Ion Conducting Membranes for Fuel Cells and other Electrochemical Devices

Ion Conducting Membranes for Fuel Cells and other ElectrochemicalDevicesKlaus-Dieter Kreuer*

Max-Planck-Institut fur Festkorperforschung, Heisenbergstrasse 1, D-70569 Stuttgart, Germany

ABSTRACT: Transport and stability issues of proton and hydroxide ionconducting separator membranes for fuel cells are critically discussed from afundamental point of view. Considerations of structure and dynamics on themolecular scale to the device level equally imply polymer-chemical andelectrochemical aspects which are closely related for this class of materials.The importance of ion/solvent, residual ion/ion, and solvent/polymerinteractions for the formation and mobility of ionic charge carriers andselective ionic transport and even as driving forces for nanoscale ordering isemphasized, and it is shown that, apart from simple electrostatics, specificchemical interactions must be considered. On the basis of this under-standing, suggestions are being made for the modification of existing and thedevelopment of new membrane types, not only for fuel cells but also forother electrochemical energy conversion and storage devices such as redox-flow 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 asfuel cells, redox-f low, and alkaline ion batteries, rely on theamazing properties of ion conducting polymer membranes.These devices can function only if the used membranesefficiently separate the electrochemically active masses (electro-des) and mediate the electrochemical reactions taking place atthe anode and cathode by conducting specific ions. These ionsmay be protons or hydroxide ions in the case of PEM-fuel cells,Li+ in lithium ion batteries, or non-electrochemically activeanions or cations in the case of redox-f low batteries. Apart fromthese key properties, many other requirements, especiallyconcerning stability, render the development of suchmembranes a formidable task!This article briefly summarizes our current understanding of

membrane structure, reactivity, and dynamics on differentlength scales with special emphasis on some surprising recentinsights; these have not only extended the understanding ofwell-established membrane materials such as Nafion andpolybenzimidazole-phosphoric acid adducts but also consol-idate the basis for the modification of existing and thedevelopment of new membrane types for the above-mentionedapplications. When done efficiently, this is a multidisciplinaryprocess comprising ab initio calculations, complex organic/inorganic synthesis, and comprehensive physicochemicalcharacterization. The great beauty of such a combined approachis exemplified by retracing the development of protonconducting multiblock copolymers for PEM fuel cell applica-tions and Li+ conducting polyelectrolytes for application in Li-ion batteries.

The present article develops a few perspectives from aspecific point of view, in that both the analysis of the state ofthe art and the presented research perspectives are biased bythe author’s own work. Because of the limited space and theincreasing abundance of the relevant literature, references arerestricted 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 recentpublication of the American Chemical Society1 can serve as asuitable starting point for accessing the comprehensiveliterature.

2. TYPES OF ION CONDUCTING POLYMERS ANDTHEIR APPLICATIONS IN ELECTROCHEMICALENERGY CONVERSION AND STORAGE

It goes without saying that common polymers such aspolyethylene (PE) or Teflon (PTFE) do not conduct ionsper se. Generally speaking, ionic conductivity occurs only in thepresence of moieties dissociating into ionic species which, tosome extent, diffuse in an uncorrelated way allowing theirreversible separation of ionic charges, which is a requirementof any ion conduction process.

Special Issue: Celebrating Twenty-Five Years of Chemistry ofMaterials

Received: August 14, 2013Revised: October 23, 2013Published: November 19, 2013

Perspective

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In the case of the best known ion conducting polymer,Nafion, sulfonic acid groups (−SO3H) are part of themacromolecular structure; i.e., they are covalently immobilized.Since these groups are super acids, they strongly interact withwater leading to hydration (solvation) and virtually completedissociation, even at low water activity (relative humidity). Thehydration water not only stabilizes the separation of cations(here protons) and anions (here sulfonic anions), it alsoenables the mobility of the hydrated protons. While initiallydeveloped as a Na+ conducting separator membrane for chlor-alkali electrolysis, Nafion in its acid form (the sulfonic acidfunctional group is sometimes termed protogenic group) laterbecame the benchmark membrane for PEM-fuel cellapplications as well (Figure 1a). According to the IUPACnomenclature,2 Nafion belongs to the family of cation-exchangepolymers because, in contact with aqueous solutions, especiallymonovalent cations can easily exchange, albeit with somepreference, for certain ions. With respect to the molecularstructure, Nafion is an ionomer, a polymer with a small butsignificant proportion of the constitutional units having ionic orionizable groups.2 In fact, Nafion is the proto-type of a class ofperfluorosulfonic acids (PFSA) which comprises a variety ofdifferent side chain architectures. In particular, short side chain(SSC) PFSAs seem to have the potential to outperform Nafionwhen it comes to increasing the operation temperature of PEM-fuel cells. Although the modification of PFSA membranes, e.g.,by dispersing nanoparticles in the membrane, becameextremely popular, there is no accepted understanding of thereported effects (e.g., conductivity enhancement). What seemsto be clear at this point is that such composites show higherconductivity at low levels of hydration only if the particles addto the total ion exchange capacity (IEC) of the membrane.Membranes hosting particles with a high degree of sulfonic acidfunctionalization seem to have systematically higher conductiv-ities.There is also a wealth of sulfonic acid functionalized

hydrocarbon membranes with conductivities similar to theseof PFSAs at high degrees of hydration, but low conductivities atlow hydration levels and unsatisfactory mechanical propertiesare unfortunately characteristic features making them inferior tothe more costly PFSA membranes for many applications. Thisapplies to hydrocarbons with a random distribution of sulfonicacid groups which are still ionomers, as opposed to di-, tri-, andmultiblock copolymers with a similar overall concentration ofionic groups. In the latter case, the protogenic groups arelocated on one kind of block of the molecular structures which

leads to a concentration of ionic groups within single domainsof their phase- separated morphology. Because of the high localconcentration of sulfonic acid groups, these domains arepolyelectrolytes,2 rather than ionomers, and show high protonconductivity even at low levels of hydration. This, together withacceptable mechanical properties emerging from the unsulfo-nated domain, renders these and related classes of ionconducting polymers an interesting alternative to PFSAmembranes.Apart from the major focus on sulfonic acid functionalized

polymers, there is an increasing interest in anion-exchangepolymers because of the promise of using them in fuel cells withnon-noble metal catalysts within their electrode structures.Most anion exchange membranes are hydrocarbons, andvirtually all contain some sort of quaternary amine as basicfunctional groups. They fall into the category of ionomers albeitwith relatively high concentrations of anion exchanging groups,making cross-linking necessary in order to prevent themembranes from dissolving in water. Recent progress in thefield notwithstanding, stability issues and the fact that anymembrane of this type readily converts into its bicarbonate(HCO3

−) form make the application of anion exchangepolymer 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), PFSAmembranes are still used because of their robustness and theavailability of compatible electrode structures, although hydro-carbon membranes generally exhibit lower water and methanolpermeation. Hydrocarbons with very polar backbones such ashighly sulfonated polyphenylene sulfones even show very highselectivity for water compared to methanol absorption.3

Traditionally, PFSA membranes are also being used asseparators in redox-flow batteries (Figure 1b), but morerecently hydrocarbon based anion exchange membranes areconsidered as well. Especially when the redox active species arecations (like in vanadium flow batteries), the high perm-selectivity of many hydrocarbon based anion exchangemembranes is a particular advantage, and it is interesting tonote that many membranes which are not stable in theirhydroxide (OH−) form show acceptable stability in their, e.g.,Cl−, Br−, or SO4

2− form under acidic conditions, which aretypical for many types of redox-flow batteries.All membrane types mentioned so far require some

hydration to conduct ions, and the diffusions of hydrationwater and ions are generally related. Approaches aiming at “dry”proton conducting polymers containing no low-molecular-

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.

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weight solvent (e.g., water) essentially make use of moietiessuch as heterocycles and phosphonic acids which combineproton donor and acceptor functions with the possibility ofcovalent immobilization as terminal functions of fully polymericstructures.4−7 In this way, it was possible to completelydecouple the long-range diffusion of protons from this of thepolymer,4 but because of the inherent conductivity limits(usually more than 1 order of magnitude lower than this ofhydrated systems, Figure 2) and interferences with the oxygen

reduction reaction at platinum surfaces,8 such ion conductingpolymers have not yet found their way into electrochemicalapplications. Although this class of materials is not discussed inthe present article, one should keep an eye on fully polymericelectrolytes for several reasons. The fact that these are not oronly moderately acidic makes them chemically compatible withless costly non-noble metal catalysts, and the use of very thinmembranes may overcome the problems associated with theirinherently low specific conductivity. While proton conductiv-ities of heterocyclic (especially imidazole, pyrazole, and triazolefunctionalized) systems are almost insensitive toward changesof the relative humidity (RH), it must be mentioned that, in thecase of phosphonic acid functionalized systems, some wateractivity is required to prevent the systems from condensation(Figure 2) which is strongly detrimental to protonconductivity.5 As expected from the related acidity increase,fluorination of phosphonic acid functionalized polymers leadsto a suppression of condensation and an increase ofhygroscopicity and proton conductivity.9

Interestingly, the problem of condensation is less severe forpure phosphoric acid, a highly viscous liquid with the highestintrinsic proton conductivity of any compound.10,11 Whilephosphoric acid fuel cells (PAFC) take advantage of theunique properties of phosphoric acid absorbed in some porousmatrixes, adducts of polybenzimidazole with phosphoric acidare used in high-T PEM fuel cells operating around T = 160°C.12 Under these conditions hydrogen rich reformates can beused as fuel without further purification. Although phosphoricacid loses much of its conductivity through the acid/base

interaction with the polymeric matrix, the remaining protonconductivity 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 foralkaline ion conducting electrolytes, e.g., Li ion conductors usedin lithium ion batteries (Figure 1c). Here, standard electrolytesare solutions of lithium salts (e.g., LiCF3SO3, LiPF6) in aproticpolar solvents such as ethylene or propylene carbonate (EC,PC) and dimethyl carbonate (DMC). Although the ionicconductivity has contributions from both cations and anions,the total conductivity is generally more than 1 order ofmagnitude lower than that of proton conducting membranesused in PEM-fuel cells. At the moment this is not a severeproblem, since the current densities of current batteries arelimited by poor electrode kinetics in most cases, and smallseparator thicknesses may compensate for low specificconductivities. There are also polymers, such as polyethylene-oxide, which dissolve Li-salts to form solid polymer electrolytes,2

but as in the case of proton conducting polymers, the ionicconductivities remain another order of magnitude below typicalconductivities of liquid systems. Therefore, most separatormaterials in Li-batteries contain some low molecular weightsolvents, even when this is trapped in gel-like structures withina stable polymer network. Nevertheless, fully polymeric,nonflammable separators are highly desirable because theiruse may solve the severe safety issues of alkaline ion batteries.

3. MEMBRANES FOR PEM-FUEL CELLSPolymer electrolyte membrane (PEM) fuel cells mostly usehydrogen or hydrogen-rich reformates as fuel, and oxygen isusually supplied as a humidified air stream within a temperaturerange constrained by the properties of the membrane material.In the case of Nafion, this is limited to about T = 90 °C becauseof the increasing hydration requirement with increasingtemperature, the decay of the morphological stability in thistemperature range, and the membrane decomposition throughthe attack of peroxo and hydroxo radicals. These radicals format the surface of the platinum electrocalalyst in the presence ofhydrogen and oxygen. Especially at OCV (open circuitpotential) conditions, highly dispersed platinum of the cathodestructure oxidizes, dissolves into the acidic membrane, andmigrates toward the anode side until precipitation occurs withinthe membrane when hydrogen crossing over from the anodeside is encountered.13 Among others, the main researchobjectives for this class of ion conducting polymers thereforeare reduced humidity dependence of the proton conductivity,reduced acidity, lower gas permeation, and improved chemicaland mechanical stability.When hydroxide ion (OH−) conducting anion exchange

membranes are used, lifetime and performance are also limitedby the membrane properties. Such fuel cells can currentlyoperate up to a temperature of only T ∼ 60 °C. As a strongnucleophile, OH− ions start to attack the chemical bonding ofthe anion exchanging group (quarternary ammonia or someother amine) of the polymeric structure, while reaction withacidic CO2 from the air converts OH− into bicarbonate(HCO3

−), which shows little electrochemical activity (seeSection 3.2).In the case of HT-PEM fuel cells operating at T ∼ 160 °C,

the commonly used PBI−phosphoric acid membranes havedemonstrated stable performance for 104 hours, but phosphoricacid leaching upon start/stop cycles and inferior mechanicalproperties at the required high phosphoric acid “doping” levels

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



are still serious downsides relevant for their use in fuel cells.Other disadvantages are the required high platinum loading ofthe oxygen electrode and the very high activation potential ofthe oxygen reduction reaction (ORR). Relatively low currentdensities actually let the fuel cell potential drop into the rangearound U = 0.7 V which is about 60% of the OCV at theoperation temperature, and the related efficiency loss may notbe acceptable when compared to other technologies.3.1. Systems Containing Sulfonic Acid. As indicated in

the Introduction, the behavior of ion conducting polymers infuel cells is affected by transport and stability issues on differentlength and time scales. Since our current understanding of thiscomplex situation is far from complete, I will focus on somerecent advances in the field in an attempt to approach a moreconsistent comprehension of the complex relationships.3.1.1. PFSA Membranes and Their Modifications. Per-

fluorosulfonic acid (PFSA) membranes generally consist of apolytetrafluorethylene (PTFE) backbone with perfluorinatedside chains of different lengths attached to the backbonethrough ether linkages and terminated by sulfonic acid(−SO3H) groups (Scheme 1). The unique properties of

PFSA membranes are most likely the immediate consequencesof the PTFE backbone to pack in an ordered way (similar tothe crystallization of pure PTFE (Teflon)), the high persistencelength of the backbone, and the fact that the polymericstructure combines the extreme hydrophobicity of the back-bone with the extreme hydrophilicity of the superacidic −SO3Hgroup (note that the latter is related to the electron withdrawalproperty of the adjacent CF2 group). Even at low RH, asignificant number of water molecules are absorbed at the−SO3H groups (the amount of water is usually expressed ashydration number λ = [H2O]/[−SO3H]), and a hydrophobic/hydrophilic phase separated morphology develops sponta-neously. The absorbed water tends to strongly bind togetherforming a continuous aqueous domain in which not onlyproton and water transport but also transport of dissolved gases(especially oxygen) take place. The accessible polymerconformations constrain the separation to the nanoscale asevidenced by the appearance of an “ionomer peak” in small-angle X-ray scattering (SAXS) patterns, as first observed byGierke et al.14

Nanomorphology. Surprisingly, details of the nanomorphol-ogy of PFSAs were under debate for a long time until a recentwork by Schmidt-Rohr and Chen who obtained seemingly thefirst quantitative picture of Nafion’s morphology fromsimulating a previously published small-angle X-ray scattering(SAXS) pattern.15 The so-called “parallel cylinder model” is nowwidely accepted and apparently supported by a recent NMRstudy on elongated membranes.16 The key feature of thismicrostructure is inverted micelle cylinders with large diameterseven at low water contents (e.g., 2.4 nm for a water volume

fraction Φwater as low as 0.2 (20%) corresponding to a relativehumidity (RH) of about 80%). But our own work17 raisedsevere doubts about this model: a constant number of cylindersrequires significant structural reorganization to adjust changesof the water content which is in contrast to extremely fastequilibration once water has entered the membrane.18,19 It isalso the severe separation of different and accumulation ofequal charges (e.g., protonic charge carriers accumulatingwithin the interior of the cylinders) which is energeticallyunfavorable when the charges are not completely screened bythe water of hydration. We have therefore suggested the waterstructures to be locally f lat and narrow which allows theprotonic charge carriers to electrostatically interact with severalsulfonic groups.17 Generally speaking, water “f ilms” may act aspositively charged “glue”, keeping together the oppositelycharged polymer structures. In the case of Nafion, this is notonly evidenced by the linear scaling of the structural correlationlength (obtained from the position of the ionomer peak) withthe polymer volume fraction but also by reasonable fits of theregime of SAXS patterns representing correlations in the 1−5nm range using water volume fractions, determined exper-imentally with a high precision. While the parallel cylindermodel was most likely biased by a large uncertainty of theexperimentally determined water content, locally flat morphol-ogies are not only consistent with the evolution of NafionSAXS patterns recorded over a wide range of water contents(up to a water volume fraction Φwater = 1 − Φpolymer ∼ 0.5), butthey also seem to be a common feature of most dissociatedionomers and polyelectrolytes (Figure 3). Details of this

nanomorphology may be different (in the case of Nafion, theremay be water films between polymer ribbons, as suggested bythe Grenoble group15), but in any case, the water films arerelatively narrow (0.3−2 nm) which has important implicationsfor transport, the visco-elastic behavior at high T and lowhumidification, and the evolution of the nanomorphology withT and RH.In long side chain PFSAs with high IEC, the hydrophobic/

hydrophilic separation is very robust; i.e., it is still observed at

Scheme 1. Molecular Structures of (a) Nafion, (b) Aquivion,and (c) the 3M-Ionomer

Figure 3. Linear scaling of structural correlation length d with thepolymer volume fraction Φpolymer as obtained from small-angle X-rayscattering (SAXS) of Nafion17 (inset), diverse sulfonated poly-ether-ketones (S-PEK), and a fully sulfonated poly-phenylene-sulfone (S-220) which indicate locally flat morphologies as illustrated for a highlyswollen situation (reprinted with permission; copyright 2013 Wiley-VCH).



very high T (up to T = 160 °C) provided RH is not too low.20

The observation that this morphological feature tends to decayat low RH (the ionomer peak is almost lost at T = 160 °C andRH = 10%) is consistent with electrostatics driving theformation of slightly ordered flat morphologies. It is expectedthat, with decreasing RH and increasing T, which dramaticallyreduces the dielectric constant of water and therefore itssolvating properties, dissociation of the ionomer, precedingelectrostatic cross-linking, is reduced. A short side chain (SSC)Aquivion membrane with the same ion exchange capacity asNafion actually loses its structural correlations at a lowertemperature around T = 100 °C for RH = 10%,20 which pointstoward other competing interactions involved in the formationof the nanomorphology, as will be explained below.Another typical structural feature of all PFSA membranes is a

certain degree of crystallinity arising from PTFE backboneordering. Of course, this “crystallinity” is more pronounced forSSC PFSAs in which the properties are more dominated by thebackbone and continuously decreases with increasing IEC, i.e.,with increasing side chain concentration. The internal structureof the ordered parts shows up as a correlation peak in theWAXS (wide-angle X-ray scattering) regime,21 and thestructural correlation between these parts is indicated by theso-called “matrix knee” which follows the ionomer peak towardlower q corresponding to correlation lengths around 15 nm.This feature appears to be quite robust compared to the verylow q part of the SAXS patterns. It is clear that even on theselarge scales (>30 nm) there is still some structure, but thecorresponding features are strongly affected by the samplehistory (including swelling, deswelling, aging, stretching, andpressing). A q−1 slope, as reported by the Grenoble group,22

indicates the existence of elongated objects, but the observationof slopes ranging from 0 to ∼3.517 (depending on the samplepretreatment and aging) strongly suggests that the PFSAstructure on this length scale (∼30−100 nm) is subject tosevere changes. These changes also seem to include the onsetof the so-called “ultrasmall angle upturn”, a distinct increase ofthe scattered intensity at very low q.15,22

Frankly speaking, there is still no satisfactory understandingof the PFSA structure neither on the molecular (sub-nanometer) scale nor on very large scales. There is onlyreasonable evidence for a few features on the nanometer scale(see above). Fortunately, these features are those most relevantfor understanding transport properties and, to some extent, alsothe visco-elastic behavior of PFSA membranes.Visco-Eleastic Properties. Apart from common stress/strain

measurements, the visco-elastic properties of PFSA membranesare generally examined by complex dynamical mechanicalanalysis (DMA).There are actually two major trends in the evolution of the

storage modulus (real part of the elastic constant): For relativehumidity values in the range RH ∼ 10−100%, water is acting asa plasticizer; i.e., the storage modulus monotonically decreaseswith increasing hydration level, and at any relative humidity,there is a severe decay of the storage modulus with temperatureup to T ∼ 140 °C (Figure 4). Surprisingly, the latter is notvisible in the SAXS regime where T dependent scattering dataare available (d = 1 − 30 nm). For a given RH, the SAXSpatterns are virtually independent of T. Only at significantlyhigher T and low RH does the intensity of the ionomer peakdecrease,20 and this perfectly correlates with another feature inthe visco-elastic behavior. In this regime (high T, low RH),water actually stif fens the membrane (Figure 4), and the decay

of the corresponding interaction shows up as a maximum of theelastic loss (tan δ in Figure 4) shifting toward higher T withincreasing RH. The interaction, fading away in this T, RHregime, is obviously identical to the interaction driving theformation of the locally flat nanomorphology (see above)which decays in the same regime. This presumably electrostaticcross-linking only accounts for a minor part of the total storagemodulus at lower T, and the decay of the modulus alreadyvisible at room temperature is most likely associated with thedecay of another interaction. This cannot be the physical cross-linking through the crystalline parts of PFSA membranes,because the corresponding structural correlation (matrix peakin the SAXS) remains unchanged far beyond T = 100 °C,20 andalso DSC (differential scanning calorimetry) shows an onset of“fusion” of the ordered parts only above T = 200 °C.23,24 Thereare actually two features observed in the DSC trace, and bothhave been interpreted as endothermal fusion. While this isundisputed for the distinct peak in the T range of PTFEmelting (>300 °C) only visible for low IEC PFSAs, the secondfeature just above T = 200 °C looks more like indicating a glasstransition. If this holds true, the corresponding glass transitiontemperature for the hydrated ionic form is expected to be lowerthan for the PFSA precursors (containing SO2F as nonionicprecursory of the sulfonic acid group) used for the DSCmeasurements. The decay of the elastic modulus withtemperature observed experimentally (Figure 4) is thereforemost likely associated with a glass transition of the amorphouspart. The fact that this shifts to higher T by ∼30 K for SSCPFSA compared to Nafion (inset Figure 4) suggestsdisintegration of the hydrophobic backbone interaction in thisregime.In short: apart from electrostatic interactions driving the

nanomorphology and leading to some ionic cross-linking,hydrophobic backbone interaction (probably in the amorphousphase) seems to govern the overall visco-elastic properties.

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



Hydration Behavior. The hydration of PFSA membranesaffects virtually all properties: transport coefficients of ions andwater and even the permeation of dissolved gases generallyincrease with the level of hydration,25 and the selectivity forexchanging cations or anions (perm-selectivity) is lost while theelastic properties usually deteriorate with increasing hydration(see above). These dependences therefore critically controltheir performance not only in PEM fuel cells but also in redox-flow batteries, as will be discussed later.Apparently, the main force driving water into PFSA

membranes is the extreme hydrophilicity of the super-acidicsulfonic acid group (−SO3H). Hydration is exothermal, and theheat of hydration decreases with increasing water uptake; in thecase of Nafion the hydration enthalpy has been measureddirectly26 by calorimetry showing that about six watermolecules per sulfonic acid group are absorbed exothermally.Accordingly, the water uptake in this regime is expected todecrease with increasing T, which recently has been provenexperimentally (Figure 5a).27 From this decrease as a functionof RH, the heat of hydration as a function of λ has beenobtained and found to be close to that of sulfuric acid. The factthat hydration isotherms for most sulfonic acid functionalizedionomers and polyelectrolytes fall into a very narrow range forlow RH, corresponding to water uptakes of λ < 6, clearlysupports the assumption that hydration in this regime isgoverned by the hydrophilicity of the sulfonic acid group. Theobserved small variations may not only result from smalldifferences of the acidity but also from hydration entropies,which depend on the degree of water dispersion. In fact, Nafionshows the highest water uptake in this regime (low RH), whichcorrelates with the fact that Nafion undergoes a verypronounced hydrophobic/hydrophilic separation leading todisordered 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 andtheir conjugated base (−SO3

−), there is much more water takenup by PFSA membranes at high RH (up to λ ∼ 20, Figure 5).The corresponding driving force has been suggested to be theentropy increase associated with the dilution of the protoniccharge carriers within the hydrated hydrophilic domain.27 Thisis nothing but osmosis, and considering the high concentrationof hydrated protons within the hydrophilic domain, the osmoticpressure may reach values limited by the correspondingcounter-pressure built up within the polymer matrix. Withthe storage modulus shown in Figure 4, internal pressures up to12 MPa (120 bar) have been calculated. While the relatedequilibrium water content is very close to the experimental oneat room temperature, the severe decay of the elastic moduluswith T leads one to expect a decreasing internal pressure andaccordingly a severe increase of hydration (swelling) with Tespecially at high water activity (RH). Indeed, there is someincrease in hydration observed, but this increase remains belowwhat is expected from the significant decrease of the bulkmodulus. This is true in the vapor phase, i.e., at RH below100%; once the dew point is reached, and water condenses onthe membrane surface and water uptake abruptly increases(Figure 5b). Observing two different water contents at thesame water activity apparently violates the Gibbs phase rulewhich is why this phenomenon is known as “Schroeder’sparadox”. But this may be resolved by introducing the surfacetension as a free parameter which has been done in variousimplicit and explicit ways.28−30 Recently we have suggested theformation of an “extended layered surface skin”, which is very

tough parallel and little permeable to water normal to themembrane surface. The surfaces of PFSA membranes are well-known to be hydrophobic in the vapor phase, and theformation of such an extended surface skin may start at thehydrophobic membrane surface under the same driving forceswhich govern the formation of Nafion’s weakly ordered layerednanomorphology (see above). As a simple consequence of theboundary condition (flat hydrophobic surface), the periodicwater “films” may arrange parallel to the membrane surface inthe near surface region. There is experimental evidence for theformation of such a highly anisotropic surface layer from gazingincident SAXS,31 and a thickness of ∼60 nm is estimated from astudy of very thin Nafion layers.32 This “skin” may then sustaina higher inner pressure than expected from the bulk modulus,and changes of the anisotropic “skin” structure in contact withliquid water are suggested as a means to release most of theinternal pressure (insets Figure 5b). This may lead to adecrease of the chemical potential of water and therefore to an

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



uptake of additional water, taking the system back to localthermodynamic equilibrium. It is also worth mentioning thatthe reorganization of the membrane surface dramaticallychanges the water exchange rate with the environment; incontact with water, PFSA membranes equilibrate about 2−3orders of magnitude faster than in humid vapor.27

Transport Properties (Water Transport and ProtonConductivity). Since the transport properties of Nafion havebeen discussed extensively in another review,25 here the majorfeatures are summarized with a focus on a few recent insightswhich have altered our general understanding of transport inPFSA membranes.Most of the transport takes place within the hydrated

hydrophilic domain of PFSA membranes, and the majorparameters to be considered are the morphology, the ionexchange capacity, and the hydration behavior, which is relatedto the viscoelastic properties (see above). The local dynamics ofwater exothermally absorbed at low RH is retarded to someextent because this water is essentially involved in ion solvation.But, interestingly, the sulfonic acid anion (−SO3

−) and thehydrated proton start to separate far before the ionic solvationshells are completed. For the model system methyl sulfonic acid(CH3SO3H)−water (H2O), 85% of the acid is found to bedissociated at a water content as low as λ = 2.33 At thishydration level, the ion concentration of the model systempasses through a maximum, and the striking observation thatthe diffusion coefficients of all species of the system passthrough a minimum at the same water content has beenexplained by an ordering of the incompletely screened ions inorder to minimize the electrostatic energy. Any molecular andionic transport then requires some extra thermally activateddisordering, and the overall transport rate is reduced.Translating this observation to the transport within hydratedhydrophilic domains of PFSA membranes, some local protonconductivity is expected even at such low water levels, but withhigher activation enthalpies than those of water diffusion andproton mobility in pure water. Activation enthalpies for localdiffusion and conduction processes are not available, but theanticipated behavior is also visible in the T dependence of themacroscopic water diffusion and proton conductivity (Figure6): the activation enthalpy of water diffusion decreases withincreasing water content approaching the value of pure water at

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

expected, the conductivity diffusion coefficient Dσ (protonmobility) closely follows the water diffusion coefficientespecially at low water content (Figure 7); i.e., as in acidicaqueous solutions, proton conductivity is brought about by avehicle mechanism, the cooperative diffusion of protonated andunprotonated molecules (here water).35,36 With increasingwater content, there is an additional conductivity contributionfrom structure dif fusion the prevailing proton conductionmechanism in pure water and dilute acids.37−39 Since thismechanism comprises intermolecular proton transfer, the terms“proton hopping” and “Grotthuss mechanism” are also used. Butuse of the term “structure dif fusion” may be more expedientbecause the protonic charge carrier just follows the hydrogenbond pattern (structure) of the solvent (water) around theexcess proton diffusing by hydrogen bond breaking andforming processes.39 In other words, this is a solvent drivenprocess involving the dynamics of many molecules displacingthe position of the protonic defect by just one molecularseparation. In particular, there is no indication for the formationof extended “Grotthuss chains” as still suggested in sometextbooks. It also should be mentioned that solvent and acidicprotons interchange their identity on a sub-picosecond timescale; 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 protonconductivity of Nafion as a function of hydration number λ:53 withincreasing λ, the corresponding values for pure water are approached.

Figure 7. Proton and water transport coefficients as a function ofwater volume fraction Φwater: DH2O (water tracer diffusion), Dσ (protonconductivity diffusion coefficient), DFick (diffusion in water concen-tration gradient), and Dp (pressure driven permeation diffusioncoefficient).25



It is the very nature of structure dif fusion that it is very sensitiveto symmetry reduction.39 In acidic solutions, this is caused bythe effect of neighboring ions biasing the hydrogen bondswithin the solvation structure of the hydrated acidic proton;within the separated morphology of an ionomer, such asPFSAs, also confinement effects reduce symmetry and thereforethe rate of structure diffusion. In PFSA membranes, confine-ment effects are relatively small, i.e., the decrease of structuredif fusion with decreasing water content resembles that of acidicaqueous solutions (e.g., ref 34).When it comes to understanding water transport and the

formation of water concentration profiles within PFSAmembranes in running fuel cells, the effect of internal pressuregradients (see section on hydration) on molecular transportmust be considered as well. In PFSA membranes, the internalpressure gradient is a driving force for hydrodynamic waterflow, which is a fundamentally different mechanism than thediffusional transport processes discussed above. Water flows ina collective way when the water structures are sufficiently wide,which is the case in PFSA membranes at high levels ofhydration. This leads to the remarkable situation that collectivepressure driven water transport increases substantially whileFickian water diffusion decreases close to the dew point ofwater (Figure 7) because of the vanishing chemical drivingforce (chemical water potential gradient). The interested readermay find the corresponding data in ref 25 and in ref 40 as a wayto implement these two different water transport processes(diffusional and hydrodynamic) into the transport equations.Interestingly, the dramatic onset of hydrodynamic watertransport is also clearly visible in the so-called electroosmoticwater drag, defined as the number of water molecules perproton dragged through the membrane as a consequence of aprotonic current. Of course, this is a key feature for anyelectrochemical application involving high current densities,which is true not only for fuel cells but also for redox-flowbatteries. In the case of Nafion, the electroosmotic dragcoefficient at room temperature is in the range of 2 below thedew point of water, but at higher swelling (Φwater > 30−40%),e.g., at higher T or in the presence of methanol,41 the dragcoefficient K increases dramatically (Figure 8). This transitionactually corresponds to a width of the flat water films around1.5 nm17 which seems to correspond to what is called “slip” influid dynamics.42

In today’s PEM-fuel cells, operating close to the dew point ofwater, hydrodynamic water transport is of paramountimportance for describing water distribution within PFSAmembranes. Apart from the boundary conditions (gashumidification, membrane surface hydrophilicity), this isessentially determined by the electroosmotic water drag fromthe anode to the cathode and the pressure driven waterpermeation (Figure 7) in the opposite direction.40

Within the family of PFSA membranes, the hydrophobic/hydrophilic separation of the SSC varieties is slightly lesspronounced which leads to slightly lower hydrodynamic watertransport for a given water volume fraction,40 but practical SSCmembranes usually have higher IECs than Nafion with higherwater uptake at a given RH, undoing this advantage.Since direct methanol fuel cells (DMFC) is still of interest in

PEM fuel cell technology, it must be mentioned that the above-discussed properties severely change in the presence of alcoholssuch as methanol. With their hydrophobic alkyl residue andtheir polar OH group, alcohols tend to behave like surfactants;i.e., they mediate interaction of the PTFE backbone with the

hydration water leading to more swelling.41 Severe increase ofhydrodynamic solvent (water and methanol) and gas transportis an immediate consequence,25 which is frequently called“crossover”. The tracer diffusion coefficient of methanol is only afactor of 2 lower compared to the water diffusion coefficient25

and electroosmotic drag coefficients are virtually identical,41 butthe presence of any alcohol dramatically reduces dissociation ofthe sulfonic acid group (−SO3H) and therefore also protonconductivity.25,41

Stability. Frequent swelling and deswelling, severe dryingclose to sealings, fuel cell stack mounting pressures of morethan 2 MPa, and the extreme chemical conditions in PEM fuelcells require the proton conducting polymer membrane to bemechanically and chemically robust.As opposed to many other membrane types, PFSA

membranes are indeed mechanically very robust, althoughtheir modulus is comparatively low. This has to do with theirextreme fracture toughness (elongation to break) characteristicfor ionomers with PTFE backbone. The severe softening athigh T and RH (Figure 4), however, clearly limits the operationwindow of PFSA membranes, and the higher glass transitiontemperature of SSC PFSAs (see above and inset of Figure 4) istherefore a distinct advantage.When it comes to chemical durability, SSC PFSAs seem to

have a slight advantage over the long side chain varietyNafion.43 This especially applies to the degradation of the sidechains, which starts from the C−S bond and is significantlyfaster than backbone degradation.44 Prior to this 19F-NMRinvestigation of membranes subjected to fuel cell test protocols,the so-called unzipping reaction starting from COOH groupsterminating the backbone had been identified;45,46 fortunately,this reaction can be suppressed by ensuring that the mainchains are terminated by −CF3 groups. The occurrence of theobserved reactions is actually the consequence of the presenceof peroxo- and/or hydroxo-radicals which form on the surfaceof the electrocatalysts in the presence of both oxygen (O2) andhydrogen (H2).

47 Therefore, gas permeability has to be

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



included into the stability considerations. At high hydrationlevels, the water in PFSA membranes dissolves comparativelyhigh concentrations of gases which are transported with thehydration water. This actually explains why PFSA degradationis highest under OCV (open circuit potential) conditions whenno protonic current and therefore no electroosmotic water fluxfrom the anode to the cathode prevents oxygen frompermeating from the cathode to the anode side.Just by considering the chemical structure of PFSA

membranes, one may also anticipate breaking of ether linkageswithin the side chains through acidic attack, but to the best ofmy knowledge there is no convincing experimental indicationfor this. Long-term hydrothermal aging experiments at T = 80°C and 100% RH show a continuous decrease of the IEC,48 butthis can be completely reversed through the usual stand-ardization procedure (e.g., 2 h in 1 M of a strong acid at T = 80°C). The authors actually found spectroscopic indication forcondensation reactions between sulfonic acid groups, but itcould well be that slow structural relaxation driven by residualelectrostatic interaction (see above) may also have preventedthe membrane from rapidly exchanging ions in the titrationprocess used to determine the IEC.3.1.2. Hydrocarbon versus PFSA Membranes. Since the

early 1990s, there has been continuously high interest inhydrocarbon membranes as potential alternatives for the wellestablished PFSA membranes. Originally research and en-gineering activities were driven by the high costs of PFSAmembranes and the putative environmental problems asso-ciated with the use of perfluorinated polymers in massproducts. Later, it became clear that hydrocarbon membraneshave distinctly different properties;49 i.e., they may be evenmore suitable than PFSAs for certain applications.Statistically Sulfonated Polyarylenes. Because of their

superior stability, initially polyarylenes have been generallychosen as starting polymers for polymer analogous sulfonation.Direct electrophilic sulfonation, e.g., in fuming sulfuric acid, wasvery popular, but soon also nucleophilic substitution routeswere used for the sulfonation of electron poor base polymers.The interested reader may refer to the review by Rikukawa andSanui,50 a chapter in the Fuel Cell Handbook,41 and theexcellent recent review by Hickner et al.51 For thesepreparation routes, sulfonation is essentially random alongthe polymer chain as it is in PFSAs, but the properties of thecast membranes differ from these of PFSA membranes in acharacteristic way.49

For similar volume densities of sulfonic acid groups (e.g., asulfonated polyarylene-ether-ketone with an IEC of 1.4 mequivg−1 corresponds to Nafion with an IEC of 0.9 mequiv g−1) anda given water volume fraction, SAXS patterns of sulfonatedpolyarylenes and PFSA membranes are similar in the regime ofthe ionomer peak, again supporting the hypothesis that someresidual electrostatic interaction between the hydrated ionsdrives the nanomorphology (see above). The influence of thebackbone properties still shows up as a slight but systematicincrease of the scattering intensity within the Porod regime anda small shift of the ionomer peak toward higher q, along withsome broadening.49 This points toward a less pronouncedhydrophobic/hydrophilic separation which naturally explainsthe significantly lower permeation coefficients for water anddissolved gases. The effect on the diffusional transport (waterdiffusion and proton conductivity) is insignificant at highhydration levels, but with decreasing hydration the decay ofwater diffusion and proton conductivity is more severe for

statistically sulfonated polyarylenes than for PFSAs.25,41 Acloser inspection of the diffusional transport coefficientsincluding their T-dependence shows that both transport onthe nanoscale and long-range diffusion are retarded insulfonated polyarylenes compared to PFSA membranes. Thisclearly suggests that the aqueous nanostructures are not onlynarrower but also less connected with a higher tortuosity onlarger scales. As expected from this nanomorphology, the onsetof excessive electroosmotic water drag is at distinctly higherwater volume fractions. The interested reader may find acomprehensive set of transport data in ref 25.Statistically sulfonated polyarylenes are ionomers just as

PFSAs, and that is why their visco-elastic properties are stilldetermined by the nature of their polymer backbones. Incontrast to low IEC PFSAs (see above), these are generallyamorphous and swell more in water which makes them verysoft under wet conditions with a low fracture toughness (smallelongation to break), while they are brittle in the dry state.Blending with a small concentration of compatible unsulfo-nated high molecular weight polymer actually leads to a distinctimprovement of the visco-elastic properties without changingthe transport properties too much.49

The mode of sulfonation has actually severe implications onthe chemical durability which is critical with respect to acid−base and redox reactions. For electrophilic sulfonation, the ruleof the thumb “easy on−easy off” applies to the hydrolyticdesulfonation as described and explained in ref 52. Not onlycan electron rich polyarylenes be easily sulfonated anddesulfonated, but they are also very susceptible against oxidativeattack, i.e., through H2O2 and even more so by oxidizingradicals. The fact that several electrophilically sulfonatedpolyarylenes survived in operating fuel cells for more than1000 h has probably to do with the limited operationtemperature (T < 80 °C) and the fact that the highly reducedgas crossover compared to PFSA membranes efficientlyprevents radical formation.

Sulfonated Polyarylenes with Block-Structure. The abovequalitative analysis naturally guides the way toward the currentapproaches for overcoming the main disadvantages of hydro-carbon membranes, namely, the steep conductivity decreasewith decreasing level of hydration, the susceptibility towardhydrolytic and oxidative attack, and the poor mechanicalstability, 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 thecharge carrier concentration and the diffusion coefficient ofwater, greatly influencing the ionic mobility. The latter hasmostly to do with the fact that the number of water moleculesabsorbed per ionic group is approximately identical for allionomers and polyelectrolytes in this regime (see above); i.e.,the water uptake is approximately proportional to the IEC. Thepercolation within the aqueous domain, however, increases in ahighly nonlinear fashion. As a result, any increase in IECdramatically increases the conductivity at low RH where waterpercolation is still low. This has been demonstrated by theextremely high conductivity of fully sulfonated polyphenylene-sulfone53,54 (Figure 2) and poly-para-phenylene.55 According tothe IUPAC nomenclature,2 these are polyelectrolytes with salt-like properties: they are extremely brittle in the dry state andsoluble in water. The high concentration of ionic groups makesany phase separation redundant; i.e., the hydration water ishighly dispersed within a morphology which is well organized



on a low nanometer scale, as evidenced by distinct X-raydiffraction patterns.53

Nevertheless, they can be used as a constituent of morecomplex polymeric structures. If the polyelectrolyte partseparates to form a continuous domain, the conductivity canbe expected to be higher than for similar materials with arandom distribution of sulfonic acid group simply because ofthe highly nonlinear increase of the local conductivity with localIEC (see above).An elegant way to form such a material is through so-called

block-co-polymers, an approach pioneered by severalgroups.56−60 Here, typically segments of a highly sulfonatedpolyarylene with a defined molecular weight of a few thousandg·mol−1 are combined with unsulfonated segments of a similardefined molecular weight through a coupling group so as toform di-, tri-, or even multi-block copolymers. Of course, suchsupramolecular structures may phase separate in a defined waywhen cast from solutions of polar aprotic solvents preferentiallyas thin membranes. This phase separation is then typically on ascale of a few tens of nanometers and should not be confusedwith the separation observed in ionomers on smaller scalesbetween water and polymer. In fact, the latter is observed as theinternal structure of the hydrophilic domain with typicallymuch lower correlation lengths than observed for ionomers(inset Figure 9).53

If the phase separation is well developed, as the one shown inFigure 9, the properties of the polyelectrolyte part are locallypreserved. Such membranes may then combine in a unique waysome of the high conductivity of the hydrated domain and thelow hydrodynamic water transport with the mechanicalproperties of the water free preferentially elastic domain. Ofcourse, this is only the case if the morphology is bicontinuous.Then acceptable compromises between transport and mechan-ical properties may still allow for proton conductivities higherthan that of Nafion (Figure 10). As in the case of purepolyelectrolytes, the transport of water and protonic charge

carriers is subjected to higher activation enthalpies,59 i.e., suchmembranes 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 correspondingmultiblock copolymers. Compared to PFSA membranes, theytherefore take up more water at high relative humidity and theexchange of water with the vapor phase is generally fast (seesection on hydration behavior of PFSA membranes).No compromise can be made with respect to hydrolytic and

oxidative stability of the sulfonated segments because, in allPEM fuel cells, they are in intimate contact with water,oxidizing gases, and radicals. A very useful way toward stablesulfonated segments is through the use of presulfonatedmonomers in the polymerization reaction, an approach whichhas been extensively promoted by the McGrath group.51 Thisallows one to sulfonate electron rich phenyl rings in the usualway before reducing the electron density, and therefore thereactivity, 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 FunctionalizedSystems. The operation of PEM fuel cells at high temperatureand low humidification is a critical issue, and thereforeunderstanding the physicochemical conductivity limits ofsulfonic acid based systems under these conditions is ofparticular interest. Since the values of the hydration number λof sulfonic acid functionalized systems fall into a very narrowrange for low water activities (RH < 65% corresponding to λ ∼6), the transport coefficients as a function of λ are informativeparameters. In this regime, the hydration water is involved inion solvation, retarding both water diffusion and protonconductivity (see above) . This also becomes visible as anincrease of the corresponding activation enthalpies withdecreasing hydration number λ (see also Figure 6). Comparingactivation 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 confinementeffects must be considered as well. For a given hydrationnumber, the activation enthalpies for water and protontransport, which are indicative of the local dynamics, exhibitthe lowest values for Nafionthe ionomer with the most

Figure 9. Morphology of a multiblock-copolymer (SU14−FS15consisting of blocks of fully sulfonated phenylene sulfones and blocksof phenylene ether sulfones with an overall IEC of 1.6 mequiv/g) asobserved by transmission electron microscopy.59 The inset shows thecorresponding SAXS pattern compared to this of S-220 (see text)(reprinted with permission; copryight 2012 Wiley-VCH).

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



developed phase separation. On the other hand, waterdispersion depends on the nature of the backbone and theIEC; i.e., it is higher for PFSA than for hydrocarbonmembranes49 and generally increases with increasing IEC.40,53

A high IEC actually leads to a higher concentration of protoniccharge carriers and a high water uptake in terms of watervolume fraction, which is associated with a more efficientpercolation of the hydrated hydrophilic domain.These are actually reasons for the very high proton

conductivity of high IEC SSC PFSAs such as DOW 858 (seeFigure 2 and ref 40) or the 3M membrane62 and explains whyhigh IEC hydrocarbons such as fully sulfonated-polyphenylene-sulfone (S-220) show proton conductivities comparable to thatof Nafion at room temperature.53 At this temperature, theretardation of transport on the local scale is compensated bythe better percolation on larger scales and the higher chargecarrier concentration. Because of the higher activation enthalpy,high IEC hydrocarbon systems may show increasingly higherproton conductivity than Nafion with increasing T.Since conductivity increase with IEC is highly nonlinear, the

accumulation of ionic groups into a distinct well percolatingphase is a suitable strategy to maximize proton conductivity.The physicochemical limit can then be estimated byconsidering the water diffusion coefficient and protonconductivity for aqueous solutions of small sulfonatedmolecules such as methyl-sulfonic acid (MSA).33 Here,percolation and confinement effects are virtually nonexistant,while the local dynamics is expected to resemble the situationin well phase separated membranes. The comparison of theconductivity of the MSA−H2O system with that of Nafionreveals a conductivity exceeding the conductivity of Nafion at λ= 3 (RH ∼ 30%) by a factor of ∼40 (Figure 11). For real

membranes with nonconducting hydrophobic volume incre-ments and some residual confinement effects, a possibleconductivity 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 IECPFSAs62 or hydrocarbon membranes, e.g., through appropriateblending, or by dispersing nonsoluble high IEC nanoparticles.In this context, preparation routes which allow the placement ofmore than one sulfonic acid group on the phenyl rings of

hydrocarbon polymers is of paramount importance, since thisprovides an additional means to increase the local density ofionic groups.63 At what density of ionic groups emergingpolyelectrolyte effects start to reduce dissociation has yet to beclarified, but in any case, the ionic groups must be concentratedin 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 tobe avoided, since this dramatically reduces both the local andlong-range transport coefficients. On the other hand themaximum swelling of such high IEC materials has to be limitedthrough a corresponding toughness of the polymer matrix. Ifthe maximum swelling were constrained to λ = 5, the hydrationlevel and proton conductivity may even become independent ofRH for RH > 50%. For such materials, the detrimental effects ofswelling and deswelling on the proton conductivity areexpected to be small while the conductivity at high RHremains reasonably high.It goes without saying that optimum membranes must be

morphologically and chemically stable. As part of membraneelectrode assemblies, they experience an external uniaxialpressure in the range 2−5 MPa within typical PEM fuel cellstacks, while fluctuations of the internal swelling pressure inspace and time lead to additional strain. The response of themembranes to the related forces is not only controlled by theirstorage modulus but also by their fracture toughness(elongation to break). Although PFSA membranes are quitesoft at T = 90 °C (see Figure 4), their enormous elongation tobreak (∼150%) prevents mechanical failure. On the other hand,most hydrocarbon membranes have a significantly higherstorage modulus under these conditions, provided hydration isnot too high. But for low hydration levels, elongation to breakis very small, i.e., the membranes are brittle; high hydrationlevels lead to softening with no significant increase ofelongation to break. For this reason the most frequent failuremode of hydrocarbon membranes is the formation of cracksand pinholes under the conditions of a running fuel cell. Theincrease of the membrane fracture toughness over a large RHrange must therefore be one of the targets of future membranedevelopments.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 specificconditions. As pointed out above, the most important types ofreaction are the attack by oxo and hydroxo radicals anddesulfonation through hydrolytic cleavage of the C−S bond.Because the intrinsic susceptibility of this bond does not varysignificantly with the chemical environment, steric shieldingseems to be the most efficient way to prevent this bond frombeing attacked. This is probably the reason why short side chainPFSA membranes are chemically more robust than the longside chain ionomer Nafion.43,44 Interestingly, the formation ofradicals depends on a membrane property, namely, the gascrossover, which is significantly higher for PFSA compared tohydrocarbon membranes. Therefore, hydrocarbon membraneshave the potential to perform stably in PEM fuel cells, eventhough virtually all of them fail Fenton’s test. In particular,membranes with locally high IECs desirable for obtaining highproton conductivity (see above) generally have a highdispersion of water and therefore a low dissolution andtransport of gases.

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



When it comes to the desulfonation reaction, the generalobservation is: “easy on easy off”; i.e., polymers which are easilysulfonated are easily desulfonated at high water activities. Ofcourse, this has to do with the free energy of the commontransition state of both reaction pathways, and in the case ofpolyphenylenes this can be increased by reducing the electrondensity of the sulfonated phenyl ring.52 The use of sulfone(−SO2−) as a strong acceptor group in ortho position to thesulfonic acid group has not only been proven to be particularlyefficient in increasing the hydrolytic durability52 but addition-ally reduces the susceptibility against radical attack.64

3.2. Anion Exchange Membranes (AEM). For variousreasons, anion exchange membranes are not simply mirrorimages of above-discussed cation exchange membranes. Themost common anion exchanging groupsquaternized aminessuch 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

sulfonic acid group is an acid (as constituent of PFSAs the latteris even a superacid). Furthermore, the hydroxide ion (OH−), asthe anion of choice for AEM fuel cell applications, not onlytends to react with acidic gases, it also behaves as a strongnucleophile in general.The OH− conductivities reported in the early literature are

significantly lower than the very high proton conductivities ofsulfonic acid functionalized systems, but there is increasingevidence that this is the consequence of CO2 contaminationleading to the conversion of OH− into carbonate (CO3

2−) andeventually bicarbonate species (HCO3

−).With a careful exclusion of carbon dioxide, OH− con-

ductivities close to the proton conductivities of PFSAmembranes have recently been reported,65 and a quantitativedetermination of water diffusion and hydroxide conductivity ofun-cross-linked polyarylene functionalized with quarternaryammonium groups (supplied by Funtech under the trade nameFAA-3) reveals clear similarities with transport in protonexchange membranes (Figure 12).66 At intermediate watercontents (Φwater ∼ 20−30%), the OH− conductivity diffusioncoefficient Dσ follows the H2O diffusion coefficient; i.e., theionic groups are highly dissociated and hydrated OH− diffusesat a similar rate as the water molecules. With increasing watercontent, the occurrence of OH− structure diffusion is evidencedby the fact that Dσ is exceeding the water diffusion coefficientapproaching the value for pure aqueous solutions (5 × 10−5

cm2 s−1) at very high water contents. As expected from thelower rate of structure diffusion in pure water, the hydroxidemobility in this regime is about a factor of 2 lower than themobility of protonic charge carriers in acidic membranes. Atvery low water contents (Φwater < 20 vol %, λ < 10), themoderate basicity of the anion exchanging group shows up as asevere decrease of Dσ compared to the water diffusion

coefficient; i.e., counterion (OH−) condensation at theconjugated acid (TMA+) excludes an increasing part of theanions from being transported within the aqueous domain. It isimportant to note that, in this regime, the water uptake in termsof λ at a given RH is lower than for sulfonic acid functionalizedsystems 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 cellapplications65 for which a maximum power density of 823 mWcm−2 has been reported.67

The reason why anion exchange membranes are not yetestablished in PEM fuel cell technology is their intrinsicchemical instability and the above-mentioned susceptibility toCO2 contamination. More than 20 years ago, Bauer et al.already determined the reaction rates of nucleophilicsubstitution and Hofmann elimination of the amine groupthrough OH− ions.68 In the first case, the attack is at the carbonnext to the amine nitrogen, while in the second case thereaction is initiated by attack of a hydrogen in the β position tothe leaving group (e.g., NR4

+) (Scheme 3). In both cases, amineis released leaving behind an alcohol group in the first case or aCC double bond and a water molecule as a second leavinggroup in the case of Hofmann elimination. It also should bementioned that even the commonly used polyarylene back-bones may decompose by reacting with hydroxide ions.69 As

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

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

Scheme 3. Splitting Off Amine Functional Groups through(a) Nucleophilic Substitution and (b) HofmannElimination68



opposed to the desulfonation of cation exchange membraneswhich can be suppressed by acceptor substitution (see above),the opposite is true for anion exchange membranes, in whichthe 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 beensuccessful as they trade increased stability through delocaliza-tion with reduced stability from dramatically reduced stericshielding due to their planar geometry.66 The rapiddecomposition of polybenzimidazole (which has a largerdelocalization range then imidazolium and guanidinium) inalkaline media70 strongly suggests that delocalization alonecannot sufficiently stabilize a quaternary amine.66 Currently, themost promising approach for stabilizing quaternary aminesseems to be through the use of spacers coiling around theammonium group, thereby achieving a high degree of stericshielding in a comparatively simple way.71 The spacer mayeither separate the amine from the polymer backbone or maybe tethered to the amine with a methyl as the other terminalgroup.Currently, claims about base stable quaternary amines are

abundant,72 but so far, the results could not be verified byindependent studies.73−75 From this literature, current anionexchange membranes in their OH− form appear to beintrinsically 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, apreselection through testing small molecules yet appears to bean efficient strategy to search for suitable candidates. On thebasis of the systematic work of Schwesinger et al.76 this has ledto the identification of various phosphonium cations as aputatively base stable anion exchanging group.69

Apart from these chemical instability issues, the susceptibilityto CO2 contamination seems to be a no-go for fuel cellapplications in air. In this environment, any anion exchangemembrane converts from its OH− into the HCO3

− form withinminutes66 leading to a dramatic decrease of the conductivity.This is not only the consequence of the approximately fivetimes lower mobility of HCO3

− compared to the rate of OH−

structure diffusion in aqueous solutions; anion exchangemembranes in the HCO3

− form also take up less water at agiven relative humidity.66

Interestingly, the oxygen reduction reaction taking place atthe 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 anelectrochemical reaction even in the presence of CO2. But thisreaction is associated with a large anodic overpotential, whichshows up as a severe voltage drop even at low currentdensities.77 At high current densities and CO2 levels < 1000ppm, as is the case for operation in air (400 ppm), theproduction of OH− in the cathode reaction and the transport ofOH− from the cathode to the anode leads to a displacement ofcarbonate species in the membrane, which is well-known as theso-called “self-purging effect”.77 But at such high currentdensities (>500 mA cm−2), the fuel cell efficiency is alreadydown to about 25% which is too low for most types ofapplications.It should clearly be stated that the detrimental precipitation

of solid carbonates characteristic for KOH based fuel cells is notan issue in fuel cells making use of anion exchange membranes

and that there is still some power generation in the presence oflow levels of CO2, but at the moment anion exchangemembranes are still far from being utilizable in PEM fuel celltechnology (for up-to-date reviews see refs 78−80). At thisstage, however, they have an immediate potential for redox-flowbattery applications, as will be discussed later.

3.3. Adducts of Polybenzimidazole and PhosphoricAcid. A fundamentally different type of membranes withproton conductivities less dependent on external humidificationtakes advantage of the properties of phosphoric acid (H3PO4).As opposed to water, phosphoric acid is an intrinsic protonconductor as a consequence of its high degree of self-dissociation and a very high mobility of protonic chargecarriers.81,82 In the nominally dry state, self-dissociation leads tothe formation of a variety of charged species:

⇌ + + ++ − + −5H PO 2H PO H PO H O H P O3 4 4 4 2 4 3 2 2 72

The fact that their concentration (about 10 mol %81) and themobility of protonic charge carriers are very high is a directconsequence of the phosphoric acid hydrogen bond networktopology. Since this is key to the unique properties ofphosphoric acid based proton conducting materials in generala brief explanation of these relations is given here: With threeout of four oxygens being protonated, the phosphoric acidmolecule (H3PO4) has a severe imbalance of the number ofpotential proton donors and acceptors. When phosphoric acidmolecules condense to form a solid or a liquid the dominantintermolecular interaction is hydrogen bonding, indeed.Individual hydrogen bonds of the type O−H···O are actuallyvery short (<260 pm compared to ∼280 pm in water) andstrong, but since the unprotonated oxygen can accept twoprotons at the most within a hydrogen bond network, there isat least one potential donor (O−H) for which no protonacceptor is available. We have used the term “frustrated” forboth the non-hydrogen bonded proton and the hydrogen bondnetwork as a whole.83 Speaking in terms of degree ofprotonation OP, there are two distinct oxygen populationswithin the hydrogen bond network of phosphoric acid (Figure13) one population around OP = 0.25 (these are oxygensexclusively accepting protons) and another around OP = 0.9.85

But the latter extends to protonation degrees OP > 1.5 whichcorrespond to O−H groups which are involved in hydrogenbonding as acceptor but only little as proton donor (Figure 13).This “frustrated” high energy state is important (i) as aprecursor for the intrinsic condensation reaction, (ii) for theextreme hygroscopicity of phosphoric acid, and (iii) for the rapidproton 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 thehydrogen bond to the “frustrated” O−H, which cannot releaseits covalently bonded proton, leads to the formation of a watermolecule (H−O−H) and breaking of the O−P bond. Ofcourse, this reaction goes along with the formation of a bondbetween this phosphorus and an unprotonated oxygen ofanother phosphate (formation of a diphosphate) and thetransfer of an additional proton (preferentially a “frustrated”proton from another molecule) to the water molecule leadingto the formation of a hydronium ion (H3O

+). As aconsequence, condensation products (e.g., diphosphates) andaqueous species always coexist within phosphoric acid, whichhas 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



any higher humidity, phosphoric acid takes up more water; forany lower humidity, the concentration of condensationproducts increases.The driving force for the uptake of additional water, i.e., the

reason for phosphoric acid’s extreme hygroscopicity, isprobably the high hydrogen bond network “frustration” aswell. 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 blendinto the hydrogen bond network by forming strong hydrogenbonds with the phosphate species.85

Apart from the condensation reaction and the formation ofaqueous species, simple proton transfer between adjacentphosphoric acid molecules also leads to the formation ofcharged 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 electrostaticproton/proton coupling within the hydrogen bonded networkwith an otherwise almost random proton “rattling” dynamics.85

Once intermolecular proton transfer takes place, the hydrogenbond structure around the overprotonated phosphate(H4PO4

+) slightly contracts, thus preparing the next protontransfer step. With some delay time, this leads to the formationof a polarized hydrogen bond chain, which may comprise up toabout five phosphate moieties with a ∼20% probability.83 Incontrast to the situation in water, this is energetically cheapbecause the proton displacements in the very short hydrogenbonds and the corresponding dipolar moments are small, andthe dielectric response of the “solvent” environment, stabilizingthe polarization, is very fast because of its protonic nature(Zundel polarizability). Of course, a chain of polarizedhydrogen 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” thenegative partial charges appearing along the dipolar chainthrough forming hydrogen bonds; hence, the chain is

interrupted and reversal of the polarization process inhibited.The charges remain separated and the protonic charge carriersdiffuse in an almost uncorrelated way before two differentdefects start to attract and neutralize.82 This process leads tothe proton conductivity of nominally dry phosphoric acid(H3PO4) which is very sensitive to any perturbation (seebelow). The only additive, which increases the protonconductivity, is water, which is an intrinsic constituent ofphosphoric acid even in the nominally dry state (see above).The addition of small amounts of water does not change theprincipal conduction mechanism, but it appears to increase therate of the elementary reactions involved.85

In the following, proton conducting adducts of phosphoricacid and polybenzimidazole (PBI) are discussed in the light ofabove-described mechanisms. Initially, PBI has been chosenbecause of its high concentration of basic nitrogen sitesassociated with the benzimidazole moiety interacting with acidseither through complete proton transfer or just by hydrogenbonding.12 How much phosphoric acid is taken up by PBI verymuch depends on the kind of PBI and the formation process ofthe adduct. In fact, a huge number of different PBIs and otherbasic polymers have been reported to form adducts withphosphoric acid (for recent reviews, see refs 86−89). Becauseof its availability, initially, meta-PBI (Scheme 4a) was

commonly used before other types of PBI such as AB-PBI(Scheme 4b) were considered as well. The higher density ofbasic imidazole groups of the latter leads to a higher uptake ofphosphoric acid for given conditions,87 which is reported to beparticularly true for iso-AB-PBI, in which the benzimidazolemoieties are oriented in alternating head to head and tail to tailconfigurations (Scheme 4c).90 Replacing the phenyl ring inconventional PBI by the moderate base pyridine not onlyimproves the capability to take up phosphoric acid, it alsoincreases the solubility of para-PBI in polar aprotic organicsolvents,88 which is relevant for the film forming process. Sincepara-PBI shows a significantly higher tensile strength for a givenphosphoric acid content, pyridine containing para-PBIcombines high proton conductivity and mechanical robust-ness.87 Polybenzimidazoles containing ether linkages (Scheme4d) are also attracting increasing interest, not because of highphosphoric acid uptakes but for its high flexibility even at lowphosphoric acid contents.

Figure 13. Degree of oxygen protonation of a single H3PO4 moleculeand phosphoric acid in the condensed hydrogen bonded liquid state asobtained from ab initio MD simulation.85 Two populations are clearlyvisible corresponding to characteristic bond patterns (insets). Notethat the frustrated proton can be considered to be part of a waterprecursor (see text).

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



Compared to the dependence of phosphoric acid uptake onthe type of polybenzimidazole, the preparation process seemsto have an even stronger influence on how much phosphoricacid is kept within the diverse PBI structures. For a long time,PBI membranes cast from solutions of, e.g., DMAc were justimmersed in aqueous solutions of phosphoric acid, but later, aprocess allowing for a more homogeneous mixing of PBI andphosphoric acid was developed by Benicewicz et al.91 Thisprocess relies on the fact that both PBI and their precursormonomers are soluble in pyro-phosphoric acid (PPA). Thepolymerization process of various kinds of PBI is carried out athigh temperature under dry conditions in PPA, before castingthe homogeneous solution of PBI in PPA onto a glass plate atambient. In the course of cooling, PPA hydrolyzes, convertinginto ortho-phosphoric acid, which is a nonsolvent for PBI. As aconsequence, PBI and phosphoric acid separate during a sol/geltransition which leads to the formation of relatively stable films.Thermodynamically speaking, it is the non-negligible wateractivity of “ortho-phosphoric acid” which makes it a nonsolventfor PBI, and the gelification of the solution, therefore, dependson the relative humidity which is sometimes controlled at ahigh level during the casting process.86 The final products arecharacterized by high phosphoric acid contents which are wellkept within the gel structure, and adducts with a phosphoricacid content of close to 90 vol % still have a surprisingmechanical strength at ambient. Since the proton conductivityvery much depends on the phosphoric acid content (seebelow), this is considered a distinct advantage, but on the otherhand, the sol/gel transition may be at least partially reversedwhen the membrane is heated to the operation temperature ofa HT-PEM fuel cell (T ∼ 160 °C). The softening going alongwith the corresponding structural rearrangements is a well-known problem of membranes produced by the PPA process.Unfortunately, ionic and covalent cross-linking strategies88

appeared to be a trade off between improving tensile strengthand avoiding the appearance of brittleness. A generalrequirement for obtaining membranes with reasonablemechanical stability is a very high molecular weight of PBIwhile crystallinity is practically no issue. Because of theplasticizing effect of phosphoric acid, the weak order(crystallinity) of pure PBI is completely lost in PBI−phosphoricacid adducts.In any case, such materials are gels at not too low water

activity; i.e., liquid phosphoric acid is just entrapped within anetwork of PBI with one phosphoric acid per benzimidazoleunit strongly interacting through hydrogen bonding or pureCoulomb forces.92 Therefore, leaching of phosphoric acid inthe presence of water is always an issue of this type ofmembrane material.Nevertheless, the apparent conductivity of PBI−phosphoric

acid gels is significantly lower than the conductivity of bulkphosphoric acid. In order to achieve high proton conductivitysimilar to this of fully hydrated Nafion, the phosphoric acidcontent must be as high as ∼80 vol %.88 For this, theconductivity is already about a factor of 5 lower than for purephosphoric acid. Apart from the small reduction of percolation,the major effect explaining this dramatic conductivity decreasecould well be the decrease of hydrogen bond network“frustration” associated with the subtraction of protons throughpolybenzimidazole.92 As pointed out above, this is expected toreduce not only phosphoric acid’s proton conductivity in thenominally dry state but also its hygroscopicity, another factoraffecting proton conductivity (see above).

In short, adducts of polybenzimidazoles and phosphoric acidare gels under most conditions. For obtaining high protonconductivity, the materials need to contain a high concentrationof phosphoric acid, which naturally tends to soften the gelstructure.Since this is most likely the consequence of the strong

interaction between the benzimidazole unit and phosphoricacid (internal salt formation92), alternative approaches may alsoconsider other polymeric matrices different from the commonlyused polybenzimidazoles. From the above considerations,backbone flexibility and kind and density of basic sites areexpected to be sensitive parameters in narrowing downinteresting membrane compositions which may show highproton conductivity at lower phosphoric acid contents. Anotherimportant parameter is the nanomorphology: apart from thepercolation of the phosphoric acid phase, progressive phaseseparation is expected to reduce the number density ofinteracting sites which already has been shown to increaseproton conductivity.93 It is definitely true that PBI−phosphoricacid membranes are the only ones that have been usedsuccessfully in PEM fuel cells at high temperature (T ∼ 160°C) without external humidification, but the severe voltagedrop of the corresponding fuel cells at rather small currentsreduces the fuel efficiency leading to the production of largeamounts of waste heat which cannot always be used insatisfactory ways. This is generally thought to be theconsequence of the adsorption of phosphate species onto theplatinum electrocatalyst at the equilibrium potential of oxygenreduction. Another challenge for further developments, there-fore, is to bind phosphoric acid in such a way that it stillsupports high proton conductivity without adsorbing onplatinum at cathodic potentials. The fact that CsH2PO4 basedfuel cells show lower over-potentials for the oxygen reductionreactions94 provides some hope that there is space for furtherimprovements.

4. MEMBRANES FOR REDOX-FLOW BATTERIESRedox-flow batteries are related to both PEM fuel cells andbatteries (Figure 1). As other type of batteries, redox-flowbatteries store electrical energy almost reversibly. Theirelectrochemically active masses are not stationary like inconventional batteries, they rather are solutions of redoxcouples pumped through the electrode compartments just ashumidified fuel and air is channeled along the gas diffusionelectrodes of PEM fuel cells. The cell designs (also stackdesigns) therefore resemble this of fuel cells, but the separatormembrane requirements are quite different. The membrane hasto separate the so-called anolyte and catholyte solutions (Figure1b) which contain the electrochemically active redox couplesand some supporting electrolyte (for the different types ofredox couples, see, e.g., ref 95). When the flow battery ischarged or discharged, an equivalent amount of ionic chargehas to cross the membrane, while the ions involved in the redoxprocess have to be efficiently separated. Apart from the obviousstability requirements (the membrane has to be durable withthe oxidizing and reducing solutions) the most important issuestherefore are selectivity and transport.Unfortunately, there is virtually no systematic membrane

development for this very demanding application. But availablemembrane types have been tested in situ analyzing poweroutput, Coulomb efficiency, and the compositional changes ofanolyte and catholyte during operation. Generally speaking,PFSA membranes such as Nafion allow for the highest power



densities, especially when “forced convection” of the activeelectrolytes is applied,96 but low Coulomb efficiency andchanges of the ion concentration in the electrolyte point towarda severe crossover of electrochemically active species and water.Diverse modifications, such as the dispersion of inorganicnanoparticles (e.g., SiO2, TiO2), blending with other polymers(e.g., PVDF), or coating the membrane surface with a moreselective (e.g., polyethylenimine) layer, are claimed to improvethe membrane’s barrier function.97 More recently, the specificadvantages associated with the use of cation and anionexchanging polyarylene membranes have clearly been recog-nized. Compared to PFSA membranes, hydration water is moredispersed in these membranes,49 and this is thought to be thereason for the generally higher perm-selectivity. In the case ofvanadium redox flow batteries (VRFB), anion exchangemembranes have the advantage of excluding electrochemicallyactive VO2

+ from entering the membrane, which also reducesthe reaction rate with this highly oxidizing species.97 For thisapplication, even moderately basic anion exchange membranescan be used, because their functional groups (commonlynonquarternized simple amines) are fully protonated under theacidic conditions of VRFBs. Unfortunately, polyarylenes notonly suffer from limited oxidation stability, but their insufficientionic conductivity under the conditions of redox flow batteriesis 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 theVRFB using vanadium redox couples for both electrolytes(VO2+/VO2

+ and V3+/V2+). In this case, any crossover reducesthe Coulomb efficiency and increases the internal heatproduction but does not lead to irreversible chemicalcontaminations.A framework which may allow for a more systematic

development of membranes for particular redox-flow systems isbriefly 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 preventedfrom entering the membrane. Therefore, anion exchangemembranes may be preferred for redox-flow batteries makinguse of cations as redox-couples (e.g., VRFB), but in many casescation exchange membranes are chosen because of power andstability considerations. Fundamentally, the point at whichDonnan exclusion (perm-selectivity) starts to decay is expectedto depend on hydration level, kind of fixed and mobile ions,nanomorphology, and chemical nature of the backbone. At lowhydration levels, at which all water molecules are involved inexothermal ion solvation (see Section 3.1.1), counterions mayexchange for other ions of the same sign, but this is not possiblefor the fixed ions of opposite sign (Figure 14a). Even in theregime where entropy is the main driving force for membranehydration (see Section 3.1.1), the uptake of ions of the samesign as the fixed ions is unfavorable albeit not completelyexcluded (Figure 14b). How Donnan exclusion decays with thedegree of hydration is expected to depend not only on thesolvation energies of all ions involved but also on thenanomorphology and solvent interaction with the backbone.For the same hydration number λ (here, water molecules perion), the dielectric constant of a sulfonated polyether ketonewas found to be significantly lower than that of a PFSAmembrane99 clearly indicating that water dispersion andbackbone interaction enters into the solvation capability of

hydration water in ion exchange membranes. Very usefulinformation is therefore quantitative data on the decay ofDonnan exclusion with increasing hydration number λ whichmay help to identify the maximum hydration for the targetedion selectivity. Of course, the latter depends on the ionmobility, which, in aqueous environment, essentially dependson ion size and charge. This is probably the reason why cationexchange membranes still show reasonable barrier function inVRFB in which the electrochemically active ions are all cations.But before addressing the issue of selectivity between differentcations or anions, it should clearly be noted that the hydrationlevel of the membrane during operation must be controlled.This can be done by controlling the external conditions, inparticular the concentration of the supporting electrolyte, whichallows adjusting the osmotic pressure difference. Provided thereis no internal membrane pressure (see Section 3.1.1), the ionconcentration within the membrane will roughly match that ofthe electrolytes. The other strategy is reinforcing the membraneso 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 simpleelectrostatic considerations, here, also specific interaction suchas hydrogen bonding or even some covalency may entersolvation and ion pairing energies. The experimentaldetermination of Donnan equilibria and electronic structurecalculations of solvation and ion pairing energies are thereforehelpful. The latter is of particular interest for understandingcounterion condensation (condensation of fixed and “mobile”ion) excluding ions from being transported through themembrane. Apart from ionic transport, the transport of solvent(in most cases water) is another critical issue. This may be the

Figure 14. Schematic illustration of ion distribution in an anionexchange membrane: (a) in the Donnan exclusion regime, where onlyanions can exchange, and (b) beyond Donnan exclusion, where alsoco-ions can enter the membrane.



consequence of osmosis (driven by the chemical waterpotential difference of anolyte and catholyte) and electro-osmotic water drag (the coupled transport of ions and water;see Figure 8). As shown above (see Section 3.1.1), bothtransport modes dramatically increase beyond a certain level ofhydration, and it is not clear yet whether this level coincideswith the decay of Donnan exclusion. By all means, it isdefinitely 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 electrochemicalcharging and discharging processes. Since this definitelyincreases with increasing hydration, like in the case of ionconducting membranes for fuel cells, the challenge is to obtainhigh ionic transport at a hydration level still acceptable for therequired selectivity. This situation very much resembles this ofhigh temperature, low humidity operation of fuel cellmembranes, and an interesting approach therefore is tomaximize the local concentration of appropriately chosenionic moieties. For the strongly acidic cation exchanging group−SO3

− high proton conductivities have already beendemonstrated (see Section 3.1.2), and basic anion exchangingamine groups may allow for high OH− conductivity at highlevels of hydration. Exchanging OH− versus anions with lowerhydration enthalpy (e.g., Cl−), as commonly present inelectrolytes of redox-flow batteries, leads to a severe drop inconductivity even at high water contents. In the case of the Cl−

form of an anion exchange membrane, the conductivitydiffusion coefficient systematically remains a factor of 2 belowthe water diffusion coefficient (Figure 15). This points toward

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

5. ELECTROLYTES FOR ALKALINE ION BATTERIESAlthough safety issues ask for fully polymeric separatormaterials in alkaline ion batteries, fully polymeric alkaline ionconducting membranes hardly found their way into high drainbattery applications yet. Since it is beyond the scope of thisarticle to address the many specific aspects associated with thisparticular application, only a few general issues of the formationand 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

propylene and ethylene carbonate, dissolve high concentrationsof Li-salts, but these salts do not fully dissociate into singlesolvated ions. The presence of contact ion pairs and even tripleions100,101 and the appearance of correlated ionic diffu-sion102,103 are characteristic for such systems. The fact that itis the oxygen of the small solvent molecules coordinating to theLi ions led several groups to dissolve salts into polymeric ethers,in particular poly(ethylene oxide) (PEO), and study their ionicconductivity more than 40 years ago.104,105 The correspondingsolvation chemistry is characterized by strong Li+ coordinationthrough the ether oxygen while the anions are only weaklyinteracting with PEO.106 As a result, such electrolytes aremainly anion conductors with Li+ transference numbers evensmaller than for liquid electrolytes.107 It must, however, benoted that in the case of liquid electrolytes there is anotherpathway for Li-ion transfer: the counter diffusion of contact ionpairs and anions.108 Low transference numbers and absoluteconductivities more than 1 order of magnitude lower than theseof liquid systems (Figure 16) has led to the development of gel-

type electrolytes, in which liquid electrolytes are simplyentrapped 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 masseswhile providing conductivities close to these of liquidelectrolytes. A comprehensive comparison of gel-type andsolid polymer electrolytes is given in ref 109. Both types ofelectrolyte conduct anions and cations, but in the case of liquidsolutions, the ions have a stable solvation shell, while for salt inpolymer solutions, e.g., Li+, it must be transported without thepolymeric solvent. The strong chemical interaction betweenpolymer and Li+ then requires a dynamical coupling betweensegmental motion and Li+ diffusion as directly evidenced byNMR.110

What is still missing for battery technology is a highlyconducting fully polymeric single Li-ion conductor containingno flammable low molecular weight solvent. Recently, singleion conductors with very high conductivity (>10−3 S cm−1)have been obtained by ion exchanging the highly protonconducting polyelectrolyte S-220 (see Figure 10) with Li+

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

Figure 16. Comparison of the total conductivities of typical Li+

conductors: mixed conducting solution of a Li-salt in aprotic polarsolvent,118 solid polymer electrolyte (salt in polymer),109 single Li-ionconducting polyelectrolyte solvated with an aprotic polar solvent,111

and a fully polymeric single Li-ion conductor.115



before solvating with dimethyl-sulfoxide (DMSO) (Figure16).111 Interestingly, Li+ diffusion is about five times slowerthan the diffusion of the solvent, which points to either strongcounterion association (incomplete dissociation) or a retarda-tion of Li+ mobility as a consequence of its heavy solvation byDMSO.112 Since the anions are completely immobilized, noconcentration polarization effects113 (formation of saltconcentration gradients) can occur in such electrolytes, andthe choice of DMSO as a solvent makes them interestingcandidates for Li/oxygen batteries.114 This is actually surprising,because the high donor number (DN ∼ 30) renders DMSO’soxidation limit to be very low. The downside of DMSO’sflammability is resolved for a very recent triblock copolymerconsisting of two PEO blocks and a segment bearing Li-sulfamide as ionic group.115 This is a true polymeric Li+ singleion 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 futureresearch.

6. FINAL REMARKSAccording to the large disparities of their levels of development,the perspectives for future research appear to be very diversefor the different types of ion conducting membranes andelectrochemical applications, albeit a few general aspects arealso clearly identified.Different strategies toward high temperature, low humidity

operation in fuel cells, the combination of high conductivityand selectivity in redox flow batteries, and polymeric single ionconductors for battery applications are adumbrated above. In allcases, high mobility of ionic charge carriers is an issue which,for most liquid electrolytes, is closely related to the solventdiffusion coefficient and the viscosity through the Nernst−Einstein and Stokes−Einstein relationships. One of thechallenges for obtaining high ionic mobility in polymericenvironment is the effective decoupling of ionic diffusion fromthe dynamics of the polymeric structure. While this is possiblefor protonic charge carriers in some hydrogen bonded liquidssuch as phosphoric acid83 (see Section 3.3) and speciallydesigned polymers containing heterocycles4 or phosphonicacid5 as protogenic groups, similar mechanisms have not yetbeen described for other ions. The separation of ionic chargecarriers, i.e., the dissociation process, is closely related to ionsolvation, which needs to be better understood beyond simpleelectrostatic consideration (Debye−Huckel approach). A lotcan be learned from the electrochemistry of solutions andpolyelectrolytes,116 for which simple empirical concepts such as“hard, softacid, base” or Manning counterion condensationare aiming at a qualitative description of ionic interactions inparticular systems. But also spectroscopic tools, especiallyNMR,117 and ab initio electronic structure calculations maysurely contribute to a better understanding of specific chemicalinteractions in processes such as charge carrier formation andmobility. Since in many cases the ionic groups are part of thepolymeric structure, such interactions affect the polymerconformations as well. Because of the close relation betweenionic transport and nanomorphology, insights into how thesepolymer conformations constrain the formation of orderednanomorphologies must be another key area of future research.This is where electrochemistry and polymer chemistry meet

with experimental, simulation, and theoretical tools applied to

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

■ AUTHOR INFORMATIONCorresponding Author*E-mail: [email protected] authors declare no competing financial interest.Biography

After receiving his Diploma in Mineralogy at the University ofCologne, Klaus-Dieter Kreuer did a PhD in the department ofChemistry at the University of Stuttgart. As a fellow of the“Studienstiftung des Deutschen Volkes” he benefited from a researchstay at the California Institute of Technology (group of R. Vaughan)and a Max-Planck award allowed him to join the MassachusettsInstitute of Technology as a visiting scientist. Later, Klaus-DieterKreuer built an R&D group within a Swiss-German company (Endress& Hauser) before joining the Max-Planck-Institute for Solid StateResearch, where he assisted J. Maier in building his new department.Since 1990 Klaus-Dieter Kreuer is lecturing at the University ofStuttgart from which he received his Habilitation degree.

■ ACKNOWLEDGMENTSThe many critical discussions with J. Melchior, A. Wohlfarth,and M. Marino, technical assistance through A. Fuchs, andproof reading by A. Kuhn and J. Jackson (all MPI-FKF) aregratefully acknowledged.

■ DEDICATIONDedicated to my colleague and friend Per Jacobsson.

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