electrophysical properties of titanates of alkaline-earth metals

ISSN 1023-1935, Russian Journal of Electrochemistry, 2007, Vol. 43, No. 8, pp. 894–900. © Pleiades Publishing, Ltd. 2007.Original Russian Text © L. A. Dunyushkina, 2007, published in Elektrokhimiya, 2007, Vol. 43, No. 8, pp. 940–948.

894

INTRODUCTION

Compounds on the basis of titanates of alkaline-earth metals (AEM) are of interest as being of promisefor assorted electrochemical devices. For example,iron-doped calcium titanate CaTi

1

–

x

Fe

x

O

3

–

δ

is stableenough in reducing atmospheres and values of its ionicand electron conductance are commensurate with oneanother. It can be used as membrane in the hydrogenproduction by means of electrochemical conversion ofnatural or technogenic combustible gas [1]. Employingoxygen-penetrable membranes makes it feasible to pro-duce hydrogen by an economic and ecological safetechnique.

Composite materials containing SrTiO

3

doped withlanthanum are viewed as an alternative to the traditionalNi/YSZ anode in solid oxide fuel cells. The point is thatthese materials are more stable in oxidizing atmo-spheres and withstand contamination by sulfur (sulfur-containing impurities), which are as a rule present innatural gas [2].

Some perovskite-structured oxide materials dopedwith acceptor admixtures exhibit high proton conduc-tion [3, 4]. Certain AEM titanates possess unipolar pro-ton conduction. Such materials are of interest as solidelectrolytes in electrolyzers and fuel cells and in sen-sors for protonated media.

Apart from having practical significance, the AEMtitanates are convenient model systems. Their perovs-kite-like structure is tolerant toward substitutions byacceptor and donor admixtures in both cationic sublat-tices, which provides for an opportunity to find out howtheir composition affects properties of materials.

This paper is a review of the results obtained duringinvestigation of the ionic, electron, and proton conduc-tion as well as oxygen penetrability of some AEM titan-ates. The review focuses on the studies carried out atInstitute of High-Temperature Electrochemistry, UralBranch, Russian Academy of Sciences (IHTE).

IONIC AND ELECTRON CONDUCTIONOF ATi

1

–

x

M

x

O

3

–

δ

(A = Ca, Sr, Ba)

The electric properties of perovskite-structuredoxide materials are to a large extent determined by thedefects in the oxygen sublattice, which depends on thetype and concentration of admixtures in either cationicsublattice, as well as on the temperature and partialoxygen pressure ( ) in the surrounding atmosphere.The way the electroconduction of single crystals of per-ovskite CaTiO

3

depends on the partial oxygen pressurewas probed into in [5]. According to the defect forma-tion model that was put forth in [5], the charge is carried

by twice-ionized oxygen vacancies and electrons.

The effect various admixtures exert on the electricproperties of CaTi

1

–

x

M

x

O

3

–

δ

was examined by theauthors of [6–16]. With cobalt and nickel substitutedfor titanium, the region of solid solutions does not exceed

x

= 0.1, and in the case of iron it reaches

ı

= 0.5 [6]. Withtitanium replaced by aluminum, single-phase solidsolutions based on CaTiO

3

form throughout the entireinterval studied (

x

= 0–0.4) [7]. Upon replacement bychromium and indium, such solutions form at, respec-tively,

ı

≤

0.2

and

ı

< 0.05 [8].

Substituting iron for titanium results in a maximumrise of the oxygen-ion and electron conduction.According to the model suggested in [6], substituting

pO2

Vo

..

Electrophysical Properties of Titanates of Alkaline-Earth Metals

L. A. Dunyushkina

z

Institute of High-Temperature Electrochemistry, Ural Branch, Russian Academy of Sciences, ul. S. Kovalevskoi 22, Yekaterinburg, 620219 Russia

Received December 20, 2006

Abstract

—Results on oxygen-ion, electron, and proton conduction and oxygen penetrability of titanatesof alkaline-earth metals doped with acceptor admixtures are briefly reviewed. The applicability of these mate-rials in electrochemical devices, in particular, as oxygen-penetrable membranes, is considered. The focus ison the studies carried out at the Institute of High-Temperature Electrochemistry, Ural Branch, Russian Acad-emy of Sciences.

Key words

: titanates of alkaline-earth metals, ionic conduction, electron conduction, proton conduction, oxygenpenetrability

DOI:

10.1134/S1023193507080071

z

Author’s e-mail: [email protected]

RUSSIAN JOURNAL OF ELECTROCHEMISTRY

Vol. 43

No. 8

2007

ELECTROPHYSICAL PROPERTIES OF TITANATES OF ALKALINE-EARTH METALS 895

ions Fe

3+

for ions Ti

4+

leads to the formation of oxygenvacancies:

CaO

+ (1 –

x

)

TiO

2

+ (

x

/2)

Fe

2

O

3

=

CaTi

1

–

x

Fe

x

O

3

–

0.5

x

.(1)

The concentration of vacancies is defined by the con-tent of ions Fe

3+

. However, it is noteworthy that theexact balance between these is defined by the generalelectroneutrality condition (using the Krueger–Vincenotation):

(2)

Here,

and

n

stand for the concentration of holes andelectrons, respectively.

The formation of vacancies in the oxygen sublatticeis a reason for the appearance of high ionic conduction

σ

i

. On the other hand, in oxidizing conditions, vacan-cies interact with oxygen of a gas phase, leading to theappearance of hole conduction:

(3)

In reducing conditions, oxide sheds oxygen with theappearance of additional vacancies and electrons e

'

:

(4)

If one ignores the concentration of holes and electronsin the electroneutrality condition (2) as compared withconcentrations of vacancies and iron, then concentra-tions of holes (5) and electrons (6) will depend on thepartial oxygen pressure to powers of 1/4 and –1/4,respectively:

p

= (1/2

ä

3

)

1/2

(5)

(6)

Here,

ä

3

and

ä

4

are the equilibrium constants for reac-tions (3) and (4).

The effect the iron concentration has on the electricproperties of CaTi

1

–

x

Fe

x

O

3

–

δ

was the subject matter of

2 Vo

..p+ FeTi'[ ] n.+=

Vo

..1/2O2+ Oo

x 2h..+=

Oox 1/2O2 Vo

..2e'.+ +=

FeTi'[ ]1/2pO2

1/4,

n 2K4( )1/2 FeTi'[ ]–1/2pO2

1/4.=

[1, 9–13]. The electroconductivity isotherms as func-tions of the partial oxygen pressure are characteristi-cally cup-shaped [1, 9–12]. In oxidizing atmospheres,the total electroconductivity

σ

tot

lowers with diminish-

ing almost proportionally to

σ

tot

~

; in an inter-

mediate region it is independent of ; and in reducingatmospheres it rises with decaying partial oxygen pres-

sure as

(Fig. 1). Dependences of that kind corre-spond to the above defect formation model and the totalelectroconductivity satisfies the equation [1]

(7)

The contribution made by the electron conduction(both

n

and

p

types) to the total conduction in the regionof the plateau in the

–

isotherm may beignored and the total conduction may be assumed to beionic. Data of various authors show the dependence ofthe ionic conduction on the iron concentration to havea maximum (

σ

i

= 0.09 S cm–1) at x = 0.2 and 1000°C[11]. The concentration dependences of the ionic con-ductivity for CaTi1 – xFexO3 – δ, SrTi1 – xFexO3 – δ, andBaTi1 – xFexO3 – δ appear in Fig. 2.

There is no doubt whatsoever that such a characterof the dependence of the ionic conductivity on the con-tent of the alloying admixture must correlate with vari-ations in crystalline structure.

Using the method of diffraction of x-rays and elec-tron microscopy, the authors of [17] showed that oxy-gen vacancies in solid solutions Ca2Ti2 – 2xFe2xO6 – x forcompositions close to CaTiO3 (0 < ı ≤ 0.4) seem statis-tically distributed with pseudocubic symmetryretained, whereas in the region of compositions 0.55 ≤ı < 1 defects undergo ordering in the (0k0) plane. Themethod of transmission electron microscopy allowedthe authors of [18] to observe a developed domainmicrotexture with a characteristic size of grains equalto 5–10 nm for the composition x = 0.38. The research-

pO2pO2

1/4

pO2

pO2

–1/4

σtot σi σhopO2

1/4 αnopO2

–1/4( ).+ +=

σtotlog pO2log

–logσion [S cm–1]

3

16 12 8 4–log pO2

[atm]

2

1

0

x = 00.10.20.30.5

CaTi1 – xFexO3 – δ, 1000°C [11]

Fig. 1. Dependences of conductivity in CaTi1 − xFexO3 – δon é2

at 1000°C [11].

σi, S cm–1

0.1

0.1 0.2 0.4 0.6x

0

1000°C

0.2

0.3

0.4SrTi1 – xFexO3 – δ, [17]

0.3 0.5 0.7

BaTi1 – xFexO3 – δ, [17]

CaTi1 – xFexO3 – δ, [17]

SrTi1 – xFexO3 – δ, [19]

Fig. 2. Dependences of ionic conductivity inÄTi1 − xFexO3 – δ (A = Ca, Sr, Ba) on the iron content at1000°C [17, 19].

896

RUSSIAN JOURNAL OF ELECTROCHEMISTRY Vol. 43 No. 8 2007

DUNYUSHKINA

ers connected its formation with the ordering of oxygenvacancies [18]. However, on the basis of these data it isdifficult to explain why the composition x = 0.2 pos-sesses maximum ionic conduction.

To expose correlation between electric propertiesand evolution of crystalline structure and oxygen nonsto-ichiometry the crystalline structure of CaTi1 – xFexO3 – δwas refined in [14] with the aid of the method of full-profile analysis (FPA) on the basis of am orthorhombicstructure. On the basis of the FPA data there was pro-posed a microscopic mechanism for the formation,ordering, and transport of oxygen vacancies. Accordingto FPA, oxygen octahedrons, which are centered by atitanium (iron) atom, are compressed along the c axis.Oxygen ions occupy two types of nonequivalent posi-tions in the CaTi1 – xFexO3 – δ structure, specifically, inthe (001) plane, which contains titanium (iron) ions,and in the (001) plane, which contains calcium ions.With ions Fe3+ substituted for ions Ti4+, oxygen vacan-cies form in the (001) planes that contain Ti(Fe), whereasthe degree of occupation of oxygen positions in the (001)planes that contain Ca is close to unity. In this case onemay presume oxygen ions in the (001) planes that con-tain calcium to be rigidly bound and the oxygen transportto be realized via the Ti(Fe)-containing planes. A cross-section of the surface of CaTi1 – xFexO3 – δ by the (001)plane, which contains titanium (iron) ions, is depictedschematically in Fig. 3. One may see that every oxygensquare borders with vertexes of four analogous squares.At low concentrations, ions of iron and vacancies ofoxygen must statistically be distributed in the perovs-kite lattice according to the requirement of the entropymaximum. At x = 0.2, iron replaces every fifth titaniumatom. Consequently, every complex comprising fivesquares that are presented in Fig. 3 contains one ironatom and, after a further increase in x, substitutionaldefects form in the neighboring squares. It is logical toassume that an oxygen vacancy that is a neighbor oftwo ions Fe3+ is bound more rigidly and its mobilityreduces, which reflects in the lowering of oxygen-ionconduction at x in excess of 0.2. Following the forma-tion of oxygen vacancies, the coordination number ofiron ions must alter and this was confirmed by results ofMoessbauer spectroscopy that were presented by theauthors of [13, 18]: iron in the CaTi1 – xFexO3 – δ latticemay occupy octahedrally, pentahedrally, and tetrahe-drally coordinated positions, and in so doing ions Fe3+

with the pentahedral coordination by oxygen ions were

discovered solely at ı ≤ 0.20, and Fe3+ with the tetrahe-dral coordination, for ı ≥ 0.20 [13].

Thus, according to [14], the fact that the ionic con-duction in CaTi1 – xFexO3 – δ increases with the iron con-tent at small values of ı (ı < 0.2) is due to an increasein the concentration of oxygen vacancies, and thedecrease in the ionic conduction at ı > 0.2 is due to theformation of complexes of defects.

Such a dependence of the ionic conductivity on theadmixture content was obtained for CaTi1 − xFexO3 – δ[7]: the maximum is observed in the vicinity of ı = 0.2.In compounds SrTi1 – xFexO3 – δ and BaTi1 − xFexO3 − δthere is observed a tendency toward an increase in theionic conductivity with increasing iron content at x =0−0.5, but against the backdrop of this tendency there isintermediate maximum at x = 0.2 in either system [15].The authors of [16, 19] noted only the increase in theionic conductivity in SrTi1 – xFexO3 – δ with increasingiron content, but the step of variation in x near x = 0.2was rather large (x = 0.15, 0.40 [18] and x = 0.1, 0.4[19]). As one can infer from Fig. 4, results [15] do notcontradict data [19] (in [15] presented are data forlower temperatures).

The region of solubility of dopants in systemsCaTi1 – xCrxO3 – δ, CaTi1 – xInxO3 – δ, SrTi1 – xMnxO3 – δ,and SrTi1 – xAlxO3 – δ is smaller than 0.2 and in the limitsof the region of single-phasedness there is observed anincrease in the ionic conductivity with increasingadmixture content [7, 8].

We deem it possible to presume that the mechanismof correlation of the concentration of an acceptoradmixture and the ionic conductivity proposed in [14]is common for perovskites on the basis of the AEMtitanates. However, this model is applicable for the casewhere the crystalline structure of a perovskite possessesorthorhombic distortions. This leads to anisotropy oftransport properties and allows us to presume the exist-ence of crystallographic planes that are preferable forthe transport of oxygen ions. It is known, however, thatSrTi1 − xFexO3 – δ has a cubic symmetry [14, 15, 19].Consequently, its structure is isotropic. The structure ofÇ‡Ti1 − xFexO3 – δ is hexagonal [15]. It is possible that,with decreasing the partial oxygen pressure in an atmo-sphere and the growth of oxygen nonstoichiometryconnected with this, structure undergoes distortion. Tocheck the universal nature of model [16] it is necessaryto perform simultaneous in situ investigations of elec-troconduction and crystalline structure of perovskites atdifferent values of the partial oxygen pressure.

OXYGEN PENETRABILITYOF THE AEM TITANATES

Owing to the presence of oxygen-ion and electronconduction, materials on the basis of the AEM titanatesare electrochemically permeable by oxygen in condi-tions of a gradient of the partial oxygen pressure. TheAEM titanates possess the required stability in a broad

Ti (Fe)

Fig. 3. Cross-section of the CaTi1 − xFexO3 – δ structure bythe (001) plane passing through ions of Ti(Fe) [16].



range of partial oxygen pressures. For example, wheninvestigating electroconduction in ë‡Ti1 − xFexO3 – δ,specimens were kept for tens of hours at of 10–16 to

10–20 atm at temperatures of 900 to 1000°C. The speci-mens did not destroy under those conditions and theirproperties were reproducible when cycling the partialoxygen pressure [11]. The production of dense andhomogeneous ceramics out of titanates is simpleenough. In connection with this, very appealing looksthe use of these materials in the role of ceramic electro-chemical membranes for the production of oxygen outof oxygen-containing gases or for the production ofhydrogen by means of electrochemical conversion of acombustible gas. As was demonstrated in [6], oxygenmay be produced out of air with the aid of mixed ion–electron conductor CaTi0.8Fe0.2O3 – δ. The efficiency ofsuch a membrane is substantially higher than that of atraditional siliconic film.

The oxygen flux through a gas-tight ceramic mem-brane is defined by the relationship [20]:

(8)

Here, σamb is the ambipolar conductance, ê0 and êL arevalues of the partial oxygen pressure on the oppositesurfaces of the membrane, and L is the membranethickness. The ambipolar conductance is defined byboth ionic and electron conductances of the material:

(9)

Equation (8) is valid in the case of a fast oxygenexchange at the membrane interfaces with the gas phasewhen the electrochemical influx of oxygen through themembrane is restricted by bulk diffusion. The oxygenflux through the membrane is inversely proportional tothe membrane’s thickness. From this viewpoint themembrane must be as thin as possible. However, hin-drances of surface exchange exist as well. Therefore, ifthe thickness of the mixed conductor turns smaller thana certain critical value, the oxygen transport from thecathodic space into the anodic one becomes limited bythe surface exchange stage an a further decrease in themembrane thickness makes now no difference. Thecritical thickness depends on the chemical composi-tion, temperature, state of the membrane surface, gasphase, because the latter affects the membrane’s prop-erties. Consequently, calculating the critical value ofthe membrane thickness is a rather daunting task butstill may be realized by experimentally investigatingthe oxygen penetrability of mixed conductors in partic-ular operating conditions.

In oxidizing conditions and at a small gradient of thepartial oxygen pressure (0.01 atm < < 0.209 atm onthe anodic surface of the membrane and air on thecathodic surface), the oxygen penetrability of

pO2

J O2( ) –RT

16F2L---------------- σamb pO2

.lnd

P0

PL

∫=

σamb

σiσe

σi σe+----------------.=

pO2

CaTi0.8Fe0.2O3 – δ is well described under the assumptionthat the ambipolar diffusion over bulk is the rate-deter-mining stage [21]. At a low partial oxygen pressure inthe anodic space, however, experimental values of theoxygen flux through a membrane turn smaller thanthose computed under such an assumption. This dis-crepancy is probably connected with a slow oxygenexchange at the anodic membrane/gas interface. Theimplication is that conditions at this interface are suchthat the rate of the oxygen influx through the membraneis limited by the surface exchange stage.

Probing into the oxygen penetrability ofSrTi1 − xFexO3 – δ showed the oxygen flux through amembrane to depend on both the ambipolar diffusion ofoxygen over bulk and the surface of exchange rate[22−24]. In air at temperatures of 700–1100°C, the holeconductance in SrTi1 − xFexO3 – δ is ten and more timesthe ionic conduction [19]. As a result, with allowancemade for equation (9), the ambipolar conductance inSrTi1 − xFexO3 – δ is close to the ionic conductance. Val-ues of the ambipolar conductance, calculated forSrTi1 − xFexO3 – δ from experimental results on the oxy-gen penetrability, are substantially smaller than those ofthe ambipolar conductance in the same temperatureinterval, which testifies to a contribution made by theslow surface exchange [22]. Using a two-layered mem-brane comprising a porous layer and a dense one, madeit possible to considerably accelerate the electrochemi-cal flux of oxygen, owing to a more developed mem-brane/gas interface and an easier surface exchange [23].Applying catalyst Pt/PrOx onto the surface of a mem-brane also amplifies the oxygen flux through the mem-brane [24]. The oxygen penetrability of a membranedepends on the technique used for modifying its sur-face. This fact confirms that the surface exchange stagehas an impact on the kinetics of electrochemical influxof oxygen. Oxygen penetrability increases with decreas-ing membrane thickness. This fact indicates that the

I, A cm–2

4 3 2 0–log pO2

[atm]5

0

0.01

0.02

1000°C

1

1050°C

1000°C

Fig. 4. Dependences of the density of oxygen currentthrough a CaTi0.8Fe0.2O3 – δ membrane 1 mm thick on

at the anodic membrane side. Solid line represents a calcu-lation for 1000°ë under assumption that bulk diffusiondefines oxygen influx [21].

pO2

898


DUNYUSHKINA

ambipolar diffusion over the bulk of SrTi1 − xFexO3 – δalso defines the oxygen influx rate [24]. At lower tem-perature, the role played by the bulk diffusion increasesas compared with the surface exchange. At tempera-tures below or equal to 1073 K, it is precisely the bulktransport that restricts oxygen penetrability [24].

Dependence of the oxygen influx current on the gra-dient of the partial oxygen pressure has a tendencytoward saturation [23]. In the authors' opinion, this phe-nomenon may be due to a concentration overvoltageconnected with the slow oxygen transport at a low oxy-gen concentration in the gas phase.

Hydrogen may be obtained in an electrochemicalreactor on the basis of a membrane made ofCaTi0.8Fe0.2O3 – δ. A scheme of such a reactor was pro-posed in [25], and its parameters were analyzed. Pro-ducts of partial oxidation of methane (CO, CO2, H2,H2O) enter the electrochemical section of the reactor.Vapor that is obtained in an evaporator is directed intothe cathodic channel. To make the membrane’s condi-tions softer—to diminish the partial oxygen pressuredrop at the opposite surfaces—the vapor is mixed witha portion of the exiting cathodic gas that containshydrogen.

A major portion of the gas mixture that exits out ofthe anodic channel is directed into an afterburner. Theresidual fuel burns in the latter with the formation ofwater and carbon dioxide. A major portion of the exit-ing cathodic mixture is directed into a condenser, wherehydrogen is purified from water.

Owing to oxygen influx through the membrane, thecomposition of gases in the anodic and cathodic chan-nels varies along the channels. The concentration ofreduced components drops in the anodic space,whereas in the cathodic channel increases the concen-tration of hydrogen. Consequently, to provide for a pos-itive emf value along the entire electrochemical section,the only acceptable configuration is the counter flow ofthe anodic and cathodic gas mixtures. The stepsinvolved in the determination of profiles of the emf andcurrent density along the electrochemical section of thereactor were described in [25]. The device productivityas a function of the composition of anodic and cathodicgases was also analyzed in [25]. It was demonstratedthat the productivity increases with the content of car-bon monoxide and hydrogen in the input anodic gas andthe humidity of the output cathodic gas. Producinghydrogen out of hydrocarbons with the aid of an oxy-gen-permeable membrane may be viewed as a promis-ing technique for production of pure hydrogen. In thisaspect it is of importance to examine the hydrogen pen-etrability of potential membrane materials as well.

PROTON CONDUCTIONIN THE AEM TITANATES

The hydrogen defects arise in oxides during interac-tion with a hydrogen-containing atmosphere, say, with

hydrogen and/or water vapor. The oxide having oxygenvacancies, the dissolution of water vapor is describedby the known reaction

(10)

Here, is the proton localized on an oxygen ion.Hence, doping perovskites with acceptor admixturesthat leads to the formation of oxygen vacancies forcounterbalancing the charge is likely to amplify theproton conduction.

The mechanism of the proton conduction is debat-able. A widely held viewpoint is that the diffu-sion is due to the proton transport between oxygen ionsvia the Grotthus mechanism, rather than through themigration of the hydroxide ion as a whole [26]. How-ever, other hydrogen defects probably make their con-tributions to the proton conduction as well. These mayinclude interstitial hydroxide ions and hydride ions

and , as well as hydroxide ions and hydride

ions and occupying positions of oxygen ions[27]. The authors of [4] presumed simultaneous exist-ence of two types of charged hydrogen defects, namely,protons ç+ and hydroxide ions éç–.

To the minds of the authors of [28], the concentra-tion of hydrogen defects in compounds produced on thebasis of strontium and barium titanates is low. Coupledwith the high thermodynamic stability, this factorattracts interest to these compounds as to potentialmaterials for protonic electrochemical devices.

The proton–oxygen conduction in perovskitesATi0.95M0.05O3 – δ, where A = Ca, Sr, Ba and M = Mg,Sc, was examined in [4, 29] in various atmospheres,specifically, in a reducing hydrogen-containing atmo-sphere and in humid air. The transport numbers for ionsand protons were measured by the emf method with theaid of the oxygen and water-vapor concentration cells.The emf method is incapable of distinguishing betweenthe transport of protons, oxygen ions, and hydroxideions without evoking some additional assumptions. Inconditions of presumably simultaneous transport ofoxygen ions and protons ç+, the temperature depen-dences of the proton conductivity exhibit each a maxi-mum, whereas under the assumption about simulta-neous transport of hydroxide ions éç– and protons thesame dependences are almost linear [4]. However, a lin-ear character of the temperature dependence of the pro-ton conduction cannot prove the transport of éç– andç+. To do this, it is necessary to examine the diffusioncoefficients for hydrogen and oxygen as a function ofthe hydrogen concentration in the oxide.

According to [4], in a reducing atmosphere, thetransport numbers for protons in titanates of Ca, Sr, andBa substituted by Sc and Mg increase with lowering thetemperature. At a temperature of 500°ë and lower,CaTi0.95Sc0.05O3 – δ is a purely proton conductor. InSrTi0.95Sc0.05O3 – δ and SrTi0.95Mg0.05O3 – δ, the proton

H2O g( ) VO

..OO

x+ + 2OHO

..=

OHO

.

OHO

.

OHi' Hi'

OHO

.HO

.



conduction is predominant. Using such materials assolid electrolytes would permit a substantial lower-ing of the operational temperature of electrochemi-cal devices. The concomitant insufficiently high pro-ton conductivity (~10–4 S cm–1 at 500°ë inCaTi0.95Sc0.05O3 – δ [4]) could be counterbalanced withthe aid of thin-film and nano technologies.

The proton conduction in CaTi1 − xFexO3 – δ was thegoal of studies performed in [9, 30]. The transport num-bers for protons were found to decrease with increasingtemperature and the iron content in the interval ı =0−0.5. The transport numbers for oxygen ions varied inthe opposite fashion. With diminishing the partial oxy-gen pressure, the isotherms of the transport numbers forprotons exhibit a tendency toward an increase in oxidiz-ing atmospheres and to a sharp decrease in the reducingatmospheres [30]. The proton conductivity inCaTi1 − xFexO3 – δ increases with ı and reaches a maxi-mum at ı = 0.05 [9]. The proton conduction is practi-cally temperature-independent. The authors of [9] con-sidered possible reasons for this fact. The activationlessproton transport may point to a metallic conductiontype. But it is more likely that the activationless trans-port is merely apparent. With decreasing temperature,protons become less mobile, while the solubility ofwater vapor in the oxide increases. These processesmay counterbalance one another.

The authors of [27] investigated the dependence of theproton conductivity in SrTi1 − xFexO3 – δ (ı = 0−0.8) as afunction of the temperature and partial pressures ofoxygen and water vapor in the oxygen and water-vaporconcentration cells. They used the emf method andmeasured the electroconductivity. They discovered thatthe proton conductivity in the titanate studied was inde-pendent of the partial oxygen pressure in oxidizing con-ditions at intermediate pressures. The sign of the emf ofthe water-vapor concentration cell corresponded to thetransport of positively charged hydrogen defects. Thedependence of conductivity on of that kind could

be caused by proton H+ hopping between the OH– ions.According to the defect formation model [27], the con-centration of these ions is independent of . In reduc-ing atmospheres, however, the sign of the water-vaporcell’s emf altered to the opposite. The implication is thehydrogen transport is realized under these conditionsby negatively charged species, hydroxide ions éç– orhydride ions ç–, rather than by protons. With elevatingtemperature the emf altered its sign at higher partialoxygen pressures. For example, in the case of SrTiO3 at700°C, the emf’s sign remained invariant until 10–22

atm. At 1000°ë, the emf altered its sign as early as at10–12 atm. At low values of the partial oxygen pressure,the proton conductivity increased with decreasing pres-sure. According to the defect formation model [27], theconcentration of hydride ions depends on the partialoxygen pressure precisely in this manner. This fact gavethe authors of [27] grounds to hypothesize that the

pO2

pO2

hydride ions dominate the hydrogen transport in reduc-ing atmospheres. In not-too-strongly reducing condi-tions, i.e. when the emf’s sign corresponded to thetransport of positively charged species, the dependenceof the proton conductivity on H2O was close to thesquare root of H2O. This fact also pointed to the trans-port of the positively charged proton in the frameworkof the defect formation model [27]. The proton conduc-tivity in SrTi1 − xFexO3 – δ increased with temperature, asopposed to ë‡Ti1 − xFexO3 – δ, where it was temperature-independent [9].The authors of [31] discovered the proton conduction inSrTiO3 doped with 2 and 10 mol % of aluminum. Aswith SrTi1 − xFexO3 – δ, the emf of a water-vapor concen-tration cell altered its sign to the opposite at low partialoxygen pressures. On this basis, the authors of [31]spoke of the participation of negatively charged hydrideions in conduction. The proton conduction inSrTi1 − xAlxO3 – δ is independent of , whereas thehydride ion conductivity increases with decreasing and reaches the level of the oxygen conductivity. Theauthors of [31] believe that this phenomenon is due tothe migration of hydride ions over oxygen vacancies.The hydride-ion conductivity at extremely small valuesof the partial oxygen pressure exceeds the oxygen con-ductivity. The transport of interstitial hydride ionssimultaneously with the transport of hydride ions occu-pying the oxygen positions might have been responsi-ble for this phenomenon.

CONCLUSIONS

Titanates of alkaline-earth metals doped with accep-tor admixtures are interesting objects both in the scien-tific sense and in the sense of practical applications.Depending on composition and external parameters,they may exhibit the oxygen-ion, electron, and protonconduction and, who knows, the hydride-ion conduc-tion as well. The titanates are cheap and stable enoughmaterials. Once realized in the form of thin films, theycould successfully be used in assorted electrochemicaldevices.

ACKNOWLEDGMENTS

Author thanks V.P. Gorelik for valuable remarks.This work was supported in part by the Russian

Foundation for Basic Research, project no. 04-03-32 377.

REFERENCES

1. Esaka, T., Fujii, T., and Suwa, K., Solid State Ionics,1990, vol. 40-41, p. 544.

2. Marina, O.A., Pederson, L.R., and Stevenson, J.W., 2ndInt. Symp. on Solid Oxide Fuel Cell Materials and Tech-nology; 29th Int. Conf. on Advanced Ceranics and Com-posites, Cocoa Beach (FL), Jan 23–28, 2005,

pO2

pO2

900


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electrophysical properties of titanates of alkaline-earth metals

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