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10

Inorganic Solid Proton Conductors

Philippe Knauth and Maria Luisa Di Vona

In this chapter, we will recall principles of ionic conduction in inorganic solids, present

general considerations on inorganic solid proton conductors, and discuss their main types.

There are a few textbooks on solid state ionics discussing this topic [1–4], but none before

1990. Several excellent review articles including inorganic solid proton conductors can also

be found [5–10]. More than 1100 articles were published in international journals on the

topic of “solid proton conductors” since the year 2005, including organic proton conductors,

and more than 400 papers on proton-conducting perovskite oxides, which appear now to be

the leading area of research within inorganic proton conductors.

10.1 Fundamentals of Ionic Conduction in Inorganic Solids

The total electrical conductivity s of a solid is the sum of the partial conductivities of ionic

and electronic charge carriers:

s ¼Xi

qiui½i� ð10:1Þ

In this equation, qi is the charge (in C), ui the mobility (in cm2s�1V�1), and [i] the

concentration (in cm�3) of the charge carrier i. The concentrations can be transformed

intomolar fractions knowing themolarmass and density of the solid. FromEquation 10.1, it

is evident that two parameters can bemodified in order to increase the ionic conductivity of a

solid: the carrier concentrations and/or their mobilities.

Solid State Proton Conductors: Properties and Applications in Fuel Cells, First Edition.Edited by Philippe Knauth and Maria Luisa Di Vona.� 2012 John Wiley & Sons, Ltd. Published 2012 by John Wiley & Sons, Ltd.

10.1.1 Defect Concentrations

The concentration of ionic defects can be increased in several ways:

(1) Deviation from stoichiometry (Figure 10.1a and 10.1b). A reduction or oxidation of the

compound produces simultaneously electronic species and point defects, thereby

leading to mixed conduction.

(2) Doping(Figure10.1c).Theadditionof aliovalent impuritieswithfixedvalence requires the

generation of ionic defectswith opposite charge in order tomaintain bulk electroneutrality.

(3) Intrinsically disordered solids. Many solids pass through an order–disorder transition

as the temperature is increased (e.g., point 4 in Figure 10.1d) with the formation of

three-dimensional disorder (e.g., a b� a transition in silver iodide AgI). In other

cases, the disorder is limited to disordered planes (intercalation compounds, such as

Figure 10.1 Defect chemistry in the bulk of AgCl. From top to bottom: (a) charge carrierconcentration as function of compound stoichiometry; (b) phase diagram: composition range ofthe homogeneous compound (c) influence of doping byCd2þ ions; and (d) Arrhenius diagramofsilver vacancy concentration. Symbols are explained in the text [16]. Reprintedwith permissionfrom Nanoionics and Soft Materials Science by Joachim Maier, Copyright (2006) SpringerScience and Business Media.

372 Solid State Proton Conductors

Na-b-alumina) or one-dimensional channels (tunnel compounds). Likewise, amorphous

phases (inorganic glasses and polymers) present a high intrinsic disorder and, in certain

cases, also exhibit ionic conductivity.

(4) Formation of space charge regions in the vicinity of interfaces. The space charge forms

to counterbalance a plane of opposite charge at the interface core [11]. The resulting

space charge region can be blocking [12] or highly conducting [13], depending on if

charge carrier depletion or charge carrier enrichment is observed.

10.1.2 Defect Mobilities

The ionic mobility depends on numerous factors [8]. Themost important is the height of the

potential barrier that ions must overcome in order to pass from onewell to an adjacent well.

The barrier height depends on several factors, including the strain energy for the ion to

“squeeze” through the bottleneck, the polarizability of the lattice, and the electrostatic

interactions between the ion and its surroundings. Perhaps the easiest to visualize is the

strain energy, and one is therefore tempted to assume that solidswith the highest free volume

should exhibit the highest mobility. While this is often the case, as in short-circuit diffusion

at extended defects (e.g., grain boundaries and dislocations), there are many examples

where solids with smaller channels support higher mobilities than those with larger

channels, due to polarization effects. As rule of thumb, it is advantageous for mobile ions

to have a high polarizability, andmeans to be soft and deformable, whichmight be related to

an intermediate valence change during transport, such as is hypothesized in the case of Cuþ

ion transport.

10.1.3 Kr€oger–Vink Nomenclature

The commonly used nomenclature for the description of defect chemical reactions was

proposed in 1956 by Kr€oger and Vink [14]. Point defects are considered as dilute species,

Table 10.1 Proton conductivity s and activation energy E of important inorganic solid protonconductors. The temperature of measurement T is indicated. The water partial pressure istypically 1 bar below 100 �C and about 30mbar above 100 �C

Formula T/ �C s/(S cm�1) E/eV

H1þ 2/3Mg2/3Al11�2/3O17 25 5 10�6 0.2–0.3a-Zr(HPO4)2�H2O 25 1 10�5 0.3HUO2PO4�4H2O 25 4 10�3 0.3–0.4H3[PMo12O40]�29H2O 25 8 10�2 0.2–0.3H3[PW12O40]�29H2O 25 6 10�2 0.2–0.3H3OClO4 25 1 10�4 0.3Sb2O5�5.4H2O 25 7 10�3 0.2CsHSO4 150 1 10�2 0.3CsH2PO4 230 2 10�2 0.4Ba(Ca1.18Nb1.82)O9 200 1 10�2 0.5BaCe0.95Yb0.05O3 600 3 10�2 0.5SrCe0.95Yb0.05O3 900 1 10�2 0.6

Inorganic Solid Proton Conductors 373

and the ideal solid plays the role of the “solvent.” Several analogies can be found between

intrinsic defect formation and self-dissociation of water [4]:

(1) A pair of charged defects is formed, which are responsible for electrical conduction.

(2) Defect formation is thermally activated. A “mass action law” constant relates the defect

concentrations (or defect activities for concentrated species) and describes the defect

equilibrium.

(3) An acidity–basicity concept can be introduced [15].

In theKr€oger–Vink notation, the subscript indicates the site of a defect; the subscript i standsfor an interstitial site. The effective defect charge, relative to the ideal lattice, is written as

superscript: a dot (.) stands for a positive charge and a dash (0) for a negative one. The

vacancy is written as V. For example, Oi00 represents a doubly negative charged interstitial

oxide ion; V..

O is a doubly positive charged oxide ion vacancy. Bulk defect chemical

reactions must obey the following:

(1) Mass balance

(2) Balance of lattice sites

(3) Charge balance (bulk electroneutrality)

However, local deviations from electroneutrality may occur near interfaces (the formation

of space charge regions; see above).

In a binary ionic compound MþX�, three main types of intrinsic ionic disorder can be

generated by permutation of the elementary defects, ion vacancies, and interstitial ions. The

predominant disorder type depends mainly on the solid’s crystal structure and the ion sizes.

(1) Cation Frenkel disorder corresponds to the formation of a cation interstitial plus a

cation vacancy, in a solid MþX�:

MM þVi>M.

i þV0M ð10:2Þ

KFr is the Frenkel equilibrium constant:

KFr ¼ ½M.

i �½V0M� ¼ A1 expð�DFrH

�=kTÞ ð10:3ÞTerms in brackets represent the defect concentrations (molar fractions); DFrH

� is thestandard enthalpy of the Frenkel reaction, and A1 a prefactor containing an entropy

term. k is Boltzmann’s constant and T the absolute temperature. Obviously, interstitial

formation is easier for small ions and/or in relatively open lattices. Therefore, cationic

Frenkel disorder is more often found than the corresponding anionic Frenkel disorder,

because anions are generally larger.

(2) Anion Frenkel disorder (sometimes also called “anti-Frenkel disorder”) corresponds to

the formation of anion vacancies and interstitial anions and can be written:

XX þVi >X0i þV

.

X ð10:4Þ

KAFr ¼ ½X0i�½V.

X� ¼ A2 exp ð�DAFrH�=kTÞ ð10:5Þ

The brackets represent again molar fractions; DAFrH� is the standard enthalpy of the

anion Frenkel reaction and A2 is the prefactor.


Due to the small radii of silver and copper ions, silver halides (such as AgCl;

Figure 10.1) and copper halides present Frenkel disorder, but also certain anion

conductors with the relatively open fluorite-type lattice (e.g., CeO2 or ZrO2).

(3) Most solids present Schottky disorder, which corresponds to the coupled formation of

cation and anion vacancy pairs:

MM þXX>V0M þV

.

X þMX ð10:6ÞMX represents ions, which have been displaced to “new” surface or interface sites. KSch

is the Schottky equilibrium constant:

KSch ¼ ½V0M�½V.

X� ¼ A3 exp ð�DSchH�=kTÞ ð10:7Þ

DSchH� is the standard enthalpy of the Schottky reaction, and A3 is the prefactor.

Schottky disorder is largely found in dense crystal lattices. For example, close-packed

alkali halides (including NaCl) show Schottky-type disorder.

In addition to intrinsic ionic disorder by point defects, one must take into account intrinsic

electronic disorder by the creation of electron–hole pairs, which can be written:

0> h. þ e0 ð10:8Þ

This process can be thermally activated or due to photons. The excess electrons e0 are in theconduction band, whereas electron holes h� are located in the valence band. The tempera-

ture-dependent equilibrium constant of this reaction is:

KelðTÞ ¼ ½h.�½e0� ¼ A4 exp ð�Eg=kTÞ ð10:9ÞThe prefactor A4 contains the effectivemass of holes and electrons, while Eg is the band gap

energy of the compound. Intrinsic electron–hole pair formation and vacancy–interstitial

pair creation (Frenkel reaction) can both be represented in level diagrams, like those used in

solid state physics.

“Anti-site” disorder describes the interchange of ions between two sublattices. While

such exchange between cation and anion sites is not observed in binary ionic systems due to

trivial electrostatic reasons, “anti-site” disorder can be observed in ternary and higher order

compounds in which cations disorder between two different cation sublattices: this is, for

example, common in solids with the spinel structure.

Extrinsic disorder. As one can easily verify in Equations 10.2, 10.4, and 10.6, intrinsic

defect reactions do not modify the composition of the exactly stoichiometric solid (the

“Daltonide”). In addition to these intrinsic disorder types, one can also observe extrinsic

disorder due to composition changes of the solid: this can be related to the presence of foreign

ions (impurities or dopants) or a deviation from stoichiometry. The latter is induced by

chemical potential changes of one of the components, for example changes of oxygen partial

pressure in the case of oxides. Oxidation or reduction of a solid leads to deviations from

stoichiometry (corresponding to a “Berthollide”) and formation of both ionic and electronic

charge carriers. For example, electron holes and interstitial oxide ions are created simulta-

neously inaFrenkel-disorderedoxidebyanoxidation reactionat highoxygenpartial pressure:

1=2 O2 ðgÞþVi > O00i þ 2 h

. ð10:10Þ


Kox ¼ ½O00i �½h.�2PðO2Þ�1=2 ¼ A5 exp ð�DoxH

�=kTÞ ð10:11ÞKox is the equilibrium constant and DoxH

� is the standard enthalpy of the oxidation reaction.Excess electrons and oxide ionvacancies are formed simultaneously in the reduction reaction:

OO > 1=2 O2 ðgÞþV..

o þ 2e0 ð10:12Þ

Kred ¼ ½V..

o �½e0�2 PðO2Þ1=2 ¼ A6 exp ð�DredH�=kTÞ ð10:13Þ

Kred is the equilibrium constant, and DredH� is the standard enthalpy of reduction. As a

consequence, the p-type or n-type conductivity increases together with the deviation from

stoichiometry. Given the much larger mobility of electronic defects, important deviations

from stoichiometry will lead to predominant electronic conductivity (n-type for reduction

and p-type for oxidation). Observation of ionic conductivity in binary compounds is

possible only near stoichiometry.

At reduced temperatures, the electrical properties of solids are determined by impurities

(or dopants) near stoichiometry, while under reducing or oxidizing conditions, defects

associated with nonstoichiometry control them. The electrical properties eventually

become intrinsic when the temperature increases to sufficiently high values. Kr€oger andVink discussed various aspects of stoichiometry deviations in inorganic compounds and

developed diagrams that show defect concentrations as a function of the chemical potential

of the components. These diagrams can be much simplified under the assumption of only

two majority defects, according to the so-called Brouwer approximation [17]. A general

discussion of these phenomena can be found in excellent standard texts [3, 4, 14, 18].

10.1.4 Ionic Conduction in the Bulk: Hopping Model

Ionic conduction in inorganic solids is due to thermally activated ion hopping. Assuming

Boltzmann statistics, the diffusion coefficient Di is a function of the jump distance a, the

characteristic attempt frequency n0, and the Gibbs free energy of migration DmigrG¼DmigrH�TDmigrS.

Di ¼ g a2n0 exp ð�DmigrG=kTÞ ð10:14ÞThe factor g takes into account geometrical and so-called correlation effects. For example,

the backward jump of an ion has a slightly higher probability than the forward jump, but, on

the other hand, cooperative motion can lead to higher diffusion coefficients than isolated

jumps. In the classical model of ionic transport in solids, typical attempt frequencies n0 areof the order of 1013 s�1 and are often identified with the Debye frequency. Enthalpies of

migration, which determine the temperature dependence of ionic mobilities, show large

variations with typical values for solid ionic conductors between 0.2 and 2 eV (about

20–200 kJ/mol). Low values are observed for interstitial mechanisms, whereas vacancy

mechanisms are generally characterized by higher activation barriers.

The Nernst–Einstein equation relates the ionic mobility ui to the diffusion coefficient Di:

ui ¼ Di qi=kT ð10:15Þ


Using Equations 10.1, 10.14 and 10.15, the ionic conductivity can be expressed as:

sion ¼ q2ikT

i½ �ga2n0 exp DmigS

k

� �exp �DmigH

kT

� �ð10:16Þ

A general equation representing the ionic conductivity can thus be written with a prefactor

s0:

sion ¼ s0Texp �DactH

kT

� �ð10:17Þ

Most crystalline and amorphous fast ion conductors (the latter below their glass transition

temperature) satisfy this equation. The activation enthalpy DactH can contain different

contributions:

(1) If the concentration of mobile ionic defects is fixed by charged background impurities

and dopants, as in the case of Y-stabilized ZrO2 (YSZ, 2 [VO€]� [YZr0]), the activation

enthalpy DactH is equal to the defect migration enthalpy DmigrH (cf. curve 2 in

Figure 10.1d).

(2) Assuming a thermally activated defect creation, the carrier concentration [i] has a

temperature dependence, such as in Equations 10.3, 10.5 or 10.7:

½i� ¼ ½i�0 exp ð�DformH�=2kTÞ ð10:18Þ

The factor 2 comes from the fact that defect pairs are always formed. In this “intrinsic”

case, DHact is the sum of the defect migration and formation enthalpies (DactH¼DmigrH þ DformH

�/2; cf. curve 1 in Figure 10.1d).

(3) If deviations from stoichiometry are observed, the activation enthalpy can be related to

reaction enthalpies, such as the oxidation enthalpy for metal-deficient oxides (cf. Equa-

tion 10.11) or the reduction enthalpy for oxygen-deficient oxides (cf. Equation 10.13).

(4) Finally, defect association can be observed at low temperatures. This issue was

discussed early on by Lidiard [19]. Dreyfus and Nowick investigated point defect

association in doped sodium chloride [20]. In the case of NaCl, doping with divalent

cations requires the formation of sodium ion vacancies for charge compensation:

MCl2 þ 2NaNa > M.

Na þV0Na þ 2 NaCl ð10:19Þ

At low temperature, defects association due to Coulomb interactions between the

oppositely charged defects begins to predominate:

M.

Na þV0Na>ðM.

Na;V0NaÞ ð10:20Þ

where (MNa.,VNa

0) refers to a bonded dopant–vacancy pair. In the temperature domain

in which association occurs, the conductivity has an effective activation enthalpy

(DactH¼DmigrH þ DassH/2, where DHass is the association enthalpy; cf. curve 3 in

Figure 10.1d). There are other important experimental consequences from such associ-

ation. The formation of vacancy–dopant pairs leads to a weaker overall enhancement of

conductivity. The pairs act as dipoles and contribute to dielectric relaxation processes.

Haven established the presence of loss peaks due to defect pairing in 1953 [21].


In addition to the formation of defect pairs, Lidiard showed that long-range defect

interactions also play a role and adapted the Debye–H€uckel theory of aqueous electrolytesto the case of ionic crystals [19]. The electrostatic interactions are also the origin of the

frequency dependence of the ionic conductivity and dielectric permittivity of structurally

disordered solid electrolytes [22, 23].

10.2 General Considerations on Inorganic Solid Proton Conductors

Now that the basics of ionic and mixed conduction in solids have been recalled, several

general issues can be discussed concerning inorganic solid proton conductors. The first issue

is the fact that a certain amount of proton conduction is observed in many, especially

amorphous, hydrated inorganic solids, but the contribution of protons to conductivity is

usually small and not very reproducible, depending a lot on the experimental conditions,

such as relative humidity and temperature. We will discuss in this chapter only inorganic

solids, where stable, reasonably large, and reproducible proton conductivity is observed.

Some observations may actually reflect poor experimental conditions: on one hand, so-

called dry conditions may often present more water than anticipated, typically tens of ppm,

unless special precautions are applied [24]. On the other hand, full water saturation may

require very high water pressure (P(H2O)) and high enough temperature for equilibration in

reasonable time scales.

A second immediately related issue is the stability of solid proton conductors. At high

temperatures, proton conductivity generally decreases, due for example, to loss of

vehicular water, salt decomposition, or proton concentration decrease in oxides. Given

that the proton mobility increases with temperature, significant proton conduction is

generally observed over a relatively narrow temperature range. In fact, few inorganic solid

proton conductors show significant proton conductivity above 400 �C, mostly oxides, in

which protons can be considered point defects and can be discussed based on point defect

equilibria.

A third issue is the proton conduction mechanism [25]. The proton is extremely small

compared to other ions. This small size and “naked” positive charge lead to an extremely

high polarizing ability of the proton on its immediate environment, and unique transport

mechanisms can be discussed: (1) lone proton transport, also called proton hopping,

translocation, or the “Grotthuss mechanism,” in which the proton jumps between

relatively stationary proton-donor and proton-acceptor sites; and (2) carried proton

transport, also called “vehicular mechanism,” where the proton rides on a carrier

molecule, such as H2O or NH3. Proton transport includes transport by any complex ion

that carries protons. According to the vehicular mechanism, protons are transported via

ions with an ordinary radius, such as oxonium (hydronium) ions (H3Oþ , 0.14 nm),

ammonium ions NH4þ , or others. In this case, counter-diffusion of the vehicle is required

for stationary conduction [26].

Although the principle of the Grotthuss mechanism was postulated a very long time

ago [27], the discovery of the vehicular mechanism is quite recent [26]. The vehicular

mechanism ismainly restricted to solidswith open structures, such as layeredmaterials with

conduction planes. In the Grotthuss mechanism, the proton jumps according to a coopera-

tivemotion from one site to another, using a complex network of hydrogen bonds extending


over large domains of the proton-conducting crystal lattice. Short hydrogen bonds of the

O–H� � �O type have an oxygen–oxygen distance of about 0.25 nm. Lone proton “Grotthuss”

transport occurs for example in potassium dihydrogen phosphate and perovskite oxides.

Proton-carried “vehicular” transport is observed e.g., in oxonium beta-alumina and

hydrogen uranyl phosphate.

An ultimate issue concerns the determination of proton conductivity. Clearly, a simple

alternating current (a.c.) conductivity measurement is insufficient, as different species

might contribute to the total conductivity of a solid, such as electronic carriers. Furthermore,

given the importance of the presence ofwater for proton conduction, indicating conductivity

values without reference of relative humidity makes no sense. The real existence of proton

conduction must be verified by one of various means: (1) direct electrochemical transport

of hydrogen through a proton conductormembrane and hydrogen evolution at the cathode at

a rate given by Faraday’s law; (2) determination of the electromotive force of a gas

concentration cell, for example, with hydrogen or water vapor, corresponding to the Nernst

law, a reduction of electromotive force (e.m.f.) indicates a partial short-circuit of the cell by

partial electronic conductivity; (3) direct current (d.c.) conductivity measurements with

blocking (Cu and Au) and reversible electrodes (Pd black saturated with H2 gas); and

(4) isotope effects, replacing protons by deuterons.

10.2.1 Classification of Solid Proton Conductors

Inorganic solid proton conductors can be classified according to different criteria: (1) the

preparation methods (direct synthesis or ion exchange), and (2) the temperature ranges in

which proton conduction is observed. Up to 100 �C, the list of “low-temperature proton

conductors” includes materials where the presence of water plays a fundamental role,

including various oxide gels and acidic compounds. The upper-temperature limit of this

group corresponds to the fact that below 100 �C, solid proton conductors can be easily

maintained fully hydrated. Between 100 �C and 250 �C, “medium-temperature proton

conductors” include some stable compounds; in particular, hydrogen sulfates and hydrogen

phosphates have been reported. Above 300 �C, “high-temperature proton conductors”

include oxides, especially of the perovskite type. The absence of high-performance

proton-conducting solids in a “gap” temperature range between 200 �C and about 350 �Chas been already noticed by Norby in 1999 [7]. These phenomenological classifications are

very useful for applications and engineering, as the interested researchers can choose easily

the appropriate material for a particular temperature range, but are not particularly

rewarding from a fundamental point of view.

Other classification schemes are based on (3) the lattice structure of the solid and the

dimensionality of proton conduction, planar (2D) versus bulk (3D); and (4) the conduction

mechanism of protons. Low-temperature inorganic proton conductors are often amorphous

solids, while intermediate-temperature ones are often crystalline acid salts with phase

transitions. High-temperature proton conductors are, instead, generally crystalline oxides

with perovskite or pyrochlore structure, presenting an important amount of point defects,

such as oxide ion vacancies. The presence of hydrogen bonds is essential for the structural

chemistry of all proton-conducting solids.

A nice cartoon by Colomban and Novak shows a very simple, but illustrative, image of

three major types of inorganic solid proton conductors (Figure 10.2) [2].


The conduction mechanism of protons is immediately related to the crystal structure of

the solid: proton defects (excess protons and proton vacancies) can be discussed in

anhydrous compounds (e.g., in KH2PO4), vehicular proton transport in loosely packed

structures (e.g., beta-alumina), and quasi-liquid conduction (e.g., in gels) [2].

In the “defect mechanism” type, encountered in oxides at high temperature, protons

can be considered extrinsic point defects and treated according to the classical

phenomenological theory of solid ion conductors, recalled above.

The “loose packed structures” include special structure types. Two-dimensional proton

disorder has been first reported in structures such as b-alumina showing two-dimensional

conduction planes where protons can move. One-dimensional disorder is observed when

Figure 10.2 Main proton transport mechanisms: protons are represented by rabbits. In case (a),interstitial rabbits move in a dense crystal with some defects (orientational disorder of hippo-potamus and a “doping” elephant). In case (b), a high concentration of mobile species movesbetween trees, representing stable structural features (suchas spinel blocks inb-alumina). In case(c), protons move in a disordered environment (with some “impurity” cows between the sheep)and many possibilities for statistical jumps [2]. Reprinted with permission from ProtonConductors: Solids, membranes and gels - materials and devices by P. Colomban CambridgeUniversity Press, Cambridge, 1992.


channels exist in the structure that can lead to enhanced proton transport along these

channels. However, in one- and two-dimensional structure types, grain boundaries are

effectively a barrier to long-range conduction, because they interrupt the conduction planes

(or channels), and a decrease of the ionic-conducting domain size actually reduces the

overall proton conductivity of the solid.

In the “quasi-liquid state,” protons can rearrange amid many possible configurations

and present a very high mobility, like in bulk-disordered solid ion conductors, such as the

high-temperature phase a-AgI, where silver ions are statistically distributed over many

sites, connected by very flat potential profiles. This structure is observed above an

equilibrium phase transition. Similar phase transitions are also observed in fast proton

(or “superprotonic”) conductors, such as CsHSO4: in its high-temperature phase, high

ionic carrier concentrations are accompanied by high and isotropic proton mobilities.

We will choose this structural and mechanistic classification to present and discuss

subsequently 2D solid proton conductors and 3D solid proton conductors. The order of

magnitude of conductivity of some representative inorganic solid proton conductors

discussed in this chapter is shown in Figure 10.3.

10.3 Low-Dimensional Solid Proton Conductors: Layered and PorousStructures

10.3.1 b- and b00-Alumina-Type

b- and b00-aluminas are famous sodium ion conductors with conduction planes situated

between spinel blocks with aluminum ions occupying tetrahedral and octahedral sites in the

Figure 10.3 Conductivity of inorganic solid proton conductors discussed in this chapter. BCN:Ba(Ca1.18Nb1.82)O9; BCY: BaCe0.9Y0.1O3�d; CHS: CsHSO4; HBA: (H3O)2Al11O17; HMP:H3[PMo12O40].29H2O; HUP: HUO2PO4. 4H2O; HWP: H3[PW12O40].29H2O; KDP: KH2PO4;LaP: LaPO4; SCY: SrCe0.95Y0.05O3�d; and ZrP: Zr(HPO4)2.H2O. The water partial pressure istypically 1 bar below 100 �C and about 30mbar above 100 �C.


cubic close-packed array of oxide ions [28]. The local electroneutrality is thus not respected

in the ideal structurewith an excess of positive charge in the conduction planes and an excess

of negative charge around the spinel block. The formula of b-alumina can be written as

(1 þ x)Na2O�11Al2O3, because it usually contains an excess of Na2O with x around 0.3.

The nonstoichiometry is believed to reduce the electrical imbalance and therefore increases

the crystal stability. The significantly more stable b00-alumina structure is very similar, but it

contains monovalent and divalent ions on tetrahedral aluminum sites.

The sodium ions are exchangeable against many other monovalent and some divalent

cations. b- and b00-structures exhibit proton conductivity when Naþ ions are exchanged

against H3Oþ or NH4

þ [29]. The crystal structures are naturally very similar to that of the

parent crystal. However, the crystals tend to cleave during ion exchange, as oxonium and

ammonium ions are much larger than sodium ions. b- and b00-gallates and b- and

b00-ferrites have similar formulas and structures and also show proton conductivity, when

sodium ions are exchanged against oxonium or ammonium. Gallium analog solid proton

conductors can be obtained from potassium b-gallates 1.3 K2O�11Ga2O3 by ion exchange

for ammonium or oxonium cations [2]. Here, the spacing of the conduction planes is larger

than in b-alumina.

Only limited information exists concerning the proton conductivity of b- and

b00-structures and large discrepancies can be noted, probably because it is difficult to

obtain large enough single crystals and suitable electrodes. Polycrystalline samples show

in general a rather low conductivity, because of the large grain boundary resistances, at

least in part due to blocking of conduction planes between crystallites of different

orientations. The conductivity of oxonium b00-alumina has been measured with blocking

and nonblocking electrodes to be 5 10�3 S/cm at 25 �C [30], but much lower values have

also been reported (5 10�6 S/cm [31]). The partly ammonium-containing b00-compound

NH4(H3O)2/3Mg2/3Al31/3O17 attains a conductivity of 10�4 S/cm at room temperature [32].

Similar data have been obtained for gallates and ferrites with activation energies around

0.2–0.3 eV [2].

Proton conductivity takes place in the two-dimensional conduction planes, but some

uncertainty remains regarding whether protons are transported via structural diffusion or a

vehicular mechanism involving oxonium or ammonium ions. The proton conductivity

increases with temperature, until loss of water or ammonia occurs above 200 �C. Seriousproblems preventing practical application are the very anisotropic properties, the difficulty

to prepare ceramic electrolytes, and the subsequent cracking, caused by changes of

composition with temperature and atmosphere [2].

10.3.2 Layered Metal Hydrogen Phosphates

Many water-insoluble metal hydrogen phosphates can be obtained as layered compounds.

Some examples are a-MIV(HPO4)2�H2O (MIV¼Zr, Ti, Sn, Pb), g-MIVPO4�H2PO4�2H2O

(MIV¼Zr, Ti), and HUO2PO4�4H2O. Given the presence of acid HPO42� or H2PO4

�

groups, some metal hydrogen phosphates exhibit good proton transport and catalytic

properties. Therefore, these compounds have been intensely investigated in the last three

decades, and copious literature is now available [33].

Zirconium hydrogenphosphate Zr(HPO4)2�H2O, in the literature often simply called ZrP,

is a proton-conducting layered compound and the “father” of a large family of similar


compounds [34, 35]. The good proton conductivity at room temperature of hydrated

amorphous zirconium phosphate (0.1–6 10�3 S/cm) was already discovered in 1968 by

Alberti [36]. Two different layered crystalline structures called a and g can be formed. In the

a-structure, zirconium cations are octahedrally coordinated by six oxygens belonging to six

different R–PO3 groups, giving a layered compoundwith the general formula Zr(RPO3)2�nS(R¼OH, C6H4SO3H. . .), where S is a solvent or other intercalated molecule. The structure

of a-ZrP (R¼OH; Figure 10.4) is appropriate for proton transport, because pendant OH

groups extend into the interlayer region and form a hydrogen-bonded network with water.

This highly stable compound has been used as a secondary phase in composite proton-

conducting ionomer membranes based on Nafion (see Chapter 8).

The distance between neighboring phosphate groups (0.53 nm) is relatively large, so

water molecules play a role as bridges; consequently, the number and arrangement of water

molecules are important. At high temperature, above 300 �C, the complete condensation of

the original acid P–OH groups can be avoided if these groups are far from each other.

However, in this case, the activation energy for the proton jump from one group to the

nearest neighbor is high.

The transport mechanism in a-ZrP at room temperature is dominated by the surface

transport and the contribution of hydrated surface groups. The surface proton mobility is at

least four orders of magnitude larger than bulk proton mobility; the crystallinity also plays

an important role. At 20 �C and 90% relative humidity, the conductivity of polycrystalline

ZrP is in the order of 10�5 S/cm, but with oriented thin samples high conductivity values

(2 10�4 S/cm [38]) can be obtained. In addition, the conductivity of a-ZrP is highly

dependent on hydration, varying by two orders of magnitude as the relative humidity is

increased from 5% to 90%, and the activation energy decreases from above 0.5 to below

0.3 eV at 20 �C. These values appear high for a pure Grotthuss mechanism; consequently,

vehicular transport also seems important. In anhydrous materials, a phase transition at

220 �C causes a reduction of interlayer distance and distance between neighboring oxygens,

Figure 10.4 Crystal structure of a-Zr(HPO4)2�H2O [37]. Reprinted with permission from SolidState Ionics, Layered metal IV phosphonates, a large class of inorganic-organic proton con-ductors by G. Alberti, M. Casciola, 97, 1–4, 177–186 Copyright (1997) Elsevier Ltd.


so that the proton jump distance is also reduced; consistently, the activation energy is

reduced from 0.65 to 0.35 eV [2, 39].

The possibility of replacing partially or totally the O3P–OH groups of a-ZrP with O3P–R

groups and the O2P(OH)2 groups of ZrPO4�O2P(OH)2 with O2POHR or O2POR0R groups

opened the way to a large group of organic derivatives of both a- and g-zirconiumphosphates, only limited by the chemist’s imagination and his or her ability to synthesize

different phosphonic or phosphinic acids [40–42]. Very thin layers can be obtained from

colloidal dispersions. Given their good thermal stability, these compounds are promising for

use in some solid state electrochemical devices at medium temperature.

Proton motion in hydrogen uranyl phosphate (or hydrated uranyl-phosphoric acid

(HUP)) HUO2PO4�4H2O has already been discovered in 1938, when Beintema noticed

that “true vagabond ions are present in the water layer of the structure” [43]. High proton

conductivity around 4 10�3 S/cm at 20 �C was reported with an activation energy of

around 0.3 eV in the “superionic” phase [44, 45]. However, thermal stability is a particular

problem, as irreversible dehydration is observed already above 100 �C, depending on

atmospheric humidity [45], and condensation of phosphate groups occurs above 150 �C.The crystal structure consists of UO2PO4 planes separated by two-level water layers. Each

water molecule can participate in four hydrogen bonds; there are more hydrogen bond

sites than protons. The oxonium ions are statistically distributed with a liquid-like

disorder in the water layer, and there is also a considerable orientational disorder of the

PO4 tetrahedra. The proton conductivity is strongly anisotropic and the conduction

mechanism is complex, which is supposed to be due to the vehicle mechanism by

oxonium ions in the water layer with a complementary Grotthuss-type contribution

through the hydrogen bond network, similar to heteropolyacids. There is an equilibrium

phase transition around 0 �C, between the high-temperature proton-conducting phase and

the low-temperature ferroelectric phase, which presents a higher activation energy of

about 0.6 eV [46].

Other isostructural solids, such as HUO2AsO4�4H2O, have a similar crystal structure and

the transition into the proton-conducting tetragonal phase is observed around 30 �C. Theactivation energy at low temperature (0.7–0.8 eV) is in favor of a vehicular transport

mechanism; at higher temperature (0.2 eV), it is consistent with a large contribution of the

Grotthuss mechanism [47].

10.3.3 Micro- and Mesoporous Structures

Micro- and mesoporous structures might also be attractive for proton conduction. After

early studies on proton-conducting zeolites [48, 49], which did not reach sufficiently

high proton conductivity, highly ordered mesoporous silicas functionalized with sulfonic

acid groups, obtained by anchoring and oxidation of thiol groups, were studied [50].

Metal organic frameworks (MOFs) couple porosity, diversity, and crystallinity [51], but

MOFs chemistry focuses on intrinsic properties of the empty frameworks. The use of

guest molecules to control functions has been essentially unexamined. Na-3(2,4,6-

trihydroxy-1,3,5-benzenetrisulfonate) conducts protons in regular one-dimensional

pores lined with sulfonate groups. Proton conduction was modulated by the loading

of 1H-1,2,4-triazole within the pores and reached 5 10�4 S/cm at 50 �C in anhydrous

H2 [52].


10.4 Three-Dimensional Solid Proton Conductors: “Quasi-Liquid”Structures

10.4.1 Solid Acids

Some hydrated inorganic acids show outstanding proton conductivities at room temperature

and can be considered among the best proton-conducting solids.

So-called crystalline heteropolyacids exist in a series of hydrated phases (see also

Chapter 2); their stability depends strongly on temperature and relative humidity. The

structure is formed of so-called Keggin anion clusters with composition [PM12O40]3�,

where M¼Mo, W, Si [53]. Hydrated molybdo-phosphoric acid H3[PMo12O40]�29H2O,

stable above 70% relative humidity, shows a proton conductivity of 8 10�2 S/cmat 25 �Cand

an activation energy around 0.3 eV [54]. The related tungsto-phosphoric acidH3[PW12O40]�29H2O, stable above 80%RH, presents a cubic structure, in which diamond-type lattices of

Keggin-type heteropolyanions [PW12O40]3� and clustered oxonium ions are interpene-

trated. The protons are clustered in a complex hydrogen bond network, spreading over entire

domains of the crystal. Proton transport on such a network is apparently vehicular with

contributions by the Grotthuss mechanism. A major disadvantage of this kind of proton

conductors is related to the significant reduction of proton conductivity in lower humidity

conditions (typically below 70% RH), because dehydration occurs and the hydrogen bond

network gets lost.

Oxonium perchlorate H3OClO4 is an early example of a solid proton conductor with

conductivity above 10�4 S/cm at 25 �C, already discussed in 1973 [55] (see also Chapter 2).It presents a first-order phase transition around �25 �C with an increase of proton

conductivity by about one order of magnitude and a decrease of activation energy from

0.4 to 0.3 eV [2]. In the low-temperature monoclinic phase, ordered layers of oxonium and

perchlorate ions are present with a strong hydrogen bond network; in the high-temperature

orthorhombic phase, the oxonium ions are disordered in a “quasi-liquid” state with

appreciable orientational disorder of perchlorate ions. Ammonium perchlorate NH4ClO4

shows somewhat similar behavior, but the orthorhombic–cubic phase transition is at notably

higher temperature, around 240 �C, and the activation energies are much higher, consistent

with vehicular diffusion of ammonium ions.

10.4.2 Acid Salts

Acid salts include compounds with stoichiometry MHXO4, M3H(XO4)2, and M2H(XO4),

whereM¼ alkali metal or NH4;X¼ S, Se, P, As [56, 57].Many of thesematerials undergo a

phase transition above which there is an increase in proton conductivity by 2–3 orders of

magnitude, reaching values as high as 10�1 S/cm. Acid salts in the high-temperature phases

are often called “superprotonic conductors.”

Alkali hydrogen sulfates and selenates are well-known ferroelectric solids in their low-

temperature phases with a higher thermal stability for the selenates. They become

frequently proton conducting in their high-temperature phase with a phase transition

generally between 80 �C and 230 �C, depending on the alkali cation [58, 59]. High proton

conductivity in ferroelectric crystals CsHSO4 and CsHSeO4 was discovered in the early

1980s byBaranov and co-workers [60]. Acid salts have been proposed as electrolytes in fuel


cell applications, although they present different problems. A difficulty is the reduction of

sulphate groups by hydrogen; furthermore, CsHSO4 is also water-soluble and must be

protected from liquid water.

Two equilibrium phase transition temperatures of around 45 �C and 141 �C were

discovered in CsHSO4, which also depend on water partial pressure. The conductivity

increases dramatically, by several orders of magnitude, at the second phase transition; the

plastic high-temperature phase shows proton conductivity of about 10�2 S/cm. The

structure of the ambient temperature phase presents infinite HSO4� chains with strong

hydrogen bonds, whereas cyclic dimers exist in the intermediate temperature phase with

weaker hydrogen bonds and increased orientational disorder of anions. In the plastic high-

temperature phase, free rotation of anions is observed together with a quasi-liquid state of

both protons and cesium ions and a disordered hydrogen bond network (Figure 10.5; see also

Chapters 3, 4 and 7) [61, 62]. Many similar cases can be discussed: in ammonium hydrogen

selenate NH4HSeO4, the phase behavior becomes very complicated with phase transitions

depending obviously also on humidity and pressure [63].

Dihydrogen phosphates and arsenates are also well-known ferroelectric materials that

undergo phase transitions into high-temperature proton-conducting phases. Potassium dihy-

drogenphosphate KH2PO4 has been extensively studied [64]. In the low-temperature ferro-

electric phase, strong hydrogen bonds link the phosphate groups. The proton conductivity is

low, and the activation energy around 0.5–0.7 eV. In the high-temperature paraelectric phase,

theprotonsare statisticallydisordered.Theprotonconductivity increases toaround10�6 S/cm.

CsH2PO4 has been discussed as a viable electrolyte for intermediate temperature fuel

cells [62, 65]. A transition to a stable, high-conductivity phase was observed at around

230 �C, with a conductivity of 2 10�2 S/cm at 240 �C and an activation energy dropping to

about 0.4 eV. Without humidification, dehydration of CsH2PO4 occurs, but the hydration

process is apparently not responsible for the high conductivity at this temperature.

A superprotonic phase transition was also found in potassium dihydrogen phosphite

KH(PO3H) at a temperature of 132 �C, reaching a proton conductivity of 4� 10�3 S/cm

at 140 �C. The compound is apparently stable against reduction, adopts a monoclinic

structure at room temperature, and transforms to a cubic one in the superprotonic phase [66].

Figure 10.5 Crystal structure of CsHSO4. (a)Monoclinic phase II and (b) tetragonal phase III. In(b), each oxygen position has half occupancy and the protons are placed in the middle ofdisordered hydrogen bonds (dashed lines) [61]. Reprinted with permission from PhysicalReview B, Superprotonic phase transition of CsHSO4: A molecular dynamics simulation studyby C. R. I. Chisholm, Y. H. Jang, S. M. Haile, W. A. Goddard, 72, 13, 134103 Copyright (2005)American Physics Society.


10.4.3 Amorphous and Gelled Oxides and Hydroxides

Proton conductivity can be found in many gelled and amorphous hydrated oxides and

hydroxides [67]. Oxide gels can be considered hydrated oxidesMO�nH2O containing water

molecules inside the oxide network. Proton conduction properties are mainly due to surface

phenomena. Water adsorption and dissociation occur at the oxide–water interface giving a

fully hydroxylated surface. TheM–OH groups at the surface of the solid can react in acid or

base equilibria:

M-OHþH2O ¼ M-OH2þ þOH� ð10:21Þ

M-OHþH2O ¼ M-O� þH3Oþ ð10:22Þ

Acid dissociation is promoted by small and highly charged cations; base dissociation is

favored by large metal cations with low charge.

Due to the small size of colloidal particles, interface properties are very important and

oxide gels usually exhibit rather high proton conductivities (s > 10�6 S/cm), but the proton

conduction at the surface of oxide particles depends strongly on the relative humidity of the

atmosphere.

Antimonic acid Sb2O5�nH2O (n¼ 5–6) is a particularly well-investigated example.

Proton conductivity up to 7.5� 10�3 S/cm for Sb2O5�5.4H2O has been reported at 20 �Cin fully hydrated conditions [68]. The crystalline formH2Sb4O11�nH2O (n¼ 0–3) presents a

framework of edge- and corner-sharing SbO6 octahedra and large water- and oxonium-

containing channels. The Arrhenius-type temperature dependence of conductivity presents

two activation energies (0.7 eV below and 0.4 eV above 60 �C [69]).

Many other hydrated oxides show relatively good proton conductivity, including

V2O5�nH2O, In2O3�nH2O, SnO2�nH2O, and TiO2�nH2O [68]. The proton conductivity is

generally strongly dependent on the hydration level: typical room temperature values are

around 10�4 S/cm. The activation energies are generally in the order of 0.2–0.4 eV [2].Most

transition metal oxides are unreactive toward hydrogen below elevated temperatures, but

proton insertion can occur at ambient temperature. In the so-called hydrogen bronzes, such

as WO3, ReO3, or MoO3 [70], proton insertion gives mixed conducting oxides with

predominant n-type conductivity for charge compensation and color change.

10.5 Three-Dimensional Solid Proton Conductors: Defect Mechanismsin Oxides

The presence of protons in solid ZnO and their influence on electrical and optical properties

were already reported in the 1950s [71, 72]. The work of Stotz and Wagner in 1966 [73]

established the existence of protonic defects in wide band-gap oxides at high temperatures

and showed that some oxides, such as Cu2O, CoO, and NiO, exhibit a certain amount of

proton conductivity in a hydrogen–water atmosphere. Stotz and Wagner derived equations

for the solubility of water vapor and hydrogen in Cu2O, CoO, and NiO considering the point

defects involved in these equilibria [73]. The determination of n- and p-type conductivity as

a function ofwater partial pressure is a goodmethod to evaluate the amount of proton defects

in an oxide. In the meantime, protons have been evidenced in a large number of binary


oxides at high temperature, including TiO2 [74, 75], Al2O3 [76, 77], SiO2, Y2O3 [78],

CeO2 [79], andmany others. The existence of proton conductivity in nanocrystalline oxides

in wet atmosphere at low temperatures is probably related to open porosity [79]. Great

quantities of protons can be incorporated in oxideswith ReO3 structure, but the predominant

electronic conductivity hides usually the significant proton conductivity.

The systematic investigation of acceptor-doped oxides known for their moderate oxide

ion conductivity provided further evidence that these oxides can be proton conducting in

hydrogen-containing environments at moderate temperature. When protons dominate, the

defect-related properties, such as conductivity, sintering, creep, and corrosion, become

humidity dependent.

10.5.1 Perovskite-Type Oxides

The discovery of a range of possible applications, including hydrogen sensors, hydrogen

pumps, steam electrolysis, and fuel cells, jump-started this domain in 1980 [80], when

Takahashi and Iwahara reported proton conductivity at elevated temperature in Y-doped

SrCeO3 and in other perovskite oxides exposed to hydrogen- and/or water vapor-containing

atmospheres [81]. The negatively charged defect corresponding to the ionized acceptor

dopants (in Kr€oger–Vink notation, YCe0) can be compensated by positively charged oxide

ion vacancies (VO€) or by interstitial protons (Hi_).The highest conductivities are observed in oxides with perovskite-type structures with

cubic or slightly reduced symmetry [10, 82]. These materials support relatively high

temperature and low humidity. The general formula of these oxides can be written as

AB1�xMxO3�d, where alkaline earth elements, such as Ba, Sr, and Ca, occupy the A site;

tetravalent elements, usually Ce or Zr, occupy the B site; M is a rare earth dopant element

with x < 0.1; and d is the oxygen deficiency. According to Matsumoto et al., cerates show

high proton conductivity, whereas zirconates present higher chemical and mechanical

stability [83].

10.5.1.1 Defect Thermodynamics

Defect chemical concepts and equations can be applied to these high-temperature proton

conductors (HTPC) [84, 85]. Protons are not part of the intrinsic oxide structure, but are

present as extrinsic defects (foreign species).

Let us discuss the example of SrCeO3, which is a p-type semiconductor in air at high

temperature, due to oxygen excess. In the presence of water vapor or hydrogen, the p-type

conductivity decreases and proton conductivity appears, because protons are incorporated

into the crystal lattice and electron holes are consumed [86–89], according toEquation 10.23

in Kr€oger–Vink notation [14]:

H2OðgÞþ 2h. > 2 H

.

i þ 1=2 O2ðgÞ ð10:23Þ

K ¼ ½H .

i �2P ðO2Þ1=2P ðH2OÞ�1½h.��2 ð10:24ÞThe formed proton can be considered as an interstitial, in Kr€oger nomenclature H

.

i .

However, the proton is an infinitely small point charge, strongly polarizing its environment,


and is inevitably attracted to the electron cloud of an oxide ion. Since an interstitial proton

interacts and associates stronglywith a neighbor oxide ion, it can be regarded as a hydroxide

ion on an oxide site, where it canmigrate by hopping between nearest-neighbor oxide ions (a

Grotthuss-type mechanism). The activation energy for proton jumps increases with

increasing oxygen–oxygen distance and is somewhat lower, as a rule of thumb, than the

activation energy for oxide ion transport.

The proton incorporation in oxygen-deficient oxides can then also be written:

H2OðgÞþVO.. þOO > 2 OH

.

o ð10:25Þ

K ¼ ½OH.

o�2½VO..��1

PðH2OÞ�1 ð10:26Þ

Protons are formed by water dissociation: the hydroxide ion fills an oxygen vacancy and the

proton forms a covalent bond with an oxide ion. The effective positive charge of the proton

defect (OH.

o ) is that of hydroxide versus oxide ion. The concentration of proton defects

competes with that of oxygen ion vacancies (V..

o ). Equation 10.25 shows immediately the

favorable conditions for proton formation: high water partial pressure and high density of

oxide ion vacancies, formed by reduction (Equation 10.12) or by acceptor doping. The

presence of acceptor dopants is more favorable for proton conduction, as the formation of

oxide ion vacancies by reduction leads to mixed or n-type conductivity (see Section 10.1).

Therefore, partial substitution of cerium by trivalent rare-earth cations, such as yttrium,

ytterbium, neodymium, and scandium, a typical example being SrCe0.95Yb0.05O3�d (see

also Chapter 4), significantly enhances the proton conductivity by providing extrinsic oxide

ion vacancies to accommodate the hydroxide defects. The typical proton concentration ½H.�is of the order of the acceptor doping level, about 1mol%. The proton mobility mH can be

determined from the equation:

sH ¼ FmH½H.� ð10:27Þwhere F is Faraday’s constant. The proton mobility is in the range of 5 10�6–5 10�5 cm2/

(sV) between 600 �C and 1000 �C with an activation energy around 0.5 eV, so that a proton

conductivity in the range of 10�2 S/cm is observed in SrCe0.95Yb0.05O3�d at 900�C [86].

The proton insertion reaction can obviously be written in terms of partial pressure of

hydrogen gas, H2(g), because there is a relation by the chemical equilibrium betweenwater,

hydrogen, and oxygen gas activities. Proton insertion can be also observed by contact with

hydrogen gas, even without oxygen vacancies, where mobile protons can be incorporated

according to the following reaction:

1=2 H2ðgÞ > H.

i þ e0 ð10:28ÞThe equilibrium constant for this reaction is:

K ¼ ½H.

i �½e0�PðH2Þ�1=2 ð10:29ÞHowever, water vapor is more favorable for observation of proton conductivity than

hydrogen gas, because it avoids a too-reducing atmosphere, which results in the predomi-

nance of n-type electronic conduction.


The dissociative adsorption of water is considered to be the main reaction leading to the

formation of protonic defects. In accordance with the point defect mechanism, the

concentration of protons is thus a function of partial pressures of H2O (or H2 and O2),

doping, and temperature. The proton and vacancy concentrations can be calculated using

thermodynamic data for proton-conducting oxides, which can be found for example in

Reference [84]. Inversely, thermodynamic data can be obtained if the proton concentrations

are experimentally determined. The entropy of reaction (10.25) has a high negativevalue (in

the order of �100 J/(Kmol)) in accordance with the loss of gaseous water. The reaction

enthalpy is also large and negative (typically�150 kJ/mol for BaCeO3 [4]). This reflects the

proton bond strength, typically related to the basicity, and the stability of the oxide. Highly

negative enthalpy values are observed for rare-earth oxides, such asY2O3, giving observable

proton concentrations at high temperature [24]. On the contrary, BaTiO3 and SrTiO3 that

showparticularly high values of protonmobility, up to 10�3 cm2/(sV) at 1000 �C [82, 90], do

not have a sufficiently high reaction enthalpy to present a high proton concentration and

show weak proton conductivity.

According to Le Chatelier’s principle, proton conduction dominates at low tempera-

tures, because reaction (10.25) is exothermic, while at elevated temperatures the reverse

reaction is favored: water desorption occurs and p-type electronic or oxide–ion conduc-

tivity is observed. The temperature at which dehydration starts depends on the oxide

composition. Given the increase of proton mobility with increasing temperature, a

maximum of proton conductivity can be expected in an intermediate range of temperature.

In conclusion, an optimized proton-conducting oxide requires appreciable acceptor

doping, intermediate temperatures, high water, and low oxygen partial pressures in the

gas phase.

Recent investigations of doped proton-conducting cerates, SrCe0.95Yb0.05O3�d (see also

Chapter 4) and BaCe0.95Yb0.05O3�d [91], and ferrates SrTi0.99Fe0.01O3�d showed that the

hydration proceeds not by chemical diffusion of H2O, but generally by decoupled, chemical

diffusion ofH andO (or ambipolar diffusion ofO2� and holes). The samemechanismworks

for zirconates as well, such as SrZr0.9Y0.1O3�d [92] (see also Chapter 4). The hydration and

dehydration kinetics of the zirconates proceeds by fast H-diffusion and slow O-diffusion,

consistent with the conductivity relaxation upon changing temperature in an atmosphere of

fixed water vapor activity and with the very slow relaxation kinetics of proton-conducting

zirconates. It is possible that someproperties,mistakenly attributed toequilibrium, including

proton conductivities, are in fact related to extremely slow transients.

As stated before, a large concentration of oxide ion vacancies is instrumental for the

formation of a large number of proton defects: it can be created by sufficiently low oxygen

partial pressure or by dopingwith lower valent acceptor cations, such as rare-earth elements,

where vacancies are formed for charge compensation. In that case, the concentration of

oxide ion vacancies is temperature independent. Accordingly, the proton concentration in

donor-doped samples is insignificant. The amount of proton defects OHO� can be deter-

mined using the electroneutrality condition:

½h.� þ 2½V..

o � þ ½OH.

o� ¼ ½M0B� þ ½e0�

whereMB0 are lower valent acceptor cations. Boundary segregation of acceptors, for instance

by a space charge mechanism [93], can lead to a reduction of the amount of compensating


oxide ion vacancies and thus the amount of protons formed by water uptake. Furthermore,

some defect ordering might appear for very large vacancy concentrations and part of the

oxide ion vacancies might then be unavailable, but this is still a matter of discussion [94].

The crystallographic structure of a protonic defect in a cubic perovskite oxide is shown in

Figure 10.6. There are eight orientations of the hydroxide ion stabilized by hydrogen bonds

with the eight nearest oxide ions in the perovskite structure. As opposed to some cases

discussed previously, for which hydrogen bonding is the dominant intermolecular interac-

tion, here hydrogen bonding is restricted to the defect region. Within this region, hydrogen

bonding interferes with other chemical interactions and, together with these, determines the

structure and dynamics of the defective region.

Other examples of perovskite-type proton conductors are KNbO3 and KTaO3, which

present slightly distorted cubic ReO3 structures and activation energies around 1 eV [95].

Their upper limit of proton conductivity is estimated to be about 10�6 S/cm.

Atomistic calculations have been done to simulate proton diffusion in numerous

perovskite-type oxides [96–98]. The energy required to break O–H bonds is about 4 eV,

so proton transfer occurs through a mechanism in which there is no significant bond

breaking. The low energy barrier for the hopping process is a result of strong interactions of

the jumping proton with both neighboring oxide ions. Simulation shows also that protons

form short and transient hydrogen bonds with all eight neighbor oxide ions. More on the

fundamentals of proton transport in perovskite oxides can be found inReferences [9, 99] and

in Chapter 7 of this book.

10.5.1.2 Chemical and Thermal Stability

Themain challenge related to high-temperature proton conductor development is to achieve

high proton conductivity while preserving thermochemical stability. Oxides combining

Figure 10.6 Proton defect inside a cubic perovskite structure. The possible proton orientationswith eight nearest-neighbor oxide ions are indicated [9]. Reprinted with permission fromChemical Reviews, Transport in Proton Conductors for Fuel-Cell Applications: Simulations,Elementary Reactions, and Phenomenology by K. Kreuer, S. Paddison, E. Spohr, M. Schuster,104, 10, 4637–4678 Copyright (2004) American Chemical Society.


high proton conductivity and high thermodynamic stability are a realistic alternative for

different types of electrochemical cells. Among the proton-conducting perovskite oxides,

BaCeO3 presents high proton conductivity [100], but suffers from poor chemical stability,

because it reacts with acidic gases, such as CO2 and SO2, and water vapor, decomposing to

CeO2 and Ba(OH)2. The decomposition is particularly deleterious for use in fuel cells or

steam electrolyzers [101] and with thin films of BaCe0.8Y0.2O3 [102]. SrCeO3 has higher

stability, due to higher basicity of the oxide (see above), but can decompose under high

partial pressures of CO2 to SrCO3 and CeO2, which is a problem for some applications [83].

Y-doped BaZrO3 also shows better chemical stability, but must be sintered at high

temperature and the proton conductivity of ceramic samples is low, due to the presence

of poorly conducting grain boundaries, and only thin films seem promising for applica-

tions [103]. As grains grow to awell-crystallized structurewith refined grain boundaries and

lower yttrium segregation at grain boundaries, proton conductivity increased at 500 �C, forsamples annealed at 800 �C, 1250 �C, and 1500 �C, from 8.7 10�5 to 2 10�3 and 4 10�3 S/cm

with grain sizes of about 10 nm, 50 nm, and 200 nm, respectively [104]. A nanometric grain

size is clearly not a panacea for increasing proton conductivity, as demonstrated previously

for solid oxide ion conductors [12]. The relatively important grain boundary resistance in

ceramics is probably related to negatively charged space charge regions and the depletion of

positively charged proton defects [105], also evidenced in SrZr0.9Y0.1O3 [106].

Generally, highly basic oxides, such as barium cerate, have better stabilized protonic

defects, anddehydrationoccurs at temperatures above600 �C,whereas lessbasicoxides, suchas barium zirconate, start already dehydrating above 400 �C. The most suitable temperature

range for proton conduction in oxides is thus a compromise between sample hydration and

proton mobility. Generally, proton conduction maxima are observed around 400–600 �C.Similar to hydrogen bond networks, a symmetry reduction of the crystal lattice reduces

the proton conductivity and enhances the activation energy. This effect can be observed by

comparison of yttrium-doped BaCeO3, SrCeO3, and SrZrO3. The orthorhombic distortion

of SrCeO3 has big effects on the arrangement of the oxide ions. The cubic oxide site

degenerates into two sites (O1, O2) with different acid–base properties (see also Figure 7.31

in Chapter 7). The time-averaged structure of a protonic defect in cubic perovskite-type

oxides shows eight orientations of the central hydroxide ion stabilized by hydrogen bond

interactions with the eight nearest oxide neighbors (Figure 10.6). Assuming that protons are

associated with these sites for the majority of the time, in SrCeO3, long-range proton

transport involves transfers between chemically different O1 and O2 sites. This was

suggested as the reason for the higher activation enthalpy and lower proton conductivity

of SrCeO3 in comparison with cubic BaCeO3 [9]. The mobility of protonic defects is very

sensitive not only to reduction of crystallographic symmetry, but also to local structural and

chemical perturbations induced by acceptor dopants or by mixed occupancy on the B site.

Traditionally, lower valent dopants with ionic radii matching those of the host ion are

chosen, but for proton conductivity in oxides, this approach clearly fails. Although Sc3þ

and In3þ have similar ionic radii as Zr4þ , BaZrO3 doped with scandium or indium shows

much lower proton mobility than BaZrO3 doped with yttrium, which has a significantly

higher ionic radius. Commonly, increasing the dopant concentration reduces proton

mobility and increases the activation energy, but for Y-doped BaZrO3, the proton mobility

and activation enthalpy are nearly independent of the dopant concentration. Electronic

structure calculations showed a significant effect of the acceptor dopant on the electron


density of the neighboring oxide ions, including their affinity for the proton. These

considerations qualitatively explain the experimental result that the highest proton con-

ductivities are observed in cubic oxides with a perovskite structure.

Quaternaryperovskite-related structureswith thegeneral formulaA2BB0O6andA3BB

0O9

and an excess of lower valent cation B0, compensated by oxygen vacancies or protons and

showing good proton conductivity, were reported byNowick and co-workers [85, 107, 108].

In particular, systems such as A3Ca1�xNb2�xO9 or Sr2(ScNb)O6 are extremely sensitive to

cation content x, with major effects on structural order, proton mobility, and mechanical

properties. The compound Ba(Ca1.18Nb1.82)O9, so-called BCN18 (see also Chapter 4), can

take up 0.18 protons per ABO3 unit and exhibits (according to Reference [85]) one of the

highest proton conductivities of these systems, about 10�2 S/cm at 200 �C. However, aproblem is again the relatively important grain boundary resistance and significantly lower

conductivities have been found in polycrystalline samples. Negatively charged, proton-

blocking space charge regions are now well established [109, 110]. Higher proton contents

maybe achieved in perovskiteswith higher oxygen deficiency and structurally ordered oxide

ion vacancies. Evidence was obtained that the Brownmillerite compound Ba2In2O5 under-

goes a reversible phase transformation upon exposure to humid atmosphere at 250 �C [111],

leading to a new proton phase Ba2In2O5�H2O that can be described as a distorted double-

perovskite structure [112]. The sheet consisting of parallel chains of InO4 tetrahedra and

parallel rows of oxygenvacancies in the parent structure ofBa2In2O5has been converted into

an InO6 octahedral perovskite-like sheet [112]. Ba2YSnO5.5 saturated with water corre-

sponds to an oxide–hydroxide Ba2YSnO5OH [86, 113].

The oxygen-deficient Ba2In2(1�x)Ti2xO5þ x(0� x� 1) compounds also react

with water vapor, and their preparation in air always leads to samples containing

protons. The conductivity of Ba2(In1�xTix)2O5þ x compounds is mainly protonic up to

450 �C, and the best proton conductivity was obtained for x¼ 0.2 with 1.1�10�3 S cm�1 at 450 �C [114]. The beginning of the dehydration process of the

Ba2In2(1�x)Ti2xO4þ 2x(OH)y [0� x� 1; y� 2(1� x)] compounds is observed around

200 �C. The highest conductivity, s180� 10�6 S cm�1, is observed for x 0.3 [112].

10.5.2 Other Structure Types

Overall, the perovskite systems continue to dominate. Other structure types [115], such as

spinels, have been studied for proton conduction, but with negative results, maybe due to a

too rigid oxide ion sublattice. Amore dynamic oxygen sublattice is useful, butwithout long-

range oxygen transport. The fluorite-related La6WO12 exhibits considerable proton con-

ductivity [116]. La1�xW1/6O2 (x¼ 0.05, 0.1) compounds are mixed ionic–electronic

conductors. Above 800 �C, p- and n-type electronic conductivity predominates under

sufficiently oxidizing and reducing conditions, respectively. Ionic conduction is dominating

at intermediate oxygen partial pressures at high temperatures and at low temperature

(<800 �C) across the range of oxygen partial pressures investigated. Protons are the major

ionic charge carrier underwet conditions below800 �C. Proton conductivity peaks at 750 �Cwith a value of 3� 10�3 S/cm in wet H2 [117].

There appears to be a growing trend toward structure types containing tetrahedral anions,

which are immobile and do not show oxide ion conductivity. Initially, there was significant

interest in lanthanide phosphates. Sr-doped LaPO4 shows respectable proton conductivity


(3 10�4 S cm�1 at 900 �C) and some p-type contribution without parasitic oxide ion

conductivity [109]. An unusual feature of these systems is that the conductivity enhance-

ment is observed for very low doping levels (�1 at %). Related structures with tetrahedral

MO4 anions, such as alkaline earth-doped LaMO4 (M¼Nb, Ta), are currently investigated.

The conductivities of LaMO4 systems are lower than in perovskites (10�3 S cm�1 at 800 �Cfor Ca-doped LaNbO4 [118]), but they show higher stability in CO2-containing atmo-

spheres [119]. For phosphate systems, NMR suggested that oxide ion vacancies are

accommodated by the formation of P2O74� units, which produce HPO4

2� units when

hydrated. Similarly, recent modeling work on LaNbO4 suggested the formation of

Nb3O117� defect clusters. Similar proton conductivities have also been observed in gallate

systems, La1�xBa1þ xGaO4�x/2 and La1�x(Sr/Ba)2þ xGaO5�x/2 (0� x� 0.2), containing

tetrahedral GaO4, although they suffer from stability problems. In this system, oxide ion

vacancies are accommodated by the formation of Ga2O78� units, which are broken up by

water incorporation.

Doped tin pyrophosphates, such as In0.1Sn0.9P2O7�d, were investigated as a new

promising class of proton-conducting oxides, but initially reported high proton conductivi-

ties might be due to remaining phosphoric acid [120]. Ga-doped pyrophosphates

Sn1�xGaxP2O7�d showed a conductivity of 5 10�2 S cm�1 at doping limits of x¼ 0.09 in

wet H2 at 175�C [121].

Pyrochlore-type oxides are well-known solid ionic conductors with the general formula

A2B2O6X, where A is an alkali ion, such as Na or K; B is a transition element, such as Ta or

W; and X¼O, F, OH. The ideal loosely packed structure relies on corner-shared BO6

octahedra with planes of A2X units, where X is situated in tetrahedral sites between four A

at hexagonal sites. Given this loose packing, especially defect pyrochlores of formula

AB2X6, are expected to be excellent oxide ion conductors. Proton conduction was first

reported in 1996 in Ln2Zr2�xYxO7�d (Ln¼La, Nd, Gd, and Sm) with conductivity around

10�3 S/cm at 800 �C [122] and also in other phases of this family, such as HTaWO6.

H2O [123].

10.6 Conclusion

The types of classification discussed in this chapter are obviously related. Low-temperature

inorganic solid proton conductors are generally based on the existence of water in the

structure, playing the role of vehicle (vehicular mechanism) or being essential for

cooperative translocation (Grotthuss mechanism). Intermediate-temperature solid proton

conductors show often phase transitions, allowing a large mobility of protons only in the

high-temperature “superionic” phase. In high-temperature solid proton conductors, protons

are extrinsic defects of crystalline, especially perovskite-type, oxides.

As already pointed out by Alberti and Casciola in 2001 [7], there are few really new

inorganic proton conductors reported in the last decades, andmorework is beingdedicated to

improving existing materials than finding new ones. Certainly, the competition with

polymeric proton conductors is difficult at low and intermediate temperatures. However,

a significant field for ulterior development can be found in the high-temperature layered

structures and oxides, especially proton-conducting perovskites,with the large development


of studies in recent years. This domain is now clearly leading inside inorganic solid proton

conductors. However, several important issues concerning materials stability must still be

overcome before inorganic solid proton conductors can be used on a large scale in practice.

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