger, M.; Wohlfarth, A.; Maier, J. Polym. Bull. 2012, 68, 317. (62) Hamrock, S. J.; Herring, A. M. Proton Exchange Membrane Fuel Cells. In Fuel Cells; Kreuer, K. D., Ed.; Springer: 2013; Chapter 17, pp 577−606. (63) Takamuku, S.; Jannasch, P. Patent Application SE1300225-8 (64) Takamuku, S.; Jannasch, P. Polym. Chem. 2012, 3, 1202. (65) Robertson, N. J.; Kostalik, H. A.; Clark, T. J.; Mutolo, P. F.; Abruna, H. D.; Coates, G. W. J. Am. Chem. Soc. 2010, 132, 3400. (66) Marino, M.; Melchior, J. P.; Wohlfarth, A.; Kreuer, K. D. J. Membr. Sci. 2013, submitted. (67) Mamlouk, M.; Horsfall, J.; Williams, C.; Scott, K. Int. J. Hydrogen Energy 2012, 37, 11912. (68) Bauer, B.; Strathmann, H.; Effenberger, F. Desalination 1990, 79, 125. (69) Noonan, K. J. T.; Hugar, K. M.; Kostalik, H. A.; Lobkovsky, E. B.; Abruna, H. D.; Coates, G. W. J. Am. Chem. Soc. 2012, 134, 18161. (70) Henkensmeier, D.; Cho, H.-R.; Kim, H.-J.; Kirchner, C. N.; Leppin, J.; Dyck, A.; Jang, J. H.; Cho, E.; Nam, S.-W.; Lim, T.-H. Polym. Degrad. Stab. 2012, 97, 264. (71) Tomoi, M.; Yamaguchi, K.; Ando, R.; Kantake, Y.; Aosaki, Y.; Kubota, H. J. Appl. Polym. Sci. 1997, 64, 1161. (72) Qiu, B.; Lin, B.; Qiu, L.; Yan, F. J. Mater. Chem. 2012, 22, 1040. (73) Hibbs, M. R. J. Polym. Sci., Part B: Polym. Phys. 2012, DOI: 10.1002/polb.23149. (74) Chen, D.; Hickner, M. A. ACS Appl. Mater. Interfaces 2012, 4, 5775. (75) Deavin, O. I.; Murphy, S.; Ong, A. L.; Poynton, S. D.; Zeng, R.; Herman, H.; Varcoe, J. R. Energy Environ. Sci. 2012, 5, 8584. (76) Schwesinger, R.; Link, R.; Wenzl, P.; Kossek, S.; Keller, M. Chem.Eur. J. 2006, 12, 429. (77) Inaba, M.; Matsui, Y.; Saito, M.; Tasaka, A.; Fukuta, K.; Watanabe, S.; Yanagi, H. Electrochemistry 2011, 79, 322−325. (78) Couture, G.; Alaaeddine, A.; Boschet, F.; Ameduri, B. Prog. Polym. Sci. 2011, 36, 1521. (79) Merle, G.; Wessling, M.; Nijmeijer, K. J. Membr. Sci. 2011, 377, 1. (80) Slade, R. C. T.; Kizewski, J. P.; Poynton, S. D.; Zeng, R.; Varcoe, J. R. Alkaline Membrane Fuel Cells; In Fuel Cells; Kreuer, K. D., Ed.; Springer: 2013; Chapter 2, pp 9−29. (81) Munson, R. A. J. Phys. Chem. 1964, 68, 3374. (82) Dippel, T.; Kreuer, K. D.; Lassègues, J. C.; Rodriguez, D. Solid State Ionics 1993, 61, 41. (83) Vilciauskas, L.; Tuckerman, M. E.; Bester, G.; Paddison, S. J.; Kreuer, K. D. Nature Chem. 2012, 4, 461. (84) Melchior, J. P.; et al. In preparation. (85) Vilciauskas, L. Ph.D. thesis. University of Stuttgart, Chemistry Department, 2012. (86) Molleo, M.; Schmidt, T. J.; Benicewicz, B. C. Polybenzimidazole Fuel Cell Technology. In Fuel Cells; Kreuer, K. D., Ed.; Springer: 2013; Chapter 13, pp 391−432. 379



Perspective

(87) Asensio, J. A.; Sanchez, E. M.; Gomez-Romero, P. Chem. Soc. Rev. 2010, 39, 3210. (88) Li, Q.; Jensen, J. O.; Savinell, R. F.; Bjerrum, N. J. Prog. Polym. Sci. 2009, 34, 449. (89) Mader, J.; Xiao, L.; Schmidt, T. J.; Benicewicz, B. C. Adv. Polym. Sci. 2008, 216, 63. (90) Gulledge, A. L.; Gu, B.; Benicewicz, B. C. J. Polym. Sci., Part A: Polym. Chem. 2012, 50, 306. (91) Xiao, L.; Zhang, H.; Scanlon, E.; Ramanathan, L. S.; Choe, E. W.; Rogers, D.; Apple, T.; Benicewicz, B. C. Chem. Mater. 2005, 17, 5328. (92) Vilciauskas, L.; Tuckerman, M. E.; Melchior, J. P.; Bester, G.; Kreuer, K. D. Solid State Ionics 2013, DOI: 10.1016/j.ssi.2013.07.003. (93) Weber, J.; Kreuer, K. D.; Maier, J.; Thomas, A. Adv. Mater. 2008, 20, 2595. (94) Boysen, D. A.; Uda, T.; Chisholm, C. R. I.; Haile, S. M. Science 2004, 303, 68. (95) Ponce de León, C.; Frías-Ferrer, A.; González-García, J.; Szánto, D. A.; Walsh, F. C. J. Power Sources 2006, 160, 716. (96) Zaffou, Z.; Li,W. N.; Perry, M. L. In Polymers for Energy Storage and Delivery: Polyelectrolytes for Batteries and Fuel Cells; ACS Symposium Series; Page, K. A., Soles, C. L., Runt, J., Eds.; American Chemical Society: Washington, DC, 2012; Vol. 1096, Chapter 7, pp 107−127. (97) Ding, C.; Zhang, H.; Li, X.; Liu, T.; Xing, F. J. Phys. Chem. Lett. 2013, 4, 1281. (98) Donnan, F. G. Z. Elektrochem. Angew. Physikal. Chem. 1911, 17, 572. (99) Paddison, S. J.; Bender, G.; Kreuer, K. D.; Nicoloso, N.; Zawodzinski, T. J. New Mater. Electrochem. Syst. 2000, 3, 291. (100) Teeters, D.; Frech, R. Solid State Ionics 1986, 271, 18. (101) Albinsson, I.; Jacobsson, P.; Mellander, B. E.; Stevens, J. R. Solid State Ionics 1992, 53−56, 1044. (102) Arumugam, S.; Shi, J.; Tunstall, D. P.; Vincent, C. A. J. Phys.: Condens. Matter 1993, 5, 153. (103) Boden, N.; Leng, S. A.; Ward, I. M. Solid State Ionics 1991, 45, 261. (104) Fenton, D. E.; Parker, J. M.; Wright, P. V. Polymer 1973, 14, 589. (105) Armand, M.; Chabagno, J. M.; Duclot, M. In Fast Ion Transport in Solids; Vashishta, P., Ed.; North Holland: New York, 1979; p 131. (106) Armand, M. Adv. Mater. 1990, 2, 278. (107) Evans, J.; Vincent, C. A.; Bruce, P. G. Polymer 1987, 28, 2324. (108) Stolwijk, N. A.; Obeidi, S. Phys. Rev. Lett. 2004, 93, 125901. (109) Meyer, W. H. Adv. Mater. 1998, 10, 439. (110) Kunze, M.; Karatas, Y.; Wiemhöfer, H. D.; Schönhoff, M. Macromolecules 2012, 45, 8328. (111) Kreuer, K. D.; Wohlfarth, A.; de Araujo, C. C.; Fuchs, A.; Maier, J. ChemPhysChem 2011, 12, 2558. (112) Smiatek, J.; Wohlfarth, A.; Holm, C. New J. Phys. (NJP) 2013, submitted. (113) Maier, J. Adv. Funct. Mater. 2011, 21, 1448. (114) Chen, Y.; Freunberger, S. A.; Peng, Z.; Barde, F.; Bruce, P. G. J. Am. Chem. Soc. 2012, 134, 7952. (115) Bouchet, R.; Maria, S.; Meziane, R.; Aboulaich, A.; Lienafa, L.; Bonnet, J. P.; Armand, M. Nat. Mater. 2013, 12, 452. (116) Marcus, Y.; Hefter, G. Chem. Rev. 2006, 106, 4585. (117) Bogle, X. S.; Vazquez, R.; Greenbaum, S.; Cresce, A. V. W.; Xu, K. J. Phys. Chem. Lett. 2013, 4, 1664. (118) Gores, H. J.; Barthel, J.; Zugmann, S.; Moosbauer, D.; Amereller, M.; Hartl, R.; Maurer, A. Liquid Nonaqueous Electrolytes. In Handbook of Battery Materials, 2nd ed.; Daniel, C., Besenhard, J. O., Eds.; Wiley-VCH Verlag GmbH & Co. KGaA: Weinheim, Germany, 2011.

380


Ion Conducting Membranes for Fuel Cells and other Electrochemical

Recommend Documents