[membrane science and technology] fundamentals of inorganic membrane science and technology volume 4...

Fundamentals of Inorganic Membrane Science and Technology Edited by A.J. Burggraaf and L. Cot

�9 1996, Elsevier Science B.V. All rights reserved

Chapter 10

Dense ceramic membranes for oxygen separation

H.J.M. Bouwmeester and A.J. Burggraaf

Laboratory for Inorganic Materials Science, Faculty of Chemical Technology, University of Twente, P.O. Box 217, 7500 AE Enschede, The Netherlands

10.1 INTRODUCTION

Dense ceramic membranes exhibiting high oxygen ionic and electronic conductivity have become of great interest as a potentially economical, clean and efficient means of producing oxygen by separation from air or other oxygen- containing gas mixtures. In addition to infinite permselectivity, notably high oxygen flux values are measured through selected mixed-conducting oxides with the perovskite structure. These may be in the range exhibited by microporous membranes, albeit that sufficiently high temperatures are required, typically above about 700~

It is generally accepted that, provided they can be developed with sufficient durability and reliability, mixed-conducting oxide membranes have great potential to meet the needs of many segments of the oxygen market. It is further expected that the Oxygen fluxes can be improved by thin film deposition on a porous substrate preferably of the same material to avoid compatibility problems. The applications envisioned range from small-scale oxygen pumps for medical applications to large-scale usage in combustion processes, e.g. coal gasification [1-4]. As oxygen, but also nitrogen, ranks among the top five in the production of commodity chemicals in the United States [5] successful development of the mixed-conducting oxide membranes could thus have clear eco- nomic benefits, at the expense of market share from more traditional supply options. Whilst the targeted membranes will be most competitive at small and intermediate scale level in which flexibility of operation is desired, they may

4 3 6 10 n DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

eventually challenge the present commercial status of cryogenics, pressure- swing adsorption (PSA) and polymeric membranes [1-4].

Another application of mixed-conducting oxide membranes is to be found in the field of chemical processing, including the partial oxidation of light hydrocarbons, e.g. natural gas to value-added products such as ethane/ethene [6--11] and syngas [12-14], waste reduction and recovery [15]. The catalyst may be either the membrane surface itself or another material deposited in particulate form on top of the membrane. Besides the controlled supply or removal of oxygen to or from the side where the catalyst and the reactants are located, a promising feature is that the oxygen flux may alter the relative presence of different oxygen species (02,0-) on the catalyst surface, thereby providing species that may be more selective for partial oxidation reactions.

This review addresses recent developments in the area of mixed ionic-electronic conducting membranes for oxygen separation, in which the membrane material is made dense, i.e. free of cracks and connected-through porosity, being susceptible only for oxygen ionic and electronic transport. Current work on different mixed- conducting oxides is reviewed using concepts from electrochemistry and solid- state chemistry. Emphasis is on the defect chemistry, mass transport and the associated surface exchange kinetics, providing some basic background knowledge which aids further development of these materials into membranes for the aforementioned applications. There is no attempt to discuss inroads against competing technologies, or to speculate on new opportunities that may result from successful development. New developments in dense ceramic membrane research could offer very economical ways of separating hydrogen such as the proton-conducting ceramics or thin Pd-foils. These are not considered in this review. For a general discussion on the topical area of membrane technology and its impact in various applications the reader is referred to specific reviews, for example, see Refs. [16-20] and other chapters in this textbook.

10.2 GENERAL SURVEY

In this section, a brief overview is given of major membrane concepts and materials. Besides membranes made from a mixed ionic-electronic conductor (MIEC), other membranes incorporating an oxygen ion conductor are briefly discussed. Data from oxygen permeability measurements on selected membrane materials are presented.

10.2.1 Major Membrane Concepts

In this chapter, a membrane is regarded as a barrier between two enclosures which preferentially allows one gas (i.e. oxygen) to permeate owing to the presence of a driving force such as a pressure or electric potential gradient.

10 m DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION 437

PO 2 ' __ O2-.~

e

(a)

P 02, ,

*e

.__ 02"_>

(b)

e ,#

P02 ' P02 " 02

(c) (d)

Fig. 10.1. Different membrane concepts incorporating an oxygen ion conductor: (a) mixed conducting oxide, (b) solid electrolyte cell (oxygen pump), and (c) dual-phase membrane. Also shown is the schematics of an asymmetric porous membrane (d), consisting of a support, an intermediate

and a barrier layer havhlg a graded porosity across the membrane.

The separation of oxygen using an MIEC membrane is schematically shown in Fig. 10.1a. The driving force for overall oxygen transport is the differential oxygen partial pressure applied across the membrane. As the MIEC membrane is dense and gas-tight, the direct passage of oxygen moleculesis blocked, yet oxygen ions migrate selectively through the membrane. Dissociation and ionization of oxygen occurs at the oxide surface at the high pressure side (feed side), where electrons are picked up from accessible (near-) surface electronic states. The flux of oxygen ions is charge compensated by a simultaneous flux of electronic charge carriers. Upon arrival at the low pressure side (permeate side), the individual oxygen ions part with their electrons and recombine again to form oxygen molecules, which are released in the permeate stream.

4 3 8 10 - - DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

Mixed conduction also plays an important role in many other processes, e.g. in improving electrode kinetics and catalytic behaviour [21]. In fact, all oxides exhibit to some degree mixed ionic and electronic conduction, and selective oxygen permeation has been reported even for dense sintered alumina above 1500~ [22,23]. Although it is common to speak of mixed conduction when the total conductivity is provided by near equal fractions (transference numbers) of the partial ionic and electronic conductivity, respectively [24], from the point of view of oxygen permeation it is more useful to relate mixed conduction to their absolute values. Volume diffusion theories treating ambipolar transport in oxides clearly indicate that higher currents (fluxes) are obtained when either the electronic or the ionic conductivity increases, or both increase simultaneously. The flux at a given total conductivity is maximum when the ionic and electronic transference numbers are equal, i.e. 0.5. In this view, alumina is not a good mixed conductor. Materials showing predominant electronic conduction may thus prove to be excellent mixed conductors when their ionic conductivity is also substantial. The general objective for optimum membrane performance therefore is to maximize the product of mobility and concentration of both ionic and electronic charge carriers in appropriate ranges of temperature and oxygen partial pressure.

Owing to the ability to conduct both oxygen ions and electrons, the MIEC membrane can operate without the need of attachment of electrodes to the oxide surface and external circuitry. The latter represents an inherent advantage over traditional oxygen pumps in which a solid oxide electrolyte is sand- wiched between two gas-permeable electrically conductive electrodes (Fig. 10.1b). An advantage of electrically-driven oxygen separation may be its ability to deliver oxygen at elevated pressures, eliminating the need for compressors [25]. Figure 10.1c shows a dual-phase membrane, which can be visualized as being a dispersion of a metallic phase into an oxygen ion conducting host or matrix, e.g, Pd metal into stabilized zirconia. This challenging approach was first described by Mazanec et al. [26] and offers an alternative use of oxide electrolytes in the field of dense ceramic membranes. Industrially important solid oxide electrolytes to date are mainly based on oxygen-deficient fluorite-related structures such a s Z r O 2 and C e O 2 doped with CaO o r Y 2 0 3 . Unless operated with an internal or external circuitry, the oxygen flux through these materials in usual ranges of temperature and oxygen pressure is negligibly low, prevent- ing their practical use as oxygen separation membrane. The existence of a non-vanishing electronic conduction in the ionic domain, and concomitant oxygen semipermeability, however, can be detrimental considering their use as solid electrolytes in fuel cells and oxygen sensors [27,28].

While past efforts were focused on expanding the electrolytic domain of oxygen ion conducting fluorite-type ceramics, more recently one has begun to introduce enhanced electronic conduction in fluorite matrices. Extrinsic elec-

10 - - DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION 4 3 9

tronic conduction in ionically conducting matrices can be obtained by dissolution of multivalent cations in the fluorite oxide lattice. Notable examples include yttria-stabilized zirconia doped with either titania [29,30] and ceria [31, 32]. Electronic conductivity in these solid solutions is reportedly found to increase with increasing dopant concentration, but may be limited by the solid solubility range of the multivalent oxide. As conduction occurs via a small polaron mechanism (electron hopping) between dopant ions of different valence charge, its magnitude will strongly vary with temperature and oxygen partial pressure. In general, the extent of mixed conductivity that can be induced in fluorite ceramics is limited, which restricts its possible use as ceramic membrane, unless very high temperatures of operation (> 1400~ and stability down to very low values of oxygen partial pressure are required as, e.g., in the production of gaseous fuels CO and H2 by direct thermal splitting of CO2 and H20, respectively, and extraction of the oxygen arising from dissociation [33].

Since the first report on high oxide ion conductivity in some of the rare earth aluminates in the mid sixties [34,35], materials with oxygen-deficient perovskite and perovskite-related structures receive much attention for the development of new solid electrolytes and mixed conductors for numerous applications [36]. Currently, extensive research is conducted on acceptor-doped perovskite oxides with the generic formula Lal_xAxCOl_yByO3_~ (A = Sr, Ba, Ca and B = Fe, Cu, Ni). Teraoka et al. [37-39] were the first to report very high oxygen fluxes through the cobalt-rich compositions, which perovskites are known to become highly oxygen anion defective at elevated temperatures and reduced oxygen partial pressure. The oxygen-ion conductivity in the given series can be 1-2 orders of magnitude higher than those of the stabilizedzirconias, though in usual ranges of temperature and oxygen partial pressure electronic conduction in the perovskite remains predominant [39,40]. Besides potential use of these perovskite compositions as catalytically active electrodes in, e.g. fuel cells, oxygen pumps and sensors, the compounds have a bright future for use as oxygen separation membrane. The precise composition may be tailored for a specific application, but this has not yet been fully developed. Structural and chemical integrity of the cobaltites, however, is a serious problem and needs to be addressed before commercial exploitation becomes feasible.

For the sake of completeness, a schematic representation of a porous ceramic membrane is given in Fig. 10.1d. The majority of porous ceramic membranes are composite or asymmetric in structure. They include materials like 0~-A1203, ~-A1203, TiO2 and SiO2, and generally consist of a thin layer of either a mesopor- OUS (2 <dp < 50 nm) or microporous (dp < 2 nm) barrier layer of a few microns thick superimposed on a mechanically strong support with a comparatively large pore diameter dp and usually a few millimetres thick. Often there are one or more intermediate layers, resulting in a graded pore structure across the membrane.

440 10 -- DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

10.2.2 Data: Oxygen Permeability of Solid Oxide Membranes

Table 10.1 lists data of steady-state oxygen permeability measurements on various solid electrolytes, mixed conducting oxides and dual phase membranes taken from various literature reports. Measurements most commonly are performed by imposing a gradient in oxygen partial pressure across the membrane, usually by passing an oxygen rich and lean gas, e.g. air and inert gas, respectively, along opposite sides of a sealed ceramic disk or tube wall, without the use of external circuitry such as electrodes and power supplies. The number of moles, volume or mass of oxygen passing per unit time through a unit of membrane surface area is measured down stream using, e.g., on-line gas chro- matography or an oxygen sensor, from which data the oxygen flux is calculated. Table 10.1 also includes data for solid electrolyte cells used in the oxygen pump mode. A graphical presentation of selected data is given in Fig. 10.2.

-5.00

-6.00

'm -7.00 o,I

E o

o -8.O0 E

N

0 . . . . ~

o'~ - 9 . 0 0 o

-10.00

-11.00

(1) ' + ~

"~ (9) ~ �9

r

. . . . . . i . . . . . .

0.65 1.00 1.35

1000/T [ l /K]

Fig. 10.2. Arrhenius plots of oxygen permeation for: (1) La0.3Sr0.K~oO3-8 [7]; (2) Ba0.9Y0.1CoO3-6 [19]; (3) YSZ-Pd (40 vol%), continuous Pd-phase [23,24]; (4) YSZ-Pd (30 vol%), discontinuous Pd-phase [23,24]; (5) La0.sSr0.sCoO34 [7]; (6) (ZrO2)0.7-(Tb203.5)0.228-(Y203)0.072 [4,5]; (7) SrCo0.sFe0.203-~ [9]; (8) BE25-Ag 40 vol% [24,29]; (9) Ba0.66Y0.33CoO34 [19]; (10) (Bi203)0.75-(Er203)0.25 (BE25) [22]; (11) BY25-Ag (35 vol%), thickness 90 wn [27]; (12) BY25-Ag (35 vol%), thickness 1.5 mm [27]. Air and inert gas are passed along opposite sides of the membrane. Unless specified otherwise, the membrane

thickness varies between I and 2 mm. For references, see footnote to Table 10.1.

10 m

DE

NS

E C

ER

AM

IC M

EM

BR

AN

ES

F

OR

OX

YG

EN

SE

PA

RA

TIO

N

447

6- ~ !

~ m

~ "Q

"~

t',.

�9 ,~

.-,,~,m

.

~ S

N

~ .

I

! .

~'~ T

"~ "~ ~

�9 ~

~

~ ~

~ ~

-~

~-

~ ~

~~

7

2~

z

~ .

~0

"i ~

'-u.

~ ~

u5 _:,_:, u~ ~ u.iu~

~::::~

~~

0~

:~'~

B

-~O

:~ 0~~


High oxygen fluxes are found through selected perovskite-structured ceramics. Table 10.1 shows for several perovskite systems the trend in permeation flux as a function of the type and concentration of applied dopants. In the range 800-900~ the highest flux was measured by Teraoka et al. [37] for SrCo0.sFe0.2034 but, as for a number of other compositions, different values have been reported by other groups. Such conflicting results reflect the experimental difficulties in measuring oxygen permeation of sealed ceramic discs at high temperatures, but may also be due to factors that influence the effective P%-gradient across the membrane, sample preparation, etc. This is further discussed in Section 10.6.7.2.

For the sake of comparison, Table 10.1 contains limited data for the oxygen flux through micro- and mesoporous membranes. As noted before, the last- mentioned category of membranes falls outside the general scope of this chapter. It is seen that the oxygen fluxes observed through membranes formed from the mixed-conducting perovskite-type oxides, such as La>xSrxCo1_yFeyO34, approach those exhibited by the porous membranes. It should be noted, however, that these types of membrane have different requirements. The high temperature needed for operation using membranes based on oxygen ion conductors may be restrictive in certain applications, but beneficial to others, e.g. coal gasification and partial oxidation of light paraffins [25].

10.2.3 Factors Controlling Oxygen Permeation

The rate at which oxygen permeates through a non-porous ceramic membrane is essentially controlled by two factors, the rate of solid state diffusion within the membrane and that of interfacial oxygen exchange on either side of the membrane. The oxygen flux can be increased by reducing the thickness of the membrane, until its thickness becomes less than a characteristic value, Lc, at which point the flux of oxygen is under conditions of mixed control of the surface exchange kinetics and bulk diffusion [41]. Below L c, the oxygen flux can only marginally be improved by making the membrane thinner.

For predominant electronic conductors like, for example, the perovskites Lal_xSrxCOl_yFeyO3~, Lc is determined by the ratio of the oxygen self diffusivity and surface exchange coefficient. Both parameters can be measured simultaneously by using 180-160 isotopic exchange techniques. Calculations show that Lc may vary from the ~tm-range to the cm-range, depending on material and environmental parameters. Modelling studies, however, show that significant increase in the rate of interfacial oxygen transfer and, hence, in the oxygen flux can be achieved by deposition of a porous MIEC layer on top of the (thin) non-porous membrane [42--44]. Since a number of simplifying assumptions is made, such as neglect of changes in material parameters with variation in the chemical potential of oxygen, the models developed are valid only in the limit of small Po2-gradients across the MIEC membrane. For a more rigorous approach,


referring to actual operating conditions of oxygen separation membranes, much more work is needed to arrive at a better understanding of the transport processes under oxygen potential gradients. In particular, our present understanding of the factors that govern the surface exchange kinetics is rather poor.

Effects related to microstructure, including grain boundary diffusion and (local) order-disorder phenomena, may also influence overall oxygen transport. Besides the processing into defect-free thin films and associated problems of compatibility between deposited membrane layer and the porous substrate material, chemical stability at high temperatures, effects induced by the presence of an oxygen potential gradient like segregation of impurities to the surface and to grain boundaries, kinetic demixing and kinetic decomposition could affect membrane performance or limit operational life. In many cases, these difficulties remain to be overcome before commercial exploitation becomes viable. All these factors are important and govern the selection of materials.

In the following sections, the emphasis is on the basic elements of mixed ionic and electronic transport through dense ceramic membranes. Due to size considerations, we shall mainly focus this chapter to mixed-conducting acceptor- doped perovskite and perovskite-related oxides. Other membrane concepts are also discussed, but only briefly. The examples chosen illustrate the fundamental factors determining the oxygen fluxes through dense ceramic membranes, which is the primary aim of this chapter.

10.3 FUNDAMENTALS

10.3.1 Bulk Transport

The basic assumption of the theory presented in this section is that the lattice diffusion of oxygen or the transport of electronic charge carriers through the bulk oxide determines the rate of overall oxygen permeation. Moreover, oxygen is transported selectively through the membrane in the form of oxygen ions, rather than molecules, under the driving force of a gradient in oxygen chemical potential. The flux of oxygen ions is charge compensated by a simultaneous flux of electrons or electron holes, which is enabled without the use of external circuitry. We only briefly review the fundamentals of solid state diffusion through mixed conducting oxides and the reader is referred to Refs. [45-47] for a more complete discussion.

10.3.1.1 Wagner Equation

Considered here is the case where the interaction of gaseous oxygen with the oxide lattice can be represented by a chemical reaction of the form 1

1 The notation adopted for defects is from Kr6ger and Vink [48].

450 10 ~ DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

1 -~ 0 2 4- W 6 4- 2 e '= O~) (10.1)

assuming that oxygen vacancies are the mobile ionic defects. These may be obtained, e.g., by doping of the oxide lattice with aliovalent cations. The intrinsic ionization across the bandgap can be expressed by

nil = e' + h. (10.2)

The single particle flux of charge carriers, wi th neglect of cross terms be tween fluxes, is given by

(3" k

jk = z~ k F2 Vnk (10.3)

where Z k is the charge number and C~k the conductivity of charge carrier k, F the Faraday constant and Vrlk the gradient of the electrochemical potential. The latter comprises a gradient in chemical potential V~tk and a gradient in electrical potential VO, for each individual charge carder k given by

VT~k = V~k + Z k F V~) (10.4)

The charge carrier diffusing more rapidly causes a gradient in the electrical potential V~, in which the t ransport of carriers wi th opposite charge is acceler- ated. At s teady state, no charge accumulat ion occurs. The fluxes of ionic and electronic defects are therefore related to each other by the charge balance

2 j v 6 = je" - jh. (10.5)

Equat ion 10.5 can be used together wi th Eqs. (10.3) and (10.4) to eliminate the electrostatic potential gradient. The flux of oxygen vacancies is then obtained in terms of the chemical potential gradients only. If it is further a ssumed that internal defect chemical reactions are locally not dis turbed by the t ransport of matter, the chemical potential gradients of individual charge species can be converted into the virtual chemical potential of gaseous oxygen, ~to2. The following differential relations hold at equil ibrium 2

2 It is tacitly assumed here that the chemical potential of lattice oxygen ~to~) is constant. The present formulation of the defect equilibrium for the formation and annihilation of oxygen vacancies and electrons by the reaction of the solid with environmental oxygen, however, is written in terms of the 'virtual' chemical potentials of the constituent structure elements. In so doing, one does not properly take into account the so-called site-exclusion effect, because the chemical potential of the oxygen vacancy 1/6 and that of lattice oxygen O~) cannot be defined ind- ependently from one another. In the present context, it suffices to say that the derived equations are in agreement with those obtained from a more rigorous thermodynamic treatment based upon the 'true' chemical potential for the building unit vacancy, i.e. (V6-O~). For further reading concerning the definition of chemical potentials, the reader may consult Refs. [45] and [46].

10 u DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION 451

1 -~ V~to2 + V~tv6 + 2V~t e, = 0 (10.6)

V~I, e, 4- V~h.---- 0 (10.7)

where ~tv6 denotes the chemical potential of the oxygen vacancy, ~te' and ~h. denoting the chemical potential of electrons and electron holes, respectively. The flux of oxygen through the membrane can be derived by combining Eqs. (10.3)-(10.7), using the relationship Jo2 = -1 /2 jr6. One finds

1 ~ ((~e'4- (~h.)(~V 6 = - J . . . . . . . . . J V~t02 (10.8)

J02 42F 2 k ((~e' 4- (~h.) 4- O'V 6 j

or in a more generalized form

1 (~elt~ion Jo2 = -- 42 F 2 V~to2 (10.9)

Gel 4- (~ion

where (~ion -- (~V 6 and (3"el ---- (~h. + (~e' are the partial ionic and electronic conductivity, respectively. The conductivity term in Eq. (10.9) is equivalent to teltiont~total =

tiont~el = telt~ion, where tel and tion a r e the fractions (transference numbers) of the total conductivity (~total provided by electronic and ionic defects, respectively. Integration of Eq. (10.9) across the oxide membrane thickness, L, using the relationship V~to2 = ORT In Po2/Ox (x = distance coordinate) and assuming no divergence in the fluxes, yields the Wagner equation in the usual form

In P"o2

RT (~elt~i~ d In Po2 (10.10) Jo2 = - 42 f 2------- ~ ~ (~el + (~ion

lnP'02

The limits of integration are the oxygen partial pressures maintained at the gas phase boundaries. Equation (10.10) has general validity for mixed conductors. To carry the derivation further, one needs to consider the defect chemistry of a specific material system. When electronic conductivity prevails, Eqs. (10.9) and (10.10) can be recast through the use of the Nernst-Einstein equation in a form that includes the oxygen self-diffusion coefficient Ds, which is accessible from ionic conductivity measurements. This is further exemplified for perovskite- type oxides in Section 10.6.4, assuming a vacancy diffusion mechanism to hold in these materials.

10.3.1.2 Chemical diffusion coefficient

The preceding theory was used by Wagner to describe oxide film growth on metals [49,50]. The driving force for diffusion is not a concentration gradient,


but rather a chemical potential gradient. An important and necessary assumption is that the internal defect reactions are fast enough to attain local chemical equilibrium so that the concentrations of involved ionic and electronic (electrons or holes) charge carriers at any distance coordinate in the oxide are fixed by the local value of the virtual chemical potential, ~to2. The effective transport is still that of neutral oxygen atoms by which the theory fits that of a chemical diffusion process in terms of Fick's first law

- 3c o jo = - D 3x (10.11)

where the driving force for diffusion is the gradient in neutral oxygen, 3Co/3X. The coefficient of proportionality, denoted by D, is called the chemical diffusion coefficient. By virtue of Eqs. (10.9) and (10.11 ), one obtains

~ _ 1 (3"el(3"i~ 3~t~ (10.12) 8 F 2 (~el + ($ion OCO

Here we note that Jo2 = 1/2jo. Because 3co/3X =-3Cv/3x, a similar expression is obtained when diffusions were dominated by neutral vacancies.

The thermodynamic factor 3~o/3Co in Eq. (10.12) can be determined directly from experiment by measuring the oxygen stoichiometry as a function of oxygen partial pressure, either by gravimetric or coulometric measurements. In view of Eqs. (10.6) and (10.7), it comprises contributions from both ionic and electronic defects, which reflect their non-ideal behaviour. For materials with prevailing electronic conductivity Eq. (10.12) may be siml~lified to yield an exact relation between the chemical diffusion coefficient D and the oxygen tracer diffusion coefficient D*:

- D* 1,/23~to2 D = ~ (10.13)

H R R T 3 In Co

Here Ha is the Haven ratio, defined as the ratio of the tracer diffusion coefficient D* D ~ to the quantity derived from dc ionic conductivity measurements,

(~ion RT D a = ~ (10.14)

c o Z 2 F 2

The Haven ratio may deviate from unity when correlation effects and possibly different jump distances and jump frequencies can not be neglected [51]. For a vacancy diffusion mechanism Ha equals the well-known tracer correlation factorf.

10 -- DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION 453

10.3.1.3 Trapping of Electronic and Ionic Defects

Equation (10.10) and those derived from it are valid as long as fully ionized oxygen defects contribute to transport. Different equations are obtained if valency changes of oxygen defects occur. Wagner [50] proposed to put the influence of reactions between ionic and electronic defect species in the cross terms of the Onsager equations. Maier [52-54] explicitly attributed individual diffusivities and conductivities to the new defect species, using the concept of a conservative ensemble accounting for free and trapped species. Following his approach, the reversible reaction between electrons and oxygen vacancies,

Vo + e ' : Vo

Vo + e ' : V ~

(10.15)

leads to the following expression for the oxygen flux,

i [ Jo2 = 42 F214s[ v o ~ +

((~e' + (~h') ((~V 6 + 4rYv6) + (~v 6 r~v 6 ] I V,o2 (10.16)

((~e" + (~h') + ((~V 6 + (~V 6) J

where we have adapted Eq. (33) in Ref. [53] (Part I) i n toa form to be similar to Eq. (10.8), in which ionic transport is by doubly ionized oxygen vacancies only. The Onsager coefficient S Wo accounts for the contribution of neutral defects, enabling oxygen transport even when the electronic conductivity of the oxide is zero. We further note that the counter-diffusion of two Vo and a single Vo would result in a net neutral oxygen flux, as reflected by the last term in the numerator of Eq. (10.16). Maier [53] also examined the case in which electronic or ionic defects are associated (trapped) with immobile centres such as dopant ions. Trapping inevitably leads to a decrease in concentrations of the charge carriers available for transport. The impact of these phenomena is that the transport equations for evaluation of data obtained from electrochemical measurements like, for example, ionic conductivity, concentration cell, permeability and Hebb-Wagner polarization experiments should accordingly be modified. It is shown by Maier how these are influenced by trapping effects observed in perovskite SrTiO3, and by the transport properties of the high-temperature superconductor YBa2Cu306+x. Because of the large oxygen excess possible in the latter material it is assumed that transport occurs by differently ionized ionic defects, partly even by neutral oxygen species. For references, see the papers by Maier [52-54].

454 10 - - DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

10.3.1.4 Empirical Equations

Evaluation of j% from Eq. (10.10) requires that data exist for the partial conductivities r~io~ and (3"el as a function of oxygen partial pressure between the limits of the integral. In what follows, some special relations for either prevailing electronic or ionic conduction are discussed.

For the sake of approximation, in defect chemical studies often an empirical power law is used for the partial conductivity of the rate determining species, O'ir s u c h a s r

odP%) = r~ ~ P~2 (10.17)

where ~o is the conductivity at standard state. The value of n can be derived from experimental data of steady-state oxygen permeation. For proper evaluation it is necessary that the Po~-gradient across a specimen is varied within the assumed range of validity of the empirical power law. Inserting Eq. (10.17) in Eq. (10.10), one finds after integration, assuming ~i << (~totab

._ ~_ __ ppYt J% [P'~)2 P 021 (10.18)

where ~ = r~~ For large positive values of n the rate of oxygen permeation is predominantly governed by the oxygen partial pressure maintained at the feed side, P02'. Likewise for large negative values it is predominantly governed by the oxygen partial pressure maintained at the permeate side, P02". For either a small value of n or a small Po2-gradient, the flux becomes proportional to In (Po~'/P02").

Provided that the electronic transference number is known, the ionic (and electronic) conductivity may be obtained by differentiation of experimental data. Assume that we have produced a data set for different P%-gradients, keeping the oxygen partial pressure P%" at the permeate side fixed. Differenti- ating Eq. (10.10) with respect to the lower integration limit yields,

/ La

RT = tel(~io n X 42 F 2L (10.19)

The ionic conductivity at a given pressure Po2 is thus obtained from the slope of the jo~- In Po2' plot at that Po2. Similarly, the ionic conductivity can be evaluated from oxygen flux values measured by varying the oxygen partial pressure at the permeate side, keeping the one at the feed side fixed. The two data sets in general will yield the ionic conductivity over a complementary range in oxygen partial pressure.


So from oxygen permeability measurements, it is possible to get information on the transport properties of oxides without making use of electrodes and external circuitry. In practice, however, experiments are plagued by the usual problems of sealing oxide ceramic discs at high temperature. A further compli- cation is that the activity of oxygen at the oxide surfaces may not be precisely known, due to the influx and effiux of oxygen. In addition to these polarization effects, the results may be significantly affected by the flow patterns that possibly exist on feed and permeate sides.

It was further assumed by Wagner that the surface reactions proceeding at the gas phase boundaries have reached a quasi-equilibrium condition relative to diffusion through the oxide. However, awareness is growing that in many cases the surface reaction may exert a partial control over the transport kinetics [41]. The extent of surface control varies with membrane thickness, temperature and oxygen pressure difference imposed across the membrane. Other limitations of the theory, some of which are briefly discussed in subsequent sections, include solid state diffusion of matter along preferred paths such as grain boundaries, porosity and, especially at large departures from the ideal stoichiometric composition, the formation of point defect clusters and ordering.

10.3.2 Surface Oxygen Exchange

The exchange of oxygen between oxide surfaces and the gas phase has been recognized to involve a series of reaction steps, each of which may be rate determining [21]. Possible steps for the reduction of oxygen include adsorption, dissociation, charge transfer, surface diffusion of-intermediate species, e.g., O2ads, Oad s and Oads, and finally incorporation in the (near-) surface layer. Generally, it is assumed that the reoxidation of oxygen anions follows the same series of steps in the reverse direction. The surface reaction is thus associated with transport of charge. The charge carriers in the interior of the oxide main- tain thermodynamic equilibrium, according to Wagner's theory, but not at the surface if rate limitations by the surface exchange kinetics come into play.

Liu [55] has presented a detailed analysis of the oxygen separation rates of mixed conductors, using the well-known Butler-Volmer formalism to model interfacial mass and charge transfer. At first glance such a description is not applicable because of the absence of electrodes. But, the adsorbate may be regarded as to replace the electrode material. Accordingly, the electrical double layer (Helmholtz layer) formed between the ionosorbed adsorbate and the oxide dominates the charge transfer kinetics, thus leading to the introduction of transfer coefficients in the rate constants for reduction and oxidation.

In general, however, even at equilibrium, there may also be a double layer, called space charge, next to the surface extending into the oxide interior (Mott- Schottky layer). The width of the space charge layer will be of the order of the

4 5 6 10 - - D E N S E C E R A M I C M E M B R A N E S FOR O X Y G E N S E P A R A T I O N

length, LD, the value of which depends on the volume concentration of all mobile charge carriers. Alteration of the space charge leads to a change in bending of the energy bands, and this influences the occupation of electronic levels near the band edges at the oxide surface. The total potential drop across the interface thus becomes distributed between the Helmholtz layer and the space charge layer [56], which complicates the analysis of the charge transfer kinetics. Furthermore, if strong electrical fields are developed in a situation where L D is greater than the size of the ionic charge carrier, which is favoured by a low concentration of mobile charges, the migration of ionic charge carriers through the space charge zone can be described by an expression similar to the Butler-Volmer equation [57,58].

The possibility of space charge in ionic crystalline solids, and the attendant redistribution of lattice components, has a great influence on the properties of boundary regions. Their defect chemistry may depart considerably from the predominant one in the bulk [59]. In oxides, there are quite a number of experimental observations to support the existence of space charge induced segregation [60,61]. Other factors have been recognized to contribute to segregation, like the misfit strain energy of the solute ion and the surface tension considering adsorption equilibria. Owing to the segregation phenomenon, the composition and crystal ordering of oxide surfaces and grain boundaries differ from that in the bulk. In some cases this leads to the formation of a second phase. It is clear that these phenomena will have a significant, often controlling influence on the properties of oxide materials, which includes the heterogeneous kinetics of the gas/solid interface.

10.3.2.1 Characteristic Membrane Thickness

A simple yet valuable criterion for candidate membrane material selection is the characteristic membrane thickness Lc [41]. In the theory, the transport equations for diffusion in the solid, and for the surface exchange are linearized. It should therefore strictly be used when small Podgradients are imposed across the membrane. In the next section, methods for measuring Lc are briefly discussed.

The starting point is to divide the membrane into a central bulk (Wagner) zone and adjacent interfacial zones, emphasizing the importance of both solid state diffusion and surface oxygen exchange to the magnitude of the oxygen flux. This is schematically shown in Fig. 10.3. The available driving force

�9 total is distributed across the membrane such that the rate determining proc- A~o 2 ess receives the greater proportion. Diffusion is rate determining, if the membrane is made sufficiently thick. Upon reducing the thickness, the flux is controlled by the limited transfer of oxygen across the interfaces. Within the approximation of linear kinetics for diffusion and interfacial exchange, the flux equation can be written as [41]

10 - - DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION 457

interracial zone

bulk in ter fac ia l

zone

PO ' l

T

I I I

"~ I I ~1 I

t -- . . . . I

I ~ . I I ' ~ . I I ~. I " , , I

I

I I ~,~ I I ,, I I I I �9

PO=

Fig. 10.3. Drop in chemical potential 1102 across bulk m~d interfacial zones of a membrane imposed to an oxygen partial pressure difference, P02' > P02". The largest drop occurs across the least permeable

zone .

A , total 1 teltion(Stotal /-l~tO 2

J02 = - 1 + (2Lc/L) 42 F 2 L (10.18)

noting that in the limit of small Po2-gradients, the exchange rates at opposite interfaces will be identical. The quanti ty Lc, has been called the characteristic thickness,

RT teltion (~total Lc = 42 F 2 x Je ~ (10.19)

and determines the transition from predominant control by diffusion to that by surface exchange. The parameter j~ (mol 02 cm -2 s -1) expresses the balanced exchange rate in the absence of oxygen potential gradients and relates to the surface exchange coefficient ks (cm s -1) accessible from data of 180-~60 isotopic exchange,

je ~ = 1/4 ks Co (10.20)

in which co is the volume concentration of oxygen anions at equilibrium. Exchange rates are customarily defined in terms of moles of anions or molecules per unit time and area. Often the geometric area is used in calculation, in spite of the fact that the true surface area on a microscopic scale may be substantially larger.

Compar ing Eq. (10.18) with the Wagner equation, we see that the diffusional flux of oxygen across the membrane is reduced by a factor (1 + 2Lc/L) -~, relative to that in the absence of transfer limitations across the interfaces. In general, the surface exchange rate will be different in, e.g., O2-N2, CO-CO2, H2-H20 atmospheres even at approximately equal oxygen partial pressures. If opposite sides of the membranes are exposed to different gas ambients, then the interface with the lesser performance dictates the overall exchange behaviour.


1 .00 ~ _. 1 .00

0 0.10 -,=,=,,t

0 .80 ),

oo . . . . . . . . . . . . .

. . . . �9 , L - " " " " 7 | "

0.10 1.00 10.00

0.60

0 . 4 0

0 . 2 0

0.01 0.00 0.01 100.00

L/L,, Fig. 10.4. Thickness dependence of the dimensionless oxygen flux j'o2 calculated using Eq. (10.18). The quantity j'o2 is defined by the ratio of the oxygen flux over the maximum achievable oxygen flux in the surface exchange limited regime. Only if L >> Lc, the oxygen flux becomes proportional to 1 / L ~with 7 = I in agreement with the Wagner theory. For smaller thicknesses, 7ranges between I and 0

(right-hand scale).

The influence of the membrane thickness on oxygen flux is shown in Fig. 10.4. When L > Lc, the oxygen flux varies inversely with L v where g = 1, in agreement with Wagner's theory (cf. Eq. (10.10)). Departures from this inverse relationship are observed when the oxygen flux becomes partly governed by the surface exchange kinetics. The value of 7, at a given L, corresponds with the negative slope in the double logarithmic plot of the oxygen flux versus membrane thickness at that L. Taking the logarithm of Eq. (10.18), partial differentiation with respect to log L shows that y is equal to the reduction factor (1 + 2Lc/L) -1, the value of which gradually decreases with decreasing thickness to become zero for L << Lc, as is shown in Fig. 10.4. The latter situation corresponds with the maximum achievable flux, which for a symmetrical membrane is given by: 1/2 "o -- total

]ex' A~l.O2 �9 Equation (10.19) may be simplified for predominant electronic conduction,

assuming that the classical Nernst-Einstein relationship can be represented as

c o D s z 2 F 2 O'i~ -- RT (10.21)

where Ds is the self-diffusion coefficient of oxygen anions with valence charge zo (= -2). Making the appropriate substitutions, the characteristic thickness Lc becomes

Ds D* Lc- ks - k--7 (tel = 1) (10.22)

10 m DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION 4 5 9

In the second part of Eq. (10.22) the fact has been used that, if correlation effects can be neglected, the tracer diffusion coefficient, D*, is equal to the self-diffusion coefficient, Ds. It is important to note once again that Eq. (10.18) is valid in the limit of small Podgradients only. Since both D* and ks for a given material are a function of its specific defect chemistry, in general, Lc will be a function of process parameters Po2 and temperature.

The picture that emerges is that, at given experimental conditions, for thicknesses below Lc no appreciable gain in the oxygen flux can be obtained by fabricating thinner membranes, unless the value of ks can be significantly increased. Similar criteria can be formulated for fuel cell electrodes as has been advanced by Steele [62,63] and Kleitz [64]. In the limit of small overpotentials (low field approximation), i.e. assuming ohmic behaviour for the relevant surface kinetics, whatever the rate controlling mechanism, the electrode resistance can be correlated with the electrolyte resistivity.

10.3.2.2. Measuring Lc

The quantity D*/ks appears to be a fundamental parameter governing tracer diffusion bounded by a limiting surface exchange between lattice oxygen and oxygen from the 1802-enriched gas phase. Experimental methods for isotopic exchange include monitoring the change of 180 2 concentration in the gas phase upon exchange in a fixed volume of 1802-enriched oxygen using a mass spec- trometer [65-69], weight measurements [70], depth probing of 180/(180 + 160) diffusion profiles using secondary ion mass spectroscopy (SIMS) after exchange at high temperature for a selected time [71-73] and combined approaches [74-76]. Fitting the acquired data to the appropriate diffusion equation allows both D* and ks to be obtained from a single experiment.

Selected data of D* and ks from 180-160 isotope exchange measurements of perovskite-type oxides are compiled in Table 10.2. High D* and ks values are reported for the ferrites and cobaltites, in which solids the anions are assumed to move via a vacancy diffusion mechanism. The mixed perovskites are notable for being excellent electronic conductors. Usually the electronic conduction is found to be predominant, in spite of the fact that the ionic conductivity may also be substantial (Section 10.6). The examples given in Table 10.2 clearly empha- size the importance of the surface exchange kinetics, relative to diffusion, in limiting overall oxygen transport through the perovskites.

Analyzing published data for D* and ks, Kilner [73] noted that the two parameters seem to be correlated. A square root dependence of ks with D* is found for perovskite oxides ABO3, albeit that the data show substantial scatter, yielding an average value for D*/ks of about 100 gm. For fluorite oxides M O 2

the two parameters are correlated almost linearly, while the corresponding value of D*/ks ranges between mm and cm. The results suggest that related


TABLE 10.2

Tracer-diffusion coefficient D*, surface exchange coefficient ks and characteristic thickness Lc for selected perovski te- type oxides

Perovskite T PO2 D* ks (b) 1 Lc (a) Ref. (~ (kPa) (cm 2 s -1) (cm s- ) (cm)

Lao.sSro.sMnO3-6

Lao.sSr0.2CoO3,5

Lao.6Cao.4Coo.8Feo.203-5

Lao. 6Sr0.4C00.8Ni0.20345

Lao.oSro.4Coo.6Nio.403.6 (c)

Lao.oSro.4Coo.4Nio.603-6

LaCoO3_8 (single crystal)

LaFeO3_5 (single crystal)

Lao.9Sro.lCoO3.8

Lao.gSro.lFeO3.6

Lao.oSro.4FeO3.6

700 70 2x10 -15 1x]0 -8 2x10 -7 1 800 8x10 -14 l x l 0 -7 8x10 -7 900 3x10 -12 9x10 ~ 3x10 -5

700 70 l x l 0 -8 3x10 ~ 3x10 -3 1 800 2x10 -8 5x10 -6 4x10 -3 900 4x10 ~ 2x10 -5 2xlO -3

700 70 2x10 ~ 4x10 ~ 5x10 -3 1 800 1x10 -7 2K10 -s 5x10 -3 900 3x10 -7 4x10 -5 7x10 -3

700 70 3x10 -8 2KIO -6 2x10 -2 2 800 lx10 -7 2x10 -6 5x10 -2 900 4XlO --7 2XlO 4 2xlO q

700 70 2X10 -9 7x10 -7 3XlO -3 2 800 6x10 ~ 3x10 -6 2x10 -2 900 3x10 -7 3x10 -6 lx10 -1

700 70 lx10 ~ 3x10 -7 3x10 -2 2 800 7x10 ~ 2x10 -6 3x10 -2 900 6x10 -7 2x10 -4 3x10 q

700 4.5 9x10 -13 1x10 -9 4x10 -4 3 800 2XlO -11 3KI0 -7 7x10 -5 900 6x10 q~ lx10 -6 4x10 -4

900 7 lx10 -i2 4x10 -8 3x10 -5 4 1000 5x10 q2 2x10 -7 3x10 -5

900 4.5 3x10 -9 lx10 -6 2XI0 -3 5 1000 2x10 ~ 2x10 ~ l x l 0 -2

900 6.5 3x10 -9 5x10 -7 6x10 -3 5 1000 lx10 -8 2x10 -6 6x10 -3

1000 6.5 6x10 -7 lx10 -5 5x10 -2 5

(a) Calculated using Eq. (10.22).

(b) Value can be equated to ]~• in accordance with Eq. (10.20), by multiplication wi th a factor 0.022. Strictly speaking, this holds for LaCoO3. (c) Authors report a two-phase mixture [2].

REFERENCES USED IN TABLE 10.2:

1. S. Carter, A. Selcuk, J. Chater, R.J. Kajda, J.A. Kilner and B.C.H. Steele, Solid State Ionics, 53-56 (1992) 597--605.

2. Ch. Ftikos, S. Car te r and B.C.H. Steele, J. European Cer. Soc., 12 (1993) 79-86.

3. T. Ishigaki , S. Yamauch i , J. Mizusaki , K. Fueki and H. T a m u r a , J. Solid State Chem., 54 (1984) 100-107.

4. T. Ishigaki , S. Yamauchi , J. Mizusaki , K. Fueki and H. T a m u r a , J. Solid State Chem., 55 (1984) 50-53.

5. T. Ishigaki , S. Yamauch i , K. Kishio, J. Mizusak i and K. Fueki, J. Solid State Chem., 73 (1988) 1 r w t - ~ . t ~ t , - n


point defect processes must be common to both diffusion and surface oxygen exchange [73]. However, mechanisms responsible for the apparent correlations remain obscure, reflecting our poor knowledge at present of the factors that control the oxygen exchange kinetics.

As discussed in previous work by the present authors [41], calculations of Lc using combined data from ionic conductivity measurements and 180-160 isotopic exchange for a number of fluorite oxides were found to be in good agreement with values estimated from oxygen permeability measurements.

Finally, it may be noted that also relaxation methods offer a useful tool for measuring Lc, without the requirement of high-temperature seals as in oxygen permeation experiments. The (re-)equilibration may follow after an instantane- ous change of the oxygen activity in the gas phase and involves the propagation of a composition gradient through a thin slab or single crystal of the oxide. The change in stoichiometry brings about a change in weight, or in electrical conductivity, which can be monitored experimentally as a function of time. It is not possible to give a full account of these techniques in this chapter and the reader is referred to, for example, Refs. [77-79].

10.3.2.3 The Effect of Surface Roughness and Porosity

The parameter Lc does not represent an intrinsic material property, but may depend through the value of ks on the roughness or porosity of the membrane surface. This has recently been exploited by Thorogood et al. [42] and Deng et al. [43], showing that the oxygen flux can be significantly improved if the thin dense membrane is coated on either one or both surfaces with a porous layer. Based on a simple effective medium model, and linearized transport equations, Deng et al. [43] calculated the oxygen flux through the modified membrane whose dense layer thickness is assumed to be small enough so that the drop in chemical potential across it can be neglected. The rate limiting step is thus ionic diffusion and surface exchange in the porous solid. The latter is modeled by a simple cubic array of consolidated spherical grains. In the limit of small Poagra- dients, the maximum enhancement in the oxygen flux through the symmetric membrane, over the non-coated membrane, is given by

= q L c S ( 1 - e ) / ~ + e (10.23)

where S is the pore wall surface area per unit volume, 0 the porosity and ~s the tortuosity of the solid phase in the porous structure. Maximum enhancement is achieved for a membrane whose porous layer thickness L >> Lp where

a p - q a c ( X - 0 ) / S T s ( 1 0 . 2 4 )


and refers to the active part of the porous layer. To achieve near full enhancement L - 3Lp suffices. With Lc = 100 ~tm and surface area S = 10 -6 cm -], the authors calculated an enhancement in oxygen flux, over the non-coated membrane, of almost two orders of magnitude. In a separate paper, Deng et al. [44] showed that the factor ~ is substantially reduced when the rate limiting step is the transport of gas molecules in the pores. Finally, we note the potential enhancement in oxygen fluxes, at thicknesses below Lc, by coating the dense membrane with an exchange active layer.

10.4 SOLID OXIDE ELECTROLYTES

10.4.1 Introduction

A key factor in the possible application of oxygen ion conducting ceramics is that, for use as solid electrolyte in fuel cells, batteries, oxygen pumps or sensors, their electronic transport number should be as low as possible. Given that the mobilities of electronic defects typically are a factor of 1000 larger than those of ionic defects, a band gap of at least 3 eV is required to minimize electronic contributions arising from the intrinsic generation of electrons and holes.

Useful solid oxide electrolytes to date are those with a fluorite or fluorite-related structure, especially ones based on ZrO2, ThO2, C e O 2 and Bi20 3 [80].

Mixed conduction occurs only at sufficiently low or high values of Po2, where electronic defects are generated for charge compensation of the excess of ionic defects relative to the stoichiometric composition. This latter mechanism proceeds only when the oxygen defects m vacancies or interstitials - - introduced by equilibration of the oxide with the gas phase have relatively low ionization energies and thus ionize under the selected conditions. Taking into account the mobilities of ionic and electronic defects, the range in nonstoichiometry can be correlated with the width of the electrolytic domain. For the stabilized zirconias

10-30 the domain width, at 1000~ typically extends to values below Po2 = atm [81]. On the other hand, the domain width of ceria electrolytes is limited, reflecting the ease of reduction of Ce 4+ to C e 3+ relative to that of the transition metal ions in stabilized zirconia. For example, at the same tempera ture

10-10 (CeO2)0.95--(Y203)0.05 has its domain boundary at about Po~ = atm, the value of which has been taken for tio n = 0.5 [82].

As is clear solid oxide electrolytes are not useful for applications as oxygen separation membrane, unless operated with external circuitry (oxygen pump) or as a constituent phase of a dual-phase membrane. Both modes of operation, classified in this paper as electrochemical oxygen separation, are briefly discussed in Section 10.4.3. But we first start with a discussion of the models that have been developed to describe the oxygen semi-permeability of solid oxide


electrolytes, originating from the residual electronic conductivity in the electrolytic domain. These models can be translated easily to mixed conductors. Examples are drawn from experimental studies on calcia-stabilized zirconia and erbia-stabilized bismuth oxide, clearly emphasizing the importance of both bulk diffusion and surface exchange in determining the rate of oxygen permeation through these solids.

10.4.2 Oxygen Semi-permeability of Oxide Electrolytes

10.4.2.1 Diffusion of Electronic Charge Carriers

In the absence of interactions between defects, the following reactions determine the defect concentrations in oxides with the fluorite structure,

O2,g 4- 2 V 6 ~ 2 0 ~ + 4h. (10.25)

O~ ~ O / '+ Vo (10.26)

nil ~ e' + h. (10.27)

with the respective equilibrium constants,

p4 [0~)]2 Kg = ]2 (10.28)

Po~ IV6

K F = [Oi"] [W~] (10.29)

K e =np (10.30)

Here n and p denote the concentrations, expressed as mole fractions, of electrons and electron holes, respectively. The anti-Frenkel defects Oi" and V6 are assumed to be fully ionized, which is usually observed at elevated temperatures. Whilst phases derived from 8-Bi203 show substantial disorder in the oxygen anion sublattice, reaching a maximum value at the disordering temperature, doping with aliovalent impurities is essential to achieve high ionic carrier concentrations in oxides such as ThO2, HfO2 and ZrO2 [83,84]. Doping sometimes serves to stabilize the cubic fluorite structure down to working temperatures, e.g. for HfO2 and ZrO2. In the following, we use notations D and A for dopant and acceptor cations, respectively.

The electronic conductivity, comprising p-and n-type contributions, is obtained by multiplying each of the concentrations with their respective charge and mobility. Upon substitution of the defect concentrations established by the above equilibria in the electroneutrality relation,

2[Vo] + p + [D'] = 2[Oi"] + n + [A'] (10.31)

the following expression for the partial electronic conductivity can be derived,

464 10 w DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

(~el --pFUh + nFue (10.32)

o D1/4 -b o O12/4 = (3"p ~ O 2 (~n P

where Uh, Ue are the mobilities of electron-holes and electrons, respectively, and r~, c~ ~ the corresponding partial conductivities extrapolated to unit oxygen partial pressure. In deriving Eq. (10.32), the concentration of oxygen vacancies [Vo in the electrolytic domain is taken to be fixed, either by the Frenkel equilibrium given in Eq. (10.26) or by the net acceptor dopant concentration: [A'] - [D]. Substitution of Eq. (10.32) into Wagner's equation (Eq. (10.10)) yields for the oxygen flux, after integration

RT Jo2 4F 2L [r~ ,o ,1/4 D --1/4, O , O , -1/4 , , -1/4 - t~o2 - "02 j - r~, t~o2 - Po2 )] (10.33)

noting that the ionic transference number, tion, in the electrolytic domain has been set to unity. If the experimental conditions are chosen such that the oxygen pressure at the permeate side is not too low, i.e. neglecting the n-type contribution to the electronic conductivity, Eq. (10.33) reduces to (cf. Eq. (10.18)),

rO ,1/4 O , ,1/4, (10.34) jo -T-o J

o RT/4F 2 where [3 = r~p Because of its great technological importance, much literature on the 'detri-

mental ' oxygen semipermeability flux occurring in solid electrolytes is available [85-98], of which a great number deals with the stabilized zirconias, but also includes electrolytes based on T h O 2 [97,98], H_fO 2 [87] and Bi20 3 [96] Results on the stabilized zirconias up to 1976 have been reviewed by Fouletier et al. [99]. For relatively thick electrolyte membranes the results are consistent with Eq. (10.34). But, often a value between 1/4 and 1/2 is found for the exponent [91-96]. Sometimes this has been taken as evidence for electronic trapping, i.e. a different mechanism for the incorporation of oxygen in the oxide lattice [91,92]. Dou et al. [95] were the first to invoke a surface reaction to reconcile the apparent conflict with a diffusion controlled mechanism, i.e. the overall kinetics is determined both by surface reactions and by bulk diffusion.

10.4.2.2 Modelling Equations

In this section, a model is presented for solid oxide electrolytes based upon two consecutive steps for oxygen permeation; one for the surface exchange process at the oxide surface on both sides of the membrane, and another for the joint diffusion of oxygen ions and electron-holes through the solid.


Considering oxygen exchange between the gas phase and the oxide surface via the reaction given by reaction (10.25), one may distinguish many steps, like adsorption, dissociation, surface diffusion, charge transfer and incorporation in the (near) surface layer, and reversed steps, each of these steps may impede interfacial transfer of oxygen. Following Dou et al. [95], the surface reaction may be represented by a simple two-step scheme:

O2,g ~ 2 O a d s

O ad s 4- V 6 ~ O~) 4- 2 h"

(10.35)

If it were supposed, for example, that the first of these reactions is at equilibrium and the second is rate-determining, the net flux of molecular oxygen through the interface at the high pressure side can be described by,

1/2 _ 1/2 k_ 2 [O~)] p2 (10.36) jo~ = 1/2 k2 K~/2 [V61 Po2

where k2 and k-2 are the forward and backward rate constant for the ionization and incorporation reaction (2nd step in (10.35)), K1 is the equilibrium constant for the adsorption reaction (1st step in (10.35)). The reverse equation holds for the rate of the surface reaction at the opposite side of the membrane (permeate side). As mentioned above, in the electrolytic domain the concentration of oxygen vacancies, [Vo] (and that of lattice oxygen, [O~)]) is constant. Accord- ingly, by applying the law of mass action for the overall exchange reaction given by Eq. (10.28) the actual concentration of electron holes p at either side of the membrane can be equated to a virtual oxygen partial pressure, i.e, the oxygen partial pressure if equilibrium were established with the gas phase. Consider- ing both surface reactions and bulk diffusion, wearr ive at the following set of equations for the oxygen flux

rD ,1/2 - 1/2] j o 2 = (X [zO2 - FOu(I )

1/4 jo 2 =~- [Po2q) - Po2(II) 1/4] (10.37)

jo 2 = 0~ [Po2(H) 1/2- Po2 ''1/2]

where ~ = 1/2 k2 K~ [V6] and Po~(~), PO~(H) indicate the virtual oxygen pressures at the feed and permeate side, respectively. The usual assumption of fast equilibration of the oxide surface with the imposed gas atmosphere would imply that PoS = Po~(i) and Po~" = PO~(H). Finally it may be noted that the set of equations given in Eq. (10.37) can only be solved numerically.

10.4.2.3 Examples

(a) Calcia-stabilized zirconia Dou et al. [95] studied the isothermal oxygen permeation through calcia-stabi-


I0

0

K, 8 x

% 4

Oo

_ 0 . 0 0 5 - - -

�9 I r

I I I I , , 0.1 0.2.

cm

Fig. 10.5. The effect of sample thickness on oxygen permeation through calcia-stabilized zirconia at 1230~ Solid lines are theoretical results calculated using Eq. (38). Reprinted from Dou et al. [95].

lized zirconia (CSZ) tubes at 960-1450~ and oxygen pressures Po2' of 10 -3 to 1 atm. The oxygen which permeated into the interior of tubes with known wall thicknesses was immediately pumped away by a diffusion pump and measured with a gas burette (10 -6 < Po2" < 10-4 atm). The data obtained qualitatively agreed with an oxygen pressure dependence in accordance with Eq. (10.18), in which the value of the exponent n is allowed to vary between 1/4 and 1/2. Figure 10.5 shows the measured effect of sample thickness on oxygen flux. Assuming Po2" = 0 atm, and with neglect of rate limitations at the low pressure interface, Dou et al. [95] arrived at the following expression for the oxygen flux,

I/4o2, o ] - 1 + ~2 ~02 -- I Jo2 20~L 2

(10.38)

which equation matched the experimental results well. The parameters c~ and showed similar activation energies" 191 + 5 kJ mole -1 and 206 + 11 kJ mole -1,

respectively. The quantity ~/offPo~') 'a has the unit of length. The meaning of it is more or less similar to that of the parameter Lc defined earlier in Section 10.3.2.1. At 1230~ and oxygen pressure Po2' of 1 atm, the transition from predominant control by bulk diffusion to that by surface exchange would occur at a sample thickness of about 2.7 x 10 -2 cm.

Dou et al. [95] performed their experiments on CSZ tubes with a homogeneous composition, having 10% pores by volume. The authors estimated that the


true surface area would be about 10% greater than the geometrical one used in calculation, and the surface exchange parameter a needed accordingly to be reduced for an ideal surface without pores. For the bulk transport of oxygen, the effect of non-connected micropores would be either to shortcut the solid state diffusion, in the case of fast surface exchange kinetics, or to enlarge the diffusion path to an extent which is a function of the pore size. For fully dense ceramics, the bulk diffusion parameter ~ was expected to be about 14% greater. Besides the possibility of fitting the experimental data to Eq. (10.38), the authors demonstrated that a more complicated mechanism for the surface exchange reaction may be invoked.

Interestingly, steady-state oxygen permeation measurements by Dou et al. [95] provided no evidence of a surface rate limitation for CSZ tubes containing a segregated impurity phase. The order with respect to oxygen remained close to 1/4. This second phase consisted of a mixture of metal silicates with a composition similar to that of a SiO2-CaO-A1203 eutectic. It was suggested that surface oxygen exchange on this second phase would be very rapid. In subsequent studies, the same authors used the 'time-lag' method (in which the transient process towards steady-state oxygen permeation is monitored) and a desorption technique to study chemical diffusion in CSZ with different impurities [100-102]. The observed kinetics scaled with the presence of Fe203, in such a way that faster equilibration rates were observed for samples containing a smaller impurity content. The results were taken to be consistent with a mechanism in which electron-holes are trapped at iron impurity sites.

(b) Erbia-stabilized bismuth oxide Indications of a limitation by a surface process have also been reported for

the oxygen flux through sintered dense ceramics of bismuth oxide stabilized with 25 mol% erbia, (Bi2OB)0.75--(Er203)0.25 (BE25) [96]. As is known, this material exhibits high ionic conductivity at a temperature distinctly lower than for the stabilized zirconias [103]. Using glass-sealed discs, the amount of oxygen which permeated into a closed reservoir, being flushed with helium gas prior to measurement, was monitored as a function of time. The oxygen flux was calculated from the corresponding slope, which was found invariant as long as

PO2' > > P o 2 " . In the range of temperatures (610-810~ and oxygen pressures (10 -4 - 1 atm)

covered by experiment, the concentration of minority charge carriers, i.e., electron-holes, in BE25 is proportional to P~2 with n = 1/4. However, the apparent value derived from experiment increases gradually from 1/4 to higher values upon decreasing specimen thickness from 0.285 cm to 200 gm, indicating permeation to be limited by two or more processes differing in order. The activation energy of the oxygen flux was found to increase too in the same direction. The observed behaviour can be attributed to the change-over from diffusion to


surface control upon decreasing sample thickness. The experimental data can be fitted well by means of Eq. (10.37), though it is necessary to adapt the kinetic order of the surface reaction with respect to oxygen to a value of 5/8. The parameters c~ and 13 obtained from numerical fitting appear to exhibit different activation energies; 136 + 4 kJ mole -] and 99 + 4 kJ mole -1, respectively, which indicates that the surface process is less limiting at higher temperatures. Iso- topic exchange measurements on sintered dense discs of BE25 showed a P~2 dependence with m = 0.60 at 550~ and m = 0.54 at 700~ for the overall surface oxygen exchange rate [67,104]. Figure 10.6 shows that the value for the surface oxygen exchange rate j~ x (= (~ Pc~2), normalized to air, obtained from the fit of the data agrees with that measured by isotopic exchange. The thickness, at which point the oxygen flux is half of that expected under conditions of pure diffusion-controlled kinetics, imposing opposite sides of discs to pure oxygen and hel ium gas, was calculated at 0.16 cm at 650~ and 0.09 cm at 800~ These values were found to be in good agreement with estimates of the parameter Lc as noted before in Section 10.3.2.2.

-6.50

= -7.50 r~

0

u . . _ . l

o

-8 .50

O

-9.50

opic exchange

oxygen . X,,,.~

0.90 1.00 1.10 1.20 1.30

IO00/T[K]

Fig. 10.6. Data for the surface oxygen exchange rate, normalized to air, of 25 tool% erbia-stabilized bismuth oxide (BE25) from (a) isotopic exchange and (b) oxygen permeation measurements. Reprinted

from Bouwmeester et al. [96].


10.4.3 Electrochemical Oxygen Separation

10.4.3.1 Oxygen Pump

The open-cell emf generated across an oxygen concentration cell such as

O2(Po2' ), Pt I CSZ I Pt, O2(Po2" ) (10.39)

with each side maintained at a different oxygen partial pressure Po2' and Po2" is given by,

_ RT Po2' Eeq = (1 - tel ) - ~ - I n ~ (10.40)

P o 2 "

where tel is defined as a mean electronic transference number. In the absence of any electronic conduction, i.e. when tel = 0, Eq. (10.40) simplifies to the Nernst equation. When the cell arrangement delivers a current I under load conditions, the cell voltage drops below the value Eeq, due to ohmic losses IRi (Ri =

electrolyte resistance) and polarization losses at both Pt electrodes. As an approximation,

E = Eeq - I R i - n (10.41)

where 1] represents the total cathodic and anodic polarization loss. Upon short- circuiting both Pt-electrodes, the emf of the cell drops to zero while oxygen is transported from the high pressure side PO2' t o the!ow pressure side PO2" . By applying an external power source, the applied dc voltage can be used to enhance the magnitude of the current but also to reverse its sign. That is, oxygen may be pumped in both directions; the rate of transport equals I/4F according tO Faraday's law. This is the basic principle of electrochemical oxygen separation.

An important phase during device development is optimization of the pump- ing rate, i.e. ohmic and polarization losses must be kept as low as possible. Much efforts have been concentrated on development, fabrication and testing of zirconia-based separators. For example, Clark et al. [105] has described the performance of a multi-stack yttria-stabilized zirconia (YSZ) based separator. Each cell contained a 125 ~tm thick YSZ layer of diameter 6.35 cm, whereas porous strontium-doped lanthanum manganite electrodes were used to eliminate the need for costly Pt. The largest of these separators, built with 20 cells, was found to be capable of an oxygen flux up to 1 1 min -1 at an operating temperature of 1000~ Factors influencing the efficiency of the oxygen separation process and systems analysis of conceptual oxygen production plants are also addressed.

A major drawback of ZrO2-based materials is the high temperature required for operation, typically 900-1000~ expressing the need for development of oxide electrolytes which exhibit significant levels of ionic conduction at modest temperatures. Several alternative materials may be considered. To provide a

4 7 0 10 m DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

reference point for discussion, the ionic conductivity of YSZ is about 0.1 S cm -1

at 950~ This value is found in bismuth oxide stabilized with dopants such as Er20 3 and Y203 and in cerium oxide doped with Gd203, Sm20 3 or Y20 3 already in the range 650-700~ [62,63] which electrolytes are less useful in, for example, fuel cells or sensor applications due to the presence of rather reducible ions Bi 3+ and Ce 4+ and, hence, a non-negligible contribution of electronic conduction. The suitability of Bi0.571Pb0.42801.285 as electrolyte membrane has been proposed for temperatures as low as 600~ [106]. This material suffices however from structural instabilities. Having its mechanical properties enhanced by incorporating ZrO 2 into the starting material, the optimized membrane is able to operate continuously up to 300 mA cm 2 at 600~ Fast ionic conduction at modest temperatures has also been reported in Bi4V2_yCUyOll (BICUVOX) 3 [107-109], which phases possess an intergrowth structure consisting of Bi2 O2+ blocks alternating with perovskite blocks. The material Bi2V0.gCu0.105.35 was found to exhibit an ionic conductivity of I x 10 -3 S cm -1 already at 240~ which is about two orders of magnitude higher than that of stabilized bismuth oxide [108]. In most cases the ability of these electrolytes for electrochemical oxygen separation has not yet been fully explored. Thus, it can not be excluded that relevant properties like, for example, oxygen ion conductivity, phase stability, gas tight- ness, mechanical strength and compatibility with electrode materials will not be affected during prolonged operation. Of course, the current-voltage characteristics and operational life are influenced not only by the quality of the solid electrolyte but also by the properties of the electrodes. For a recent review on oxygen electrode kinetics, see Ref. [64].

10.4.3.2 Dual-phase Composites

As seen from Table 10.1 impressive oxygen fluxes have been reported through 25 mol% yttria-stabilized bismuth oxide (BY25) [110] and 25 mol% erbia-stabilized bismuth oxide (BE25) [111,112], which oxide electrolytes were rendered electronically conductive by dispersion with silver metal. A prereq- uisite is that both constituent phases in the composite membranes do form a continuous path for both ionic and electronic conduction, having their concentrations above the critical (percolation threshold) volume fraction ~)c. The latter quantity determines the minimum volume fraction in which conduction is possible and is a function of, for example, the relative dimensions and shape of the particles of both constituent phases [113]. In actual composite materials,

3 It may be noted that BICUVOX represents only one member of a family of Bi203-based solid electrolyte phases, whichmay be derived from Bi4V2Ollby substitution of copper for vanadium. Many cations may be substituted for vanadium and the general acronym BIMEVOX was given to these materials, which have been claimed for electrochemical oxygen separation at temperatures as low as 500 K [109]. Besides copper, high oxide ion conductivity is reported for substituents titanium and niobium [212].


however, the interconnectivity between particles will not be ideal. These may be linked up to form so-called dead-ends or isolated clusters, which do not contribute at all to the conductance of the percolative system. Accordingly, conduction is expected to proceed through a significantly smaller fraction of consolidated particles or grains, which implies that the actual volume fraction of each phase should always be somewhat in excess of ~c. The optimum volume ratio is just above ~c of the high conducting phase, i.e. the metal phase, in order to have the highest effective ionic conductivity of the composite.

Dual-phase membranes made of BY25-Ag [110] and YSZ-Pd [114] behave quite similar in having their conductivity threshold at about 33-35 vol% of the metal phase. These membranes were made by conventional ceramic processing techniques. The value of ~c obtained for these composite materials agrees well with the high concentration limit predicted by simple effective medium theory in which the composite is described as a three-dimensional resistor network [115]. The effective ionic conductivity is reduced relative to that what is expected purely on the basis of the volume fraction of the ionic conducting phase, which originates, at least partly, from the enhanced tortuosity of the migrating path for the oxygen anion due to partial blocking by the metal phase. It is therefore expected that a further gain in the oxygen flux can be realized through proper design of the microstructure [112,114]. The optimum situation would correspond with the one in which the particles of each phase line up in strings (or slabs) parallel to the applied gradient in oxygen partial pressure. Even though, theoretically, the critical volume fraction of the metal phase could be reduced in this way to a value practically equal to zero, such an approach is bounded by the additional requirement for practical membranes of fast surface exchange kinetics, especially for very thin membranes.

The exchange reaction at the composite surface is confined to the three-phase boundary (tpb) between the gas, metal and electrolyte formed by particle grains being connected to the percolative network. Fast oxygen transfer can be sus- tained only if the corresponding length or area available to oxygen exchange is large enough, where it should be noted that the exchange reaction can only take place at a point remote from the tpb line which is shorter than the spill-over distance of electro-active species across the surface. The electrical field necessary to guide the current becomes distorted in the vicinity of the surface of a coarse-grained composite, where the separation between adjacent tpb lines is too large and, hence, only part of the surface is effective towards oxygen exchange. This contribution is stressed in the SOFC literature and is known as the constriction effect [116]. Often, it is the synergism between electrode and electrolyte material that leads to fast exchange characteristics. The oxygen flux through disc membranes made of BE25-Au (40 vol%) was found to increase almost one order of magnitude by substituting gold for silver in the composite [112]. This observation can be related to the higher activity of silver in the

4 7 2 10 m DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

oxygen exchange reaction on BE25, compared with gold, imposing less limitations on overall oxygen transport.

Materials like, for example, Bi2CuO44 [117], TiN [112], MgLaCrOg,s [26] have been proposed to replace the inert metals. Even though, in the examples chosen, ionic and electronic transport are confined to separate phases, mixed ionic-electronic conductors could be useful. A systematic evaluation of dual phase membranes, however, is too new so far to come to definite conclusions. Besides simple modelling in terms of a short-circuited oxygen concentration cell, to our knowledge no one has yet described oxygen permeation through dual-phase membranes, taking into account the distinct three-dimensional aspects of the microstructure that may arise in practical composite materials. Besides high values for the oxygen flux (and permselectivity), commercial use of membrane systems will demand chemical, mechanical and structural integrity of applied materials in appropriate ranges of temperature and oxygen partial pressure. Dual-phase membranes have the obvious potential to distribute specific requirements among the system components.

10.5 INTRODUCING ELECTRONIC CONDUCTION IN FLUORITE-TYPE OXYGEN ION CONDUCTORS

10.5.1 Introduction

Stimulated by the search for candidate materials for electrodes in solid oxide fuel cells (SOFC) and oxygen separation membranes, researchers have explored the possibility of introducing electronic conductivity in oxygen-ion conducting fluorite-type matrices by doping with multi-valent dopants. The major factors which establish electronic conduction in the mixed-conducting oxide solid solutions obtained are, at a given temperature, (i) the multi-valent dopant fraction, (ii) its redox characteristics and (iii) oxygen partial pressure. The suggested mechanism for electronic conduction is the hopping of electrons between adjacent dopant ions of different valence charge. However, experimental data of oxygen permeation is still scarce. In the following sections, we briefly focus on the defect chemistry, which includes some fundamentals of mass transport, and shall summarize relevant work on selected oxides.

10.5.2 Defect Chemistry

The topic of 'mixed conduction in nonstoichiometric oxides" was reviewed by Tuller [24], and his comprehensive paper is recommended to the reader inter- ested in more detail concerning the role of multivalent dopants on the defect chemistry of fluorite and fluorite-related oxides, and corresponding transport properties. Equations which express the oxygen flux in solid solutions of, e.g.,


ceria in stabilized zirconia, as a function of temperature, oxygen partial pressure and dopant concentration have been developed recently by Ling et al. [118] and Marques et al. [ 119 ].

In addition to the defect reactions given in Section 10.4.2.1, one extra reaction needs to be considered, i.e. the ionization of the multivalent cation. On using the general notation N for the multivalent cation one may write,

N M' ~--- N~ + e (10.42)

with equilibrium constant,

[N~] n KN= [NM'] (10.43)

Mass conservation requires that

[NM' ] + [N~4 ] - [NM]tota 1 (10.44)

Electroneutrality relation Eq. (10.31) must be rewritten to include the charged species NM':

2[Vo] + p + [D'] = 2[00" ] + n + [A'] + [NM' ] (10.45)

With the aid of experimentally derived equilibrium constants, Eqs. (10.42)- (10.45) may be used to construct the Kr6ger-Vink defect diagram, from which expressions for the partial conductivities of the mobile ionic and electronic defects can be derived [24].

The defect diagram obtained by Marques et al. [119] for lightly ceria-doped ZrO2-Y203, ignoring the possibility of defect association, is shown schematically in Fig. 10.7a. The corresponding conductivity diagram (Fig. 10.7b) can be obtained by multiplying each of the mobile species by their respective charge and mobility, which leads to the following expression for the total electronic conductivity,

Gel = (~p 4- (~n 4- (~h (10.46)

where (~n and o v represent the intrinsic n-type and p-type conductivities, respectively, and Oh is the extrinsic electronic conductivity owing to the multivalent cations. As distinct from the n- and p-type contributions, for which a band-like mechanism is assumed, the extrinsic contribution to electronic conduction is assumed to proceed via a small-polaron mechanism, involving the activated hopping of electrons between adjacent dopant cations of different valence charge. As the small polaron mobility includes the fraction of sites not already occupied by electrons [24], the extrinsic electronic conductivity Oh depends on both [Ce~r] and [Cezr'] and is given by

(~h = F[Cezr ' ] [Ce~r] u~ e x p ( - E H / k T ) (10.47)


(a)

A

|

E 0

,,- 0

"10

0

23

21

19

17

Vo . . . . . . , / / FM

- - - - N - . - ' - " - " = - - " - - a : : ~ - " ~ - - . - - - - . . . . .

i ",, o - / ,

�9 I ..... I I I , I

-30 -20 -10 0 10

(b)

I o g ( P o = / P a )

.-;-, E 0

6O

b

0

-4

-8

-12

- ~ ~ ~ " " " ~ " - " - , r ~ b - - , , , , , , ,, ,

, .~~ ~

/ . / ' / P1/2

!., I, ' , . I

-30 -20 -10 0

I o g ( P o = / P a )

Fig. 10.7. (a) Defect and (b) conductivity diagram for ceria-doped YSZ at 1000~ The relevant parameters to construct the diagrams are given in Ref. [119]. The theoretical dependence of ionic transference number tion and oxygen permeability jo~ are given in (c). Dashed lines in (c) refer to YSZ. FM' (Yzr') represents the aliovalent dopant used. Reproduced (slightly adapted) from Marques et al.

[119].


(c) 0 . 2 5

";', 0 . 2 0

|

E o 0 . 1 5

0

E 0 . 1 0

0 �9 -., 0 . 0 5

II [I

I I I III II

, , , , = , i ,,,, , , , i , , , , , , , , ,, , , , i , ,

I I I I

- I

i J i . . . . ,, | , .

0.8

0.6

0 . 4

0.2

i , . -

0

-20 -10 0

Iog(Po2 /Pa)

Fig. 10.7c. C a p t i o n oppos i t e .

where u ~ is the pre-exponential of the mobility and EH the hopping energy. Accordingly, r~h displays a maximum at a critical oxygen pressure, P~/~, charac- terized by equal concentrations of [Ce~.r] and [Cezr']. On assuming that the oxygen vacancy concentration remains fixed by the aliovalent dopant concentration, one may derive from Eqs. (10.42)-(10.45) that the Po2-dependence of the small polaron conductivity at given temperature takes the form,

K ~ D1/4 J-O 2

c~ h - F D1/4 + 1)2 u~ exp(-EH/kT) (10.48) (K, ~o~

-" [Cez r ] t o t a l and K~ where K~ K~ 2 21/2Ke 1/4 ,]1/2). = /(KceKg [Yzr The partial ionic and electronic conductivities may be substituted into the Wagner equation (Eq. (10.10)) to derive an expression for the oxygen flux. Typical results of such calculations are given in Fig. 10.7c, in which the oxygen flux at given temperature is plotted against Po~", assuming air to be present at the feed side of the membrane. It can be seen that the oxygen flux saturates upon lowering Po2". An inflection point occurs at Po2" = P1/2. At the lowest values of Po2", the curve bends upwards again due to the onset of the intrinsic electronic conduction. For extended discussion we refer the reader to the original papers [118,119].

10.5.3 Examples

(a) Ceria-doped ZrO2-Y203 Electrical properties of solid solutions ZrO2-CeO2-Y203 have been investi-


gated thoroughly by Cal6s and Baumard [31,32]. The main features have been confirmed by others [120,121]. The amount of 10 mol% yttria used by Cal6s and Baumard ensured a minimum concentration of oxygen vacancies in a wide range of experimental conditions. For all ceria dopant levels and temperatures (1000-1400~ ionic conduction is found to predominate at high Po2 values. Doping with ceria decreases the ionic conductivity up to (ZrO2)o.45--(CeO2)o.45- (Y203)0.1, beyond which it increases again up to the composition as high as ( C e O 2 ) 0 . 9 - - ( Y 2 0 3 ) 0 . 1 .

For not too low ceria contents the total electrical conductivity displays a maximum at reduced P02 values, shifting to lower P02 values as the temperature is decreased. It thereby follows the predictions of the preceding section, which is generally taken as evidence that the electrons in the ceria-based solid solutions move by a hopping mechanism. The maximum can be correlated with the presence of nearly equal concentrations of Ce 4+ and Ce 3+ if one takes into account the concomitant change in ionic conductivity with decreasing P02" Cal6s and Baumard deduced that, for (ZrO2)o.8r(CeO2)o.09-(Y203)0.1 at an oxygen pressure of about 10-13-10-14 atm and temperature 1200~ (~el is of the same order of magnitude as Oion and approximates 0.06 S cm d. The contribution of the electronic to the total conductivity, at a given P02 and temperature, increases with increasing ceria content, albeit at the expense of the ionic conductivity (up to the composition (ZrO2)0.45--(CeO2)0.45--(Y203)0.1).

For the most reducing conditions the conductivity becomes predominantly ionic again, albeit that the corresponding value is significantly less than that observed at high Po~. This may cause surprise knowing that the major fraction of the cerium ions under the reduced conditions adopts the trivalent state and, hence, the concentration of oxygen vacancies will be enhanced. The reduced ionic conductivity at low Po2 is attributed to enhanced defect interactions and/or lattice distortions. This type of behaviour is reminiscent to that of zirconia and ceria based electrolytes, for which it is observed that the ionic conductivity increases with the extent of aliovalent doping up to a certain limit beyond which defect ordering or formation of defect associates lowers the ionic conductivity [73]. In solid solutions with a high ceria content, for example, (ZrO2)0.45--(CeO2)0.45-(Y203)0.1, the ionic conductivity at 1100~ decreases rapidly at an oxygen partial pressure below about 10 -11 atm due to the formation of ordered pyrochlore-type domains, as confirmed by XRD measurements.

The more recent work on 10 mol% ceria-doped YSZ (5.8 mol% yttria) by Ramanarayanan et al. [122] showed that the ionic transference number tio n decreases with reduction in grain size. This observation suggests that the preferred path for electronic conduction is via the grain boundary. TEM imaging confirmed a strong tendency for cerium cations to segregate to the grain boundary, showing enrichments up to about 20 mol% compared with the value of 11 mol% observed in the lattice.


To the best of our knowledge literature reports from oxygen permeation measurements on solid solutions ZrO2--CeO2-Y203 are not available. Recent data from measurements on the related system ZrO2--CeO2-CaO are included in Table 10.1.

(b) Titania-doped ZrO2--Y203 Pure titania has the rutile structure and therefore has limited solubility in YSZ.

The observed linear decrease in lattice parameter with increasing titania concentration in these solid solutions suggests that titanium cations enter the lattice substitutionally for zirconium. Concordant with the data from XRD measurements [29,30,123] the cubic fluorite structure is retained upon addition of 12-20 mol% titania, above which a second phase appears, claimed to be ZrTiO4 [123]. The spread in data of the solubility limit produced by different authors may be due to slight differences in, e.g., yttria concentration, sample processing, sintering temperature and impurity content in the cited studies. Microstructural investigations based on SEM and TEM indicated that precipitates of the second phase actually may appear already at lower titania contents [123,125].

Contrary to the earlier observations [29,30], recent studies on electrical conductivity by Marques et al. [124] and Lindegaard et al. [125] indicate that the lattice ionic conductivity decreases with the extent of incorporation of titania into YSZ. Results confirm that the ionic conductivity of 10 mol% titania-doped YSZ in air, at a typical temperature of 1000~ is about ten times less than that of undoped YSZ. The hopping electronic conductivity at this temperature is estimated to be ---10-7 S cm -1 [121]. For similar dopant levels addition of titania appears to be more effective in enhancing the electronic conductivity of YSZ than ceria, which is not expected considering the redox behaviour of pure ceria and titania. Using thermogravimetric measurements on YSZ with ceria and titania addi- tions up to 10 mol%, Marques et al. [124] confirmed that Ce 4+ cations in these solid solutions are more easily reduced than Ti 4+. The observed increase in grain boundary conductivity with increasing titania concentration [29,30, 126] and decreasing Po~ [125] would indicate that electronic conductivity occurs at the grain boundaries. Liou and Worell [29,30] presumed segregation of Tizr to the grain boundary region, but within experimental uncertainty of EDS (Energy Dispersive Spectroscopy of X-rays) no evidence was found for titania-rich grain boundaries in the already cited study by Marques et al. [124]. The higher electronic conductivity of titania-doped YSZ was interpreted to reflect the formation of highly mobile electronic defects (large polarons) in the bulk, by comparison with the low mobility of small polarons formed in ceria-doped specimens.

At 1000~ significant levels of electronic conduction in titania-doped YSZ, as in ceria-doped specimens, are found only under strongly reducing atmospheres. Data of oxygen permeability have been presented for the ZrO2-Y203- TiO2 system by Arashi and Naito [127] (see also Table 10.1). By virtue of its

478 10 m DENSE CERAMIC M E M B R A N E S FOR OXYGEN S E P A R A T I O N

stability up to 2000 ~ it is proposed to be used as a membrane for direct hydrolysis of water to produce hydrogen [128].

(c) Miscellaneous materials Terbia has been dissolved in pure ZrO2 to form mixed-conducting solid

solutions with Tb203. 5 concentrations as high as 50 mol% by Iwahara et al. [129]. Briefly, the electrons in these mixed oxides move by hopping between Tb 3+ and Tb 4+ ions, the coexistence of which ions has been confirmed using the XANES (X-ray absorption near edge structure) technique [130]. The relative contribution of the electronic to the total conductivity measured by Iwahara et al. increases with increasing terbia concentration, in spite of the fact that the relationship between the latter quanti ty and (Itotalr at a given temperature and oxygen pressure, turns out to be very complex. At 900~ and atmospheric pressure, (~total for (ZrO2)0.7-(Tb203.5)0.3 is 1.8 x 10 -2 S cm -2, and tio n is 0.30. Oxy- gen from permeation, at 900~ was found at the argon side of a 2-3 m m thick disc membrane of this composition, at a rate of about 5 x 10 -9 mol c m -2 s -1, the value of which was measured with ambient air maintained at the feed side.

When terbia is dissolved in YSZ, this ensures a m in imum value for the oxygen vacancy concentration, which is then fixed by that of yttria. Cao et al. [131,132] examined the electrical conductivity and oxygen permeat ion of selected compositions, including (ZrO2)07-(Tb2035)03-y--(Y2OB)y with y = 0, 0.025, 0.05 and y = 0.072. At 900~ (~total decreases f r o m l . 2 x 1 0 - 2 S c m -2 for y = 0 to 0.86 • 10 -2 S cm -2 for y = 0.072, where tioniS 0.37 and 1, respectively. The oxygen flux, at 900~ passing from the air to the helium side of 2 m m thick disc-shaped membranes varied in the range 2.6-3.7 x 10 -11 mol cm -2 s -1. Hardly any effect of the yttria-content on oxygen fluxes was measured. Based upon these results, amongst some additional experimental facts, e.g., the Po2-dependence of the oxygen flux, it is concluded that the surface exchange reaction is the rate limiting step for oxygen permeation. Regrettably, no account is given as to why the data of oxygen permeation are almost two orders of magni tude lower than the one claimed for y = 0 by Iwahara et al. [129]. Dense thin films of several microns could be grown successfully on different porous ceramic substrates by electrochemical vapour deposition (ECVD) [133,134]. An oxygen permeation flux of 7 x 10 -1~ mol c m -2 s -1

at 953~ was measured for a film (ZrO2)0.86--(Tb2OB.5)0.10--(Y203)0.04 of thickness 8 ~tm deposited on a coarse a-alumina substrate, which value increased to 3 x 10 -8 mol c m -2 s -1 if the helium line was switched to CO/CO2 having Po2-- 5 x 10 -16 atm. In these experiments, air was supplied to feed side of the membrane.

Finally, we briefly describe the observations recently made in the present authors ' laboratory in an at tempt to increase the oxygen permeat ion flux through stabilized bismuth oxide by substitution of the 8-Bi203 host with 40 mol% terbium on the bismuth sites (BT40). Measurements using the concentration cell method and ac impedance confirmed that BT40 exhibits good p-type

10 n DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION 4 7 9

conductivity and is an excellent mixed conductor with ionic transference numbers, tio n = 0.74 at 650~ and t ion = 0.85 at 800~ in air [135]. Using ambient air as the source of oxygen and helium as the sweep gas on the other side of dense BT40 disc membranes, in the range of thickness 0.07-0.17 cm and temperature 600-800~ did not yield the expected increase in the oxygen flux, over BE25 [136]. Isotopic exchange measurements in the relevant range of oxygen partial pressure and temperature showed that both oxides exhibit an almost equal activity in oxygen exchange [104], which is in support of the conclusion made from oxygen permeation measurements that the oxygen fluxes through BT40, at the conditions covered by the experiments, are limited by the surface exchange kinetics.

Additional attempts have been presented to render hosts with the fluorite and the related pyrochlore structure electronically conductive by doping with mixed-valence and/or shallow dopants. The list of dopant materials examined includes oxides of elements of, for example, Ti, Cr, Mn, Fe, Zn, Fe, Sn, Ce, Pr, Gd, Tb and U. In general, however, the extent of mixed conductivity that can be obtained in fluorite-type ceramics is rather limited, by comparison with the corresponding values found in some of the perovskite and perovskite-related oxides considered in the next section.

10.6 ACCCEPTOR-DOPED PEROVSKITE A N D PEROVSKITE-RELATED OXIDES

10.6.1 Introduction

The general trend observed from the pioneering studies on oxygen permeation through perovskites of the type L n l _ x a x C O l _ y B y O 3 _ a (Ln = La, Pr, Nd, Sm, Gd; A = Sr, Ca, Ba; B= Mn, Cr, Fe, Co, Ni, Cu) by Teraoka et al. [37-39] is that higher oxygen fluxes are facilitated by increased A-site substitution, and a lower thermodynamic stability of the particular perovskite.

Clearly, not all these perovskite compositions are useful for oxygen delivery applications. For example, ceramics based on Lal_xAxCrOB-a (x = Sr, Ba, Ca), Cal_xSrxCrl_yMnyO3_ a and Cal_xCaxCrl_yCoyO3_ a have been proposed for use as interconnection material (separator) in solid oxide fuel cells (SOFC), and therefore should be dense and impermeable in order to prevent burning off of the fuel without generating electricity [137,138].

Selected perovskite compositions are also targeted in basic SOFC research for use as potential electrode material for the cathodic reduction of oxygen. The most promising cathode materials to date are the manganites Lal_xSrxMnO34 [137,138]. The composition with x = 0.15 scarcely permeates oxygen up to 900~ as was measured by feeding air and helium to opposite sides of a dense sintered membrane of I mm thickness [136]. The observed behaviour is consis-

480 1 0 - DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

tent with the low value of the oxygen self-diffusivity in La0.sSr0.5MnO3~, determined by 180-160 isotopic exchange, and can be attributed to the small negative departure from oxygen stoichiometry exhibited in the range of temperature and oxygen pressure covered by experiment [139]. On the other hand, oxygen transport is actually predicted to be quite fast under conditions of high oxygen- deficiency, i.e. low oxygen partial pressures, as the oxygen vacancy diffusion coefficient of La1_xSrxMnO34 was found to be comparable in magnitude with that of Fe- and Co-based perovskites [140].

Emerging from the first of these studies by Teraoka et al. [37] is that in the series La1_xSrxCo1_yFeyO34 the oxygen fluxes increase with Co and Sr content, the highest flux being found for SrCo0.sFe0.2034. Data were obtained with air on one side of a 1 mm thick disc specimen, using helium as sweeping gas on the other side, up to a maximum temperature of 1150 K. The observed oxygen fluxes were found to be roughly proportional to the ionic conductivity of the perovskites, which is in agreement with the fact that the electronic conductivity of compositions in this series can be extremely high, typically in the range 102103 S cm -1 [40]. Four-probe dc measurements using electron blocking electrodes showed that the ionic conductivity at 800~ in air can be 1-2 orders of magnitude higher than that of stabilized zirconia [40]. These findings have been confirmed by others, apart from scatter in the published data, which partly reflects the experimental difficulties in measuring the ionic conductivity in these predominantly electronic conductors [141-144].

In a subsequent study, Teraoka et al. [38] investigated the influence of A and B site substitution on oxygen permeation through La0.6A0.4Co0.sFe0.2034 (A = La, Na, Ca, Sr, Ba) and La0.6Sr0.4Co0.8B0.203,s (B = Cr, Mn, Fe, Co, Ni, Cu). As seen from Figs. 10.8 and 10.9, the oxygen permeability in the two series increases in the respective orders La < Na < Sr < Ca < Ba and Mn < Cr < Fe < Co < Ni < Cu, which differ from trends in the periodical system, as far as comparison is meaningful. Results from ionic and electronic conductivity measurements of La0.6A0.4Co0.8Fe0.203_ 8 (A = La, Ca, Sr) and La0.6Sr0.4Co0.8B0.203.8 (B = Fe, Co, Ni, Cu) suggest that oxygen permeation is governed by the ionic conductivity [39]. In the homologous series Ln0.6Sr0.4CoO3_a, the oxygen flux was found to increase in the order La 3+ < Pr 3+ < Nd 3+ < Sm 3+ < Gd 3+ which corresponds with a decrease in radius of the lanthanide-ion [38].

Since the initial observations by Teraoka et al., a considerable number of studies have appeared. Selected perovskite compositions have been re-examined, while a few others have been adapted in an attempt to optimize the oxygen fluxes. The list of materials for which oxygen permeation data are presently available has been extended to include, LaCoO34 [145], La1_xSrxCoO34 [146-149], La1_xSrxFeO34 [150,151], Lal_xAxCol_yFeyO34 (A = Sr, Ca) [12,139,144, 152], SrCo0.8Fe0.2034 [13,41,153,154], SrCoo.sBo.203_8 (B = Cr, Co, Cu) [154], SrCol_• (B = Cr, Mn, Fe, Ni, Cu, x = 0...0.5) [155], and Y1_xBaxCoO34 [156]. In general, fair agreement


1 . 5 0

A 1 . 0 0 ,-;.,

E o

o E

,,, 0.50 O

�9 Ba

�9 Ca

Sr

�9 Na

�9 La

0.00 300 500 700 900

temperature (~ Fig. 10.8. Temperature variation of the oxygen permeation rate from the air to the helium (30 cm 3

min -1) side of disc membranes La0.6A0.4Co0.8Fe0.2034 (A = Na, Ba, Ca, Sr), 20 m m in diameter and 1.5 m m thick, after Teraoka et al. [38]. Reproduced (data re-scaled) from Teraoka et al. [38].

is obtained with data produced by Teraoka et al., albeit that in a number of studies the observed oxygen fluxes are reportedly found to be significantly lower [12,153,154].

The pioneering studies by Teraoka et al. [37-39] have opened a very challenging research area as the perovskites, e.g. Lal_~SrxCo1_yFeyO34, have a bright future for use as oxygen separation membrane. The precise composition may be tailored for a specific application, but this has not yet been fully developed. One of the important issues is considered to be the low structural and chemical stability of the perovskites, especially in reducing environments, which remains to be solved before industrial applications become feasible. In order to meet this challenge, it is necessary first to understand the factors that limit and control the quality criteria for any given application. The perovskite and related oxides exhibit a great diversity of properties, like electrical, optical, magnetic, catalytic properties, which have been studied extensively. In the following sections, we mainly focus on those properties affecting the magnitude of the oxygen fluxes through these materials.

482 10 n DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

1 . 5 0

~ " 1 . 0 0

.'~~ 0.50

0.00 500

�9 Cu �9

�9 Ni

�9 Co

�9 Fe �9 Cr + Mn

~ ~ , I , I ,

600 700 800 900

t e m p e r a t u r e (~ C)

Fig. 10.9. Temperature variation of the oxygen permeation rate of La0.6Sr0.4Co0.8B0.203-8 (B = Cr, Mn, Fe, Co, Ni, Cu) after Teraoka et al. Experimental conditions are specified in the legend of Fig.

10.8. Reproduced (data re-scaled) from Teraoka et al. [38].

10.6.2 Structure and Defect Chemistry

10.6.2.1 Perovskite Structure

The ideal perovskite structure ABO3 consists of a cubic array of corner-shar- ing BO 6 octahedra, where B is a transition metal cation (Fig. 10.10). The A-site ion, interstitial between the BO 6 octahedra, may be occupied by either an alkali, an alkaline earth or a rare earth ion. In many cases the BO 6 octahedra are distorted, or tilted, due to the presence of the A cation, which is generally larger in size than the B cation. Alternatively, the perovskite structure may be regarded as a cubic close-packing of layers AO 3 with B cations placed in the interlayer octahedral interstices [157]. The latter turns out to be more useful in distinguishing different structural arrangements (stacking sequences) of perovskite blocks. The tolerance limits of the cationic radii in the A and B sites are defined by the Goldschmidt factor, which is based on geometric considerations: t = (ra + ro) / (~-(rB + ro)), where rA, rB and ro are the radii of the respective ions [158]. When the distortion becomes too large, other crystal symmetries such as

~ A @0 QB

::::iii i:i::..


Fig. 10.10. Ideal perovskite structure.

orthorhombic and rhombohedral appear. Nominally, the perovskite structure should be stable between 1.0 < t < 0.75. The ideal perovskite lattice exists only for tolerance factors t very close to one.

Clearly, it is the stability of the perovskite structure that allows for large departures from ideal stoichiometry, resulting either from the substitution with aliovalent cations on the A or B-site or from redox processes associated with the presence of transition metal atoms which can adopt different formal oxidation states. Oxygen vacancies are free to move among energetically equivalent crystallographic sites as long as the perovskite structure exhibits ideal cubic symmetry. The degeneracy between sites disappears upondis tor t ion of the lattice towards lower symmetries. The onset of electronic conductivity mainly depends on the nature of the B-site cation. The total electrical conductivity can be either predominantly ionic as in the acceptor-doped rare earth aluminates or predominantly electronic as in the late transition metal containing perovskites considered below.

10.6.2.2 Nonstoichiometry

Important contributions to the area of defect chemistry of the acceptor-doped Lnl_xAxBO3, perovskites, where B is selected from Cr, Mn, Fe or Co, have been made by a number of investigators. Particular reference is made to reviews provided by Anderson [159,160] and Mizusaki [161]. The substitution of divalent alkaline-earth ions on the A-site increases the concentration of oxygen vacancies. Temperature and oxygen partial pressure determine whether charge compensation occurs by an increased valency of the transition metal ion at the B-site or by the formation of ionized oxygen vacancies. Thermogravimetric studies have indicated that in, for example, LaCrO3, YCrO3 and LaMnO3 the


native nonstoichiometric ionic defects are cation vacancies, leading to oxygen- excess stoichiometries [160]. For simplicity, it is assumed here that extrinsic ionic defects generated by A-site substitution prevail, i.e. only oxygen-deficient stoichiometries are considered. Furthermore, crystallographic sites available for oxygen are taken to be energetically equivalent.

For the purpose of our discussion, LaFeO3 is considered to be the host for substitution. The dissolution of SrFeO 3 into this material can be represented by,

LaFeO 3 SrFeO3 ) SrLa' + FeFe + 30~) (10.49)

The incorporation of Sr 2§ thus leads to charge compensation by the formation of Fe 4+ ions, which is in accord with the Verwey principle of controlled ionic valency [162]. The extent of oxygen non-stoichiometry is established by the following defect chemical reactions,

2FeFe + O~) ~ 2Fete + V' o + 1/2 0 2 (10.50)

2Fete ~ FeFe' + FeFe (10.51)

with the corresponding equilibrium constants,

1/2 [Fete] 2 [V'o] Po 2 Kg = [FeFe] 2 [O~)] (10.52)

[FeFe'] [FeF e] K d = [Fete] 2 (10.53)

The oxygen vacancies formed at elevated temperatures and low oxygen partial pressure are assumed to be doubly ionized. The thermally activated charge disproportionation reaction given by Eq. (10.51) reflects the localized nature of electronic species and may be treated as equivalent to the genera.tion of electrons and electron holes by ionization across a pseudo band gap (cf. Eq. (10.27)). The associated free enthalpy of reaction may be taken equal to the effective band gap energy.

At fixed A/B site ratio the following condition must be fulfilled

[FeFe'] + [Fete] + [FeFe ] = 1 (10.54)

and the condition of charge neutrality is,

[SrLa'] + [FeFe'] = 211/'O] + [FeFe] (10.55)

In the absence of extended defects, i.e. no interaction between point defects, Eqs. (10.52)-(10.55) may be used with the aid of experimentally determined equilibrium constants to construct the Kr6ger-Vink defect diagram, from which ex-


pressions for the partial conductivities of the mobile ionic and electronic defects can be derived.

Oxygen nonstoichiometry of the perovskites Lal_xSrxBO3.s (B = Cr, Mn, Co, Fe) and its relationship with electrical properties and oxygen diffusion has been studied extensively [159-161]. Typical nonstoichiometry data for La1_xSrxFeO34 and for some other perovskites as obtained from gravimetric analysis and coulometric titration are given in Fig. 10.11. At small oxygen deficiency, acceptor dopants are the majority defects. The charge neutrality condition then becomes,

[SrLa'] = [FeFe] (10.56)

In this region, one finds for the oxygen non-stoichiometry 5,

~/2 (10.57) oc p o 2

noting that 8 = [V6], by definition. A plateau is observed around the point of electronic stoichiometry, 8 = x/2, where the charge neutrality condition reads,

[Srca'] = 2[V6] (10.58)

3 . 0 5

2.95

o 3

2.9

2.85

&

&

A . . . . . -

/ 0 Lao.9Sro.lCo03.6

�9 Lao.oSro.lFeO3.s

o Lao.3Sro.7Cr03-8

zx Lao.2Sro.sMn03-8

.8 lu ~ I I I I I , I , I I I , . I I I I I I I I '

-20 -16 -12 -8 -4 0

Iog(P 0 z/atm)

Fig. 10.11. Data of oxygen nonstoichiometry of Lao.75Sro.25CrO3-~, Lao.gSro.lFeO3-~, Lao.gSro.lCoO3-~, and Lao.sSro.2MnO3-~ at 1000~ as a function of oxygen partial pressure. Solid lines are results from a fit of the random point defect model to the experimental data. Reproduced (slightly adapted) from

Van Hassel et al. [185].


corresponding with a minimum in the electronic conductivity of La1_xSrxFeO34 [163,164]. In this region, the oxygen non-stoichiometry is virtually constant. As the oxygen activity decreases further, oxygen vacancies are again generated, down to the oxygen activity at which decomposition of the perovskite structure occurs. The onset of the different regions depends on the nature of the transition metal B-cation. The incentive of B-site substitution can therefore be to optimize oxygen transport in appropriate ranges of oxygen partial pressure and temperature. As discussed below, doping may also increase stability or suppress cooperative ordering of oxygen vacancies.

10.6.2.3 Localized versus Delocalized Electrons

Given the relative success of the above point defect scheme to model the experimental data of oxygen nonstoichiometry and electrical conductivity for Lal_vSrxFeO34 [165,166] and Lal_xSrxCrO3.s [167], its use is less satisfactory for Lal_xSrxCoO34 and Lal_xSrxMnO3~, which compounds show notably high values for the electronic conductivity.

Nonstoichiometry of the compounds La1_xSrxCoO34 (x = 0, 0.1, 0.2, 0.3, 0.5 and 0.7) in the range 10 -5 < Po2 < 1 atm and 300< T < 1000~ was investigated by Mizusaki et al. [168] using thermogravimetric methods. At 800~ 5 in Lal_xSrxCoO3_ s varies almost proportional to Po2 n with n = - 1 / 2 for x = 0 to n = - 1 / 1 6 for x = 0.7 (see Fig. 10.12). No plateau is observed around 8 = x/2. Fitting the 5-Po2 relationship in accord with the random point defect model leads to very large concentrations of disproportionation reaction products Coco" and Coco. A corollary is that the pseudo bandgap must be very small. The model fit, however, is less satisfactory for high Sr substitutions [169]. A similar exp lana t ion holds for La1_xSrxMnO34, d i s regard ing the oxygen-excess stoichiometries seen in this system at high oxygen partial pressures.

At high oxygen deficiency of the perovskite, the validity of the ideal mass action equations (based upon dilute solution thermodynamics) cannot be assumed a priori. In addition, interaction and association between defects are expected at high defect concentrations. A further limitation concerns the nature of electronic defects. The general assumption, that in the first row transition metal perovskites changes in the oxygen content leads to changes in the 3d electronic configuration, may be too naive. It is based implicitly on the idea that oxygen is strongly electronegative and, by comparison, the 3d electrons can be easily ionized. There is substantial evidence from soft-X-ray absorption spectroscopy (XAS) based studies that the electron holes introduced by doping with divalent earth-alkaline ions go to states with significant O 2p character [170]. This has also been reported for the perovskite-related oxide YBa2Cu306§ [171]. In a localized description, i.e. assuming a narrow bandwidth of the hole band derived from the O 2p band, this would imply that 0 2- is effectively converted into O-.


0 ,-4

- 1

- 2

- 3

I i ' ' I - " ~ ' I "' - i ....

m

.3

a

0

m i

.... I 1 1 1,, I l _

-4 -2 0

log (P /atm) 0 2

Fig. 10.12. Oxygen pressure dependence of 6 in Lal-xSrxCoO3-6 for different strontium contents

at 800~ Reprinted from Mizusaki et al. [168].

A proper description of electronic defects in terms of simple point defect chemistry is even more complicated as the d electrons of the transition metals and their compounds are intermediate between localized and delocalized behaviour. Recent analysis of the redox thermodynamics of La0.sSr0.2CoO34 based upon data from coulometric titration measurements supports itinerant behaviour of the electronic charge carriers in this compound [172]. The analysis was based on the partial molar enthalpy and entropy of the oxygen incorporation reaction, which can be evaluated from changes in emf with temperature at different oxygen (non-)stoichiometries. The experimental value of the partial molar entropy (free formation entropy) of oxygen incorporation, Aso2, could be


fitted by assuming a statistical distribution among sites on the oxygen sublatfice,

(3 -8 ) Aso2 = s ~ - 2k In 8 (10.59)

where s o is a constant. That is, no entropy change associated with electron annihilation can be identified. The partial molar enthalpy (free enthalpy of formation of vacancies) associated with oxygen incorporation was found to decrease almost linearly with 8. A first inclination might be to assume that the mutual repulsion between oxygen vacancies increases with increasing oxygen deficiency. But this interpretation immediately raises the question why such a behaviour is not found in the case of La1_xSrxFeO34 [166,167]. Instead, the experimental data are interpreted to reflect the energetic costs of band filling. With increasing oxygen nonstoichiometry in La0.sSr0.2CoO34 the two electrons, which are needed for charge compensation of a single oxygen vacancy, are donated to an electron band broad enough to induce Fermi condensation characteristic of a metallic compound. The average density of electron states at the Fermi-level is determined to be 1.9_+0.1 eV -1 per unit cell. The physical significance of the work is that the defect chemistry of La0.8Sr0.2CoO3-4 cannot be modeled using simple mass action type of equations. An empirical model for the oxygen non-stoichiometry of La0.sSr0.2CoO3-8 is proposed, which demon- strates that the density of states is related to the slope of the log-log plots of 8 versus Po~. In support of these interpretations, it is noted that XAS has not been successful in detecting charge disproportionation in LaCoO34, due to localiza- tion of electrons, in the temperature range 80-630 K [173]. The nonstoichiometry data obtained for La0.8Sr0.2CoO3-4 are found to be in good agreement with earlier results from gravimetric analysis in the series La1.~rxCoO34 obtained by Mizusaki et al. [168], which authors arrived at more or less similar conclusions regarding the role of electronic states in the energetics of oxygen incorporation into these compounds.

10.6.3 Oxygen Desorption and Perovskite Stability

As seen from Fig. 10.11, the value of (3-8) in Lal_xSrxCoO34 falls off with decreasing oxygen activity much more rapidly than for the other compounds shown. The general trend at which the perovskites become nonstoichiometric follows that of the relative redox stability of the late transition metal ions occupying the B-site, i.e. C r 3+ > Fe 3+ > Mn 3+ > C o 3+. The reductive nonstoichiometry of the cobaltites increases further by partial B-site substitution with copper and nickel.

The reductive (and oxidative) nonstoichiometry and the stability in reducing oxygen atmospheres of perovskite-type oxides was reviewed by Tejuca et al. [174]. Data from temperature programmed reduction (TPR) measurements indicate that

10 E DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION 489

the stability (or reducibility) of the perovskite oxides increases (decreases) with increasing size of the A ion, which would be consistent with the preferred occupancy of the larger Ln 3+ ion in a 12-fold coordination. The trend is just the reverse of that of the stability of the corresponding binary oxides. The ease of reduction increases by partial substitution of the A ion, e.g., La 3+ by Sr 2+. Trends in the thermodynamic stabilities of perovskite oxides have been systematized in terms of the stabilization energy from their constituent binary oxides and the valence stability of the transition metal ions by Yokokawa et al. [175].

The stability of the undoped perovskites LaBO3_a, at 1000~ expressed in terms of Po2 decreases in the order LaCrO34 (10 -20 atm) > LaFeO34 (10 -17 atm) > LaMnO3_~ (10 -15 atm) > LaCoO34 (10 -7 atm), noting that the cited value for LaCrO34 corresponds with the lowest limit in a thermogravimetric study by Nakamura et al. [176]. The same trend was found by means of TPR [174].

Tabata et al. [177] and Seyama [178] both described significant differences in the chemical composition of the surface, due to Sr segregation, compared with the bulk composition in a series of powders Lal_~SrxCoO3_~. This indicates a behaviour of the surface different from that of the bulk in these compounds. Not only can this account for a number of observations made in the total oxidation of CO and C H 4, as discussed by the authors, but it is also considered to be an important factor when one tries to correlate the composition of a perovskite with its activity in surface oxygen exchange.

The sorpfion kinetics of oxides is certainly influenced by their corresponding defect structure. A number of interesting observations were made by Yamazoe and co-workers [179,180], showing that for perovskites LaMO3_~ (M = Cr, Mn, Fe, Co, Ni), Lal_vqrxCoO3_~ (x = 0, 0.2, 0.4 and 1) andLa0.8Ao.2CoO3_8 (A = Na, Ca, Sr and Ba), two distinct types of oxygen are desorbed upon heating in a helium stream after a pre-treatment step in which the oxide was saturated in an oxygen-rich atmosphere at high temperature, followed by slow cooling to room temperature. The oxygen desorbed in a wide range at moderate temperatures, referred to as c~-oxygen, was found to be correlated with the amount of partial substitution of the A ion. The onset temperature of the so-called ~-desorpfion peak observed at high temperature was correlated with the thermal decomposition temperature of the corresponding transition metal oxides. Accordingly, the ~-peak corresponds with the reduction of the transition metal ion from B 3+

to B 2+. The partial substitution of Co by Fe in the series Lal_xSrxCOl_yFeyO3_ ~ stabilizes the Co B+ oxidation state (no -peak observed), while shifting the o~-type of desorpfion to lower temperatures [181,182].

10.6.4 Equations for Oxygen Transport

Equations for oxygen transport can be derived from the point defect equilibria discussed in Section 10.6.2.2. This provides us with some general insight


into the transport behaviour of oxygen-deficient perovskites. Strictly speaking, the equations presented below are valid at low defect concentrations only, i.e. assuming oxygen defects to be randomly distributed.

Oxygen transport in the perovskites is generally considered to occur via a vacancy transport mechanism. On the assumption that the oxygen vacancies are fully ionized and all contribute to transport, i.e., oxygen defects are not associated, the Nernst-Einstein equation reads,

4 F 2 [ V o ] Dv (10.60) Gi~ ----- RTVm

where Dv is the vacancy diffusion coefficient and Vm is the perovskite molar volume. Since electronic conduction in the perovskites predominates, i.e. Gel >

Glow, the integral in the Wagner equation (Eq. (10.10)) involves only Gion over the applied oxygen partial pressure gradient. Using Eq. (10.60), we may rewrite the Wagner equation, to give

In P"o2

Dv (10.61) j o 2 = 4VmL ~ ~) d In Po2

In P'o2

by virtue of 6 = [Vo]. Evaluation can be performed numerically provided that Dv and the 5-1n(Po2) relationship are known. The ability of Eq. (10.61) to quantitively fit experimental data of oxygen permeation is illustrated for La0.9Sr0.1FeO3~ in Fig. 10.13. Similar results have been presented for, e.g., La0.75Sr0.25CrO3~ [183] and La0.70Ca0.30CrO3~ [184]. The analytical solution of the integral given by Eq. (10.61) incorporating random point defect chemistry has been given by Van Hassel et al. [185].

When data of oxygen nonstoichiometry follows a simple power law 8 ~ P" 02 , integration of Eq. (10.61) yields an expression similar to that of Eq. (10.18) having ~ = DvS~ n. Examination of the data from oxygen permeability measurements on disc specimens of thickness 2 mm in a series Lal_xSrxCoO3_ 8 (0_ x _< 0.8) in a study by Van Doom et al. [148] indicate that the results, at 1000~ can be fitted well by this equation, the validity of which is usually restricted to a small range in oxygen partial pressure. For compositions x _< 0.6, the values of n obtained from fitting were found to be in excellent agreement with the corresponding slopes of the In 5--ln Po2 plots derived from data of thermogravimetry in the range 10-4_< Po2 -< 1 atm. For compositions x = 0.7 and x = 0.8 the agreement obtained is less. Measurements made as a function of specimen thickness, down to a minimum value of 0.5 mm, suggest that this may be due to a partial rate control of the oxygen fluxes for these compositions by the surface reaction.

In writing Eq. (10.61), Dv was taken to be constant. Strictly speaking, Dv may decrease slightly when the oxygen-deficiency increases (i.e., towards decreas-


0.3

Ishigaid et al.

E o E E

O4

O "~ Dv~ =6.0.1 -e

-2.5 -2 -1.5 -1 -0.5 0

log (po2/bar) Fig. 10.13. Theore t ica l fit of feed-s ide Po2-dependence of oxygen pe rmea t ion th rough La0.9Sr0.1FeOa~, at 1000~ The best fit is obtahled when Dv equals 6 x 10 .4 cm 2 s -1, which slightly deviates from the corresponding value obtahled from isotopic exchange. Reprinted from Ten Elshof

et al. [1511.

ing oxygen partial pressure), to an extent depending on the particular solid. In accord with classical diffusion theory, the probability of vacancy hopping, hence, Dv is proportional to a factor (1-5/3), which represents the site occupancy of lattice oxygen anions [186]. Moreover, complications due to local stresses resulting from a change in cell volume with decreasing oxygen partial pressure may need further consideration, especially if the nonstoichiometry becomes relatively large.

With the help of Eq. (10.60) the vacancy diffusion coefficient Dv can be determined from ionic conductivity measurements provided that data of oxygen nonstoichiometry are available. The direct measurement of (~ion in these materials requires the use of auxiliary electrolytes such as doped zirconia or ceria to block the electronic charge carriers. The ionic current that is passed through the sample is measured by the electric current in the external circuit. The problems faced are that of interfacial charge transfer between the blocking electrodes and the mixed conductor and that it is very difficult to effectively block all of the electronic current. Short-circuiting paths for oxygen transport can occur such as diffusion along the oxide surface or via the gas phase through rapid exchange, leading to overestimates of the ionic conductivity [143]. Also, there is the possibility of an interfacial reaction between the perovskite and the blocking electrode material or the glass used for sealing to suppress the para- sitic contributions to oxygen transport [187].

The relation between Dv and the tracer diffusion coefficient D* can be expressed as

8 D* =f (3 L 8i Dv (10.62)

492 10 m DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

where f is the correlation factor for diffusion of oxygen vacancies across the ideal perovskite anion sublattice: f = 0.69 (for small values of 8) [75]. Data of Dv thus obtained have been published by Ishigaki, Mizusaki and coworkers for single crystals of LaCoO34 [74,188] and LaFeO34 [189]. The value of D* at 950~ in both oxides is found to be proportional to P~2, with n =-0.34 + 0.04 and

-V2 n = -0.58 + 0.15, respectively, where a Po2-dependence is expected in the high Po2-regions covered by experiment. These results provide firm evidence that diffusion in the perovskites occurs by a vacancy mechanism. Though the values of D* observed in these oxides differ about 3 orders of magnitude, the corresponding values of Dv are nearly the same. Additional data have been reported for single crystals of La0.gSr0.1CoO34 (x = 0.1) [75], La1_xSrxFeO34 (x = 0.1, 0.25 and 0.4) [75], and polycrystaIIine phases La1_xSrxFeO34 (x = 0, 0.4, 0.6 and 1.0) [190] and La0.70Ca0.30CrO3.8 [184], revealing that the diffusivities at elevated temperatures may be similar to those observed for fluorite and fluorite-related oxides, albeit that the associated activation energy generally tends to be slightly higher in the perovskite structure (see Ref. [161]).

When electronic conduction predominates, one may derive the following relationship between Dv and the chernical diffusion coefficient D (by combining Eqs. (10.12) and (10.60)),

N D~ O In Po~ (10.63) D =--~-- 31n8

Measurements of the weight change following a sudden change of the oxygen activity in the gas phase were carried out for determining D in La1_~SrxCoO34 (x = 0 and x = 0.1) at 800-1000~ in the range 10 -5 < Po2 < 1 atm [191]. The calculated values of D v agree with those obtained from the above studies, which suggests that the random point defect model holds well for the cobaltites at low Sr contents. Fair agreement is also obtained with results from relaxation experiments on La1_xSrxCoO34 (x = 0 and 0.2) in which the time change of the conductivity was traced [192]. A typical value of D taken from these studies is 10 -s c m 2 s -1 at 900~ Relaxation experiments were also carried out to study chemical diffusion in, for example, SrCo1_yFeyO34 (y = 0.2, 0.5 and 0.8) [187,193], Lal_xSrxCOl_yFeyO3_ ~ (0.2 < x < 1.0 and 0 < y < 1.0) [194], Lal_xCax_ CrO3_ ~ (x = 0.1, 0.2 and 0.3) [195], Lal_~SrxMnO3,s (x = 0.05--0.20) [140] and Lal_ xSrxMnO34 (x = 0.20 and 0.5) [196,197].

10.6.5 Electronic Conductivity

The late transition metal-containing perovskites exhibit high electronic conductivities. In the materials which receive prime interest for oxygen delivery applications, the electronic contribution at high temperature of operation is usually predominant. The values for e.g., La1_~SrxCol_yFeyO34 at 800~ in air


range between 102-103 S cm -1, whilst 10-2-1 S cm -1 is found for the ionic conductivity [40]. The ionic transference numbers in this series vary between 10-4-10 -2. Departures of the above behaviour can occur at reduced oxygen partial pressures, due to the loss of p-type charge carriers [194,198]. This is most serious at the point where the hopping-type of conductivity goes through a minimum, to become n-type, e.g., in La1_xAxFeO34 (A = Sr, Ca) [163,164,199], provided that the correspondingly low oxygen pressure is maintained during experiment. At 1000~ the minimum in La0.75Sr0.25FeO3-s occurs below Po2 values of 10 -12 atm [164]. In the following, we discuss a number of characteristics which are considered to control electronic conductivity in this class of oxides and give some examples of their behaviour.

Electronic conduction in usual ranges of temperature and oxygen pressure is reported to be p-type, and is commonly explained by assuming a small polaron mechanism with a thermally activated mobility [159,160]. This behaviour may be masked by substantial oxygen loss, and concomitant decrease in the concentration of p-type charge carriers, seen at the highest temperatures and at reduced oxygen partial pressures. The direct overlap between transition-metal d orbitals is known to be small, being across a cube face. The hopping transport of mobile charge carriers between two neighbouring B-cations in the perovskite lattice is mediated by the O 2p orbital,

B n+ - 0 2 - - B (n - l )+ --4 B (n - l )+ - O - - B (n - l )+ --4 B (n - l )+ - 0 2 - - B n+

which is known as double exchange, first discussed qualitatively by Zener [200]. This process is favoured by a strong overlap of empty or partly filled cation orbitals of the d-manifold (involving e~ and t2~ orbitals) with the filled 0 2p orbital of neighbouring anions, and reaches a maximum for a B-O-B angle of 180~ corresponding with ideal cubic symmetry.

Lal_xAxCrO3_~ (Sr, Ca) and Lal_xAxFeO3_ s (Sr, Ca) are typical examples of which the data of small-polaron transport can be explained in terms of simple defect chemistry, including the thermally activated charge disproportionation among the B-cations. The predominant mechanism of hopping in Lal_xAxFeO3_s (Sr, Ca) in the p-type region is between Fe 4+ and Fe 3+ valence states, changing to that between Fe 2+ and Fe 3+ in the n-type region, upon lowering the oxygen partial pressure. The electrical conductivity thereby passes through a minimum at the point of electronic stoichiometry, where the concentrations of Fe 4+ (FeFe') and Fe 2+ (FeEd)are equal [163,164,199]. Charge disproportionation has also been used to account for results from electrical conductivity and thermopower measurements of selected substitutionally mixed oxides LaMnl_xCOl_xO3-s [201], LaMnl_xCrl_xO3_ ~ [202], Lal_xCaxCOl_yCryO3_~ [203] and Lal_xSrxCOl_yFeyO3_s [204]. Preferential electronic charge compensation may occur in these compounds, i.e. the charge carrier may (at low temperature) be temporarily trapped at the small polaron site which is lower in energy, thereby decreasing the electrical conduc-

4 9 4 10 ~ DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

tivity. This effect disappears at high temperature, when the thermal energy is sufficient to surpass the barrier between the traps and the more conductive hopping sites, or when the population of low energy sites exceeds the percolation limit. In the latter case the electrical conductivity is controlled by the short-range hopping among the lower energy sites. Based upon results from electrical conductivity and thermopower measurements, in conjunction with some other techniques, Tai et al. [204] concluded that in La0.sSr0.2COl_yFeyO3. 8 compensation of Fe 3+ --~ Fe 4+ is more likely than Co 3+ --~Co 4+.

The electrical conductivity and the spin-transition state of LaCoO34 and La1_xSrxCoO34 has been studied extensively. The covalent mixing of orbitals induces itinerant behaviour of the charge carriers in La1_~SrxCoO34 [205]. Racah and Goodenough [206] claimed a first order localized-electron to collective- electron phase transition in LaCoO3 at 937~ in air, though the electrical conductivity was found to be continuous at the transition. Electrical conductivity and differential thermal analysis (DTA) behaviour of ACoO3. 6 (A = Nd, Gd, Ho, Y, La) were investigated by Thornton et al. [207]. Evidence was adduced that the transition previously noted in LaCoO3 (and in the other cobaltates) may be caused by the presence of binary cobalt oxides. The endothermic heat effect observed in DTA can be correlated with a similar feature obtained from a sample Co30 4. Instead, a gradual semiconductor-to-metal transition with increasing temperature is suggested in compounds ACoO3,s. For LaCoO3& which adopts a rhombohedrally distorted perovskite structure at room temperature, this range was found to extend from 110 to 300~ Analysis was based in part upon the temperature independent magnetic susceptibility observed at high temperature, which was attributed to Pauli-paramagnetism (with possible Van Vleck type of contributions) of the conduction electrons. Data produced by Mizusaki et al. [208] suggest a close relationship between the transition temperature of the conductivity and the rhombohedral angle, which gradually decreases with increasing x in Lal_xSrxCoO3_~. The conduction becomes metallic when the rhombohedral angle becomes smaller than 60.3 ~ noting that the value of 60 ~ corresponds with ideal cubic symmetry. Another factor is the extent of oxygen non-stoichiometry. In the metallic region, the electrical conductivity of Lal_xSrxCoO3. 5 decreases almost linearly with increasing 8. Apart from changes in d-band occupancy, it is assumed by the authors that band narrowing takes place with increasing oxygen nonstoichiometry.

Finally, the data of electrical conductivity and thermopower of Lal_~Sr• at elevated temperature has been treated by a hopping mechanism for x < 0.2, and by a band model for the semi-metallic behaviour observed at x > 0.3 by Mizusaki [161]. On the other hand, to explain their results of Lal_xSrxMnO3_~ and Lal_xSrxMnO34 for compositions with 0.30 < x _< 0.80, Stevenson et al. [209] included the thermally activated charge disproportionation of Mn 3+ into Mn 2+ and Mn 4+ pairs.


10.6.6 Extended Defects and Vacancy Ordering

As is known for fluorite and fluorite-related oxides, increased defect interactions are likely if the oxygen vacancy concentration exceeds 1 mol% [81]. The interaction between defects and defect association effectively lower the concentration of 'free' oxygen vacancies available for oxygen transport. In the perovskites, it is not uncommon to have an oxygen deficiency of 10 tool% or more. As noted before, the assumption of randomly distributed point defects at such large vacancy concentrations probably is an oversimplified picture. Van Roosmalen and Cordfunke [210] showed that by assuming divalent transition metal ions in undoped perovskites LaBO3_~ (M = Mn, Fe, Co) to be bound to oxygen vacancies, forming neutral defect clusters of the type <B'-V6-B'>, the model fit to the experimental data was greatly improved. Similar, but probably more complicated, extended defects were suggested to be the structural building elements of highly defective perovskites.

In general, the tendency to form ordered structures progressively grows with increasing defect concentrations. As a matter of fact, ordering of oxygen vacancies in the perovskite and perovskite-related structures seems more common than their random distribution between the perovskite slabs. Sometimes the two limiting cases are linked for the same composition by an order-disorder transformation, driven by the gain in configurational entropy of oxygen vacancies in the disordered state at elevated temperature. It has also been suggested that any ordering of the oxygen vacancies in the defective perovskite and perovskite-related oxides, thereby confining vacancy transport to two-dimensional layers, may give rise to fast ionic conductivity at significantly reduced temperatures [211,21], a point to which we return below.

10.6.6.1 Static Lattice Simulation

Attempts were made by Kilner and Brook [213] to model the ionic conductivity in perovskites LnA103 using static lattice simulation techniques. Here, the minimum energy positions for the mobile ions, and the activation energy barrier that they must surmount to migrate through the rigid crystal lattice, are calculated by minimization of the total lattice energy. The results show that aliovalent dopants might act as trapping centres for oxygen vacancies through the formation of defect associates, e.g., V6-SrLa'. It is further found that the size proportion of A and B cations is of significant importance in determining the minimum migration enthalpy for oxygen transport in the ABO3 structure. During diffusion, the migrating 0 2- ion must pass the saddle point formed by two A ions and one B ion, as shown in Fig. 10.14. The associated energy barrier to migration decreases with increasing size of the B cation and decreasing size


B cation A cations

Fig. 10.14. Saddle-pohlt configuration for oxygen anion migration.

of the A cation. This work has been extended recently by Cherry et al. [214] to include perovskites LaBO3 (Cr, Mn, Fe, Co), showing the importance of relaxation effects at the migration saddle point, which were not invoked in the previous study by Kilner and Brook. Profile mapping shows that the migrating 0 2- ion does not prefer a linear path between adjacent sites of a B O 6 octahedron, but rather follows a curved route with the saddle point away from the neighbouring B site cation.

The pioneering work by Kilner and Brook was further expanded by Cook and Sammells [215] and Sammells et al. [216] to include additional empirical relationships for predicting oxygen ionic conductivity in perovskite solid solutions, such as the average metal-oxygen bonding energy, the lattice free volume and the overall lattice polarizibility towards anion migration. A marked correlation found is that between the activation energy of oxygen anion migration in the perovskite lattice and the free volume. A smaller activation energy is apparent in ABO3 perovskites which possess an inherently larger free volume. For fluorite-structured oxygen ion conductors, again linear but opposite correlations are found. This contradiction might be solved if there is optimum value of the lattice free volume at which the Coulombic, polarization and repulsive contributions to the migration enthalpy in both type of structures are best balanced. The results obtained led to the identification of perovskite oxide electrolytes BaTb0.9In0.1034 , CaCe0.9Gd0.1034 and CaCe0.9Er0.1034, which indeed exhibit low activation energies varying between 35-53 kJ mo1-1 for the ionic conductivity [216]. Even, though, the suggested correlations guide the selection towards materials exhibiting a low value of the ionic migration enthalpy for oxygen anions and protons (if assumed to occur via OH-), none of the above authors did address the problem how to optimize the magnitude of the pre-exponential term in the expression of the overall ionic conductivity. The pre-exponential term is related in part to the density of mobile oxygen anions and the

10 m DENSE C E R A M I C M E M B R A N E S FOR OXYGEN SEPARATION 497

availability of sites (e.g., vacancies) to which they might jump. Its value will thus be determined by the state of order in the oxygen sublattice 4.

10.6.6.2 Vacancy Ordering

X-ray diffraction, electron diffraction and HRTEM have provided ample evidence that, particularly at low temperature, nonstoichiometry in oxygen-deficient perovskites is accommodated by vacancy ordering to a degree which depends both on oxygen partial pressure and on the thermal history [219,220]. A family of intergrowth compounds is observed for, e.g., the ferrites, which can be regarded as being composed of perovskite (ABO3) and brownmillerite (A2B205) structural units stacked along the superlattice axis. The intergrowth structures fit into a homologous series expressed by AnBnO3n_l, with end members n = (perovskite) and n = 2 (brownmillerite). For non-integral values of n, disordered intergrowths are observed between nearby members of the ideal series.

The concept of 'perovskite space' has been introduced by Smyth [221,222] in order to systematize the intergrowth structures exhibited by various oxygen- deficient and oxygen-excess perovskite systems. In the proposed diagram, the close structural relationship between the parent and intergrowth structures is expressed by plotting the value of n along with the compositional excursion from ideal perovskite stoichiometry. The perovskite systems are found to be quite specific in their tendency to form ordered structures in the sense that only selected values of n are found for a particular system. The observed vacancy patterns and three-dimensional structures vary along with lateral shifts in the stacking sequence of successive layers AO3_ 6. Anderson-et al. [223] showed the driving force towards ordering to be strongly dependent upon the size and electronic configuration of the B-cation, in addit ion to the size and coordination preference of the A-cation. In general, compounds that contain ordered vacancies are found for n = 5, 4, 3, 2,1.5,1.33, and 1, that is for overall oxygen contents 2.8, 2.75, 2.67, 2.5, 2.33, 2.25 and 2 [223]. Unfortunately, the structural studies are usually carried out at room temperature and little is known about the extended defects at elevated temperatures, and their behaviour dur ing quenching and cooling.

4 Ordering reduces the number of free ionic charge carriers and, in general, has a negative impact on the magnitude of the ionic conductivity. An exception to this rule is apparent in selected pyrochlores with composition Ln2Zr207 (Ln = La-Gd), and is best illustrated for Gd2Zr207 [217,218]. The pyrochlore structure can be derived from that of fluorite by ordering of both cations and vacancies on their respective sublattices. Electron microscopy has shown that the actual microstructure of these solids consists of ordered pyrochlore domains embedded in a disordered fluorite matrix. The degree of ordering can be varied by thermal annealing. Ordering causes both a lower pre-exponential term and activation enthalpy, and leads to an overall ionic conductivity for well ordered Gd2Zr207 competitive with values as measured for stabilized zirconia. The observed phenomena have been interpreted to reflect the presence of high diffusivity paths in the pyrochlore structure.


10.6.6.3 Microdomain Formation

The structural rearrangements accompanying the redox phenomena sometimes lead to a texture of the sample consisting wholly of microdomains, which may be of varying size and composition. The superstructures observed depend upon local composition, while their orientation within each microdomain may be randomly distributed. As the domains are usually smaller than the sampling size (<500 A) of X-ray and neutron diffraction, such investigations require the use of high resolution transmission electron microscopy and electron diffraction.

As an example, we briefly describe here the observations made for Lal_xCax_ FeO3_ ~. In this solid solution system, a single perovskite-brownmillerite intergrowth structure (n = 3) is found at composition x = 2/3 [224]. The intergrowth structure exhibits ideal stoichiometry, i.e. LaCa2Fe308, at reduced oxygen partial pressures near the minimum observed in the electrical conductivity [199]. For any other composition, a disordered intergrowth is observed.

Oxidizing LaCa2Fe3Os in air at 1400~ and quenching to room temperature produced a material which appeared cubic to X-rays. The most simple explanation would have been that oxygen vacancies become disordered in a highly defective perovskite structure, but transmission electron microscopy and diffraction revealed the existence of three-dimensional intergrowths of domains of the parent phase [224]. Oxygen excess is therefore considered to be accommodated at the domain walls, suggesting that the size of the domains is closely related to the oxygen excess, i.e. the greater the oxygen excess the smaller the domains.

In the regions 0 < x < 2/3 and 2/3 < x < 1, disordered intergrowths are observed with members n = 2 and n = ~, respectively. In the reduced samples, it is assumed that the local composition is very heterogeneous and a variable amount of oxygen can be accommodated as the Ca/La ratio changes. Oxidizing the disordered intergrowth structures again led to quenched samples with a microdomain texture having an estimated size of domains of about (100 A) 3. In each of these cases up to six types of microdomains were observed, three for each of the end members, e.g. LaCa2Fe308 and Ca2Fe205 for the region 2/3 < x < 1. Since the fringe separation within each of the microdomains was found to be very regular, the oxygen excess is probably again located at the domain walls [225]. In fact, the domain texture appeared to be not very stable either when kept in air at room temperature or under reducing conditions (as present in the electron microscope). After long operation in the electron microscope up to nine sets of microdomains were observed in the oxidized composition LaCa2Fe3Os.2~, three for each member in the ideal series, n = 2 (Ca2Fe2Os), n = 3 (LaCa2Fe3Os) and n = ~ (LaFeO3).

Other materials exhibiting a three-dimensional microdomain texture include CaMnl_xFexO3_~ [226,227], BaxLal_xFeO3_~ [228] and Sr2CoFeO5 [229], while a few


more examples are given below. Although it may be anticipated that mass transport is strongly influenced by the presence of a microdomain texture, the question is open as to whether these microdomains are also present at high temperatures. In a study by M6ssbauer spectroscopy, X-ray diffraction and magnetic susceptibility, Battle et al. [229] found a microdomain texture in partially oxidized samples Sr2CoFeOs~ quenched in air from 1200~ but not in substoichiometric samples obtained either after quenching in liquid nitrogen from 1200~ or after annealing under argon at 800~ Furthermore, the brownmillerite microdomains thus produced in air-quenched Sr2CoFeO54 were found to be significantly larger in size in the centre of a pellet, while oxidation at the surface had produced a new phase, presumed to be perovskite. The authors arrived at the conclusion that the microdomains were produced by the accommodation of excess oxygen in the domain walls, arising from the diffusion of oxygen into the oxide lattice during quenching, and therefore would not be in thermodynamic equilibrium.

10.6.6.4 Brownmillerite Structure

The well-known brownmillerite structure with ideal stoichiometry A2B205 can be derived from the ideal perovskite lattice by regularly removing one sixth of the oxygen atoms, resulting in an orthorhombic ~ q2-a x 4a x ~-a supercell. The ordering of oxygen vacancies thus creates layers of corner-shared B O 6

octahedra (O) alternating with layers of corner-shared B O 4 tetrahedra (T) stacked in an ...OTOTOT... sequence, as shown in Fig.t0.15. No less than six unique vacancy patterns are known for this type of structure [223]. The structures tend to disorder at elevated temperature, so as to produce a material with a high intrinsic concentration of mobile oxygen vacancies [36], i.e. not achieved by aliovalent doping. Indeed, Ba2In20 5 which has the brownmillerite structure of Ca2Fe205 becomes a fast oxygen conductor, with ionic conducfivities exceed- ing those of stabilized zirconia, after a first-order transition at about Tt = 930~ at which the oxygen sublattice disorders [211]. The situation is analogous to that of 5 - B i 2 0 3 , which undergoes a first-order order--disorder transition at T t - 730~ [230]. On the other hand, no increased ionic conductivity is found after a similar first-order transition in the isostructural Sr2Fe205 at Tt = 700~ which suggests an ordered arrangement of oxygen vacancies above Tt, while Ca2Fe205 remains stable up to at least 1100~ [231,232]. Examination of Sr2Fe205 by M6ssbauer spectroscopy showed that tetrahedral coordination of Fe persists even in the apparently disordered material. It was therefore suggested that the ordering breaks up into microdomains of a size not detectable by XRD measurements [233].

Seeking for solutions to lower Tt in Ba2In205, Goodenough et al. [211,212] extended the foresaid investigations to include Ba3In2MOs (M = Zr, Hf, Ce). Ideal ordering of the oxygen vacancies in these compounds would render them

500 10 m DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

OT

O

T

[110]c i

~

(a) (b)

Fig. 10.15. Idealized structure of (a) cubic perovskite (CaTiO3) and b) orthorhombic brownmiUerite (Ca2Fe205) lattice with ordered oxygen vacancies if-I) along the cubic [110] direction.

isostructural with Ca3Fe2TiO8, in which two octahedral (O) layers alternate with one tetrahedral (T) layer in an ...OOTOOT... sequence along the c-axis [234]. Accordingly the structure may be regarded as an intergrowth (n = 2) of perovskite layers alternating with brownmillerite layers. Enhanced ionic conduction of, in particular, Ba3In2Zr208 at modest temperatures was interpreted to reflect the disordering of the In 3+ and Zr4+-cations over the cation sublattice, which would lead to interstitial oxygen in the T-layers and correspondingly to oxygen vacancies in the O-layers, due the preference of the Zr4+-cations for octahedral coordination versus tetrahedral coordination. In a subsequent study [235], however, these high values could not be reproduced. In fact, long-range ordering does not occur, and the residual low-temperature conduction is due to facile insertion of oxygen and/or water even below 400~ In moist air, the ceramic discs crumbled on repeated thermal cycling. Only short-range ordering of the oxygen vacancies is anticipated. It is further suggested that any ordering of the oxygen vacancies in the oxygen deficient perovskites so as to create layers of corner-shared B O 4 tetrahedra (T), as in Ca3Fe2TiO8, makes more facile the insertion of oxygen and/or water at modest temperatures.

10.6.6.5 High Temperature NMR

A point defect model has been proposed for the brownmillerite structure, based on a study of electrical conductivity and emf measurements of Ba2In20 5 [236]. The experimental results matched well with modelling, while evidence was found for protonic conduction at low temperatures. In modelling, the unoccupied oxygen sites relative to the perovskite structure, below Tt, a r e

regarded as structural units and therefore are potential interstitial sites for

10 - -DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION 501

oxygen. Charge neutrality is then dominated by intrinsic anion Frenkel disorder, i.e. [Oi"] = [Vo']. Above Tt, charge neutrality would be determined by the condition 2[Vo'] - [ I n l n ] .

Partially at variance with these findings, Adler et al. [237] arrived at the conclusion that even above the order-disorder transition observed at T t = 925~ local order in the oxygen sublattice persists. Vacancy transport in Ba2In205 would be confined mainly to two-dimensional layers, whilst oxygen anions are bounded and thus immobile in adjacent layers. Using high-temperature 170- NMR and X-ray diffraction, it was demonstrated that the material retains its orthorhombic brownmillerite structure until ~1075~ at which point it becomes cubic. The density of mobile oxygen anions was found to increase continuously between 925~ and ~1075~ and only above -~1075~ the full population of oxygen anions becomes mobile.

The investigations were extended by Adler et al. [238] to include BaIno.67Zro.3303. & BaIno.67Ceo.3303_ 6, and the cobalt-containing perovskites La0.5Ba0.sCo0.?Cu0.3034i and La0.6Sr0.4Co0.sCu0.2034 i. Whilst room temperature X-ray and neutron powder diffraction revealed simple long-range cubic symmetry, HRTEM images and electron diffraction patterns indicated that these materials possess microdomains with a layered-like structure on a length scale 50-500 A. Measurements of high-temperature 170-NMR made apparent that in all these phases only a few oxygen vacancies are mobile below 800~ suggesting that the remainder is trapped in locally ordered layers. Owing to intrinsic paramagnetism of the cobalt-containing compounds, the application of 170- NMR is not possible at room temperature. Only by rapid motion of oxygen, the spectra become visible. For both cobalt-containing perovskites, the signal inten- sity above 800~ was found to increase steadily with temperature, up to the maximum temperature of 950~ in the experiments by Adler et al. [238], suggesting a concomitant increase in the number of mobile oxygen anions. In the case of La0.6Sr0.4Co0.8Cu0.2034, the value of the ionic migration enthalpy for oxygen transport estimated from NMR, viz., 1.13 + 0.10 eV, showed good agreement with the 1.04 eV obtained from 4-point ionic conductivity measurements.

It is clear that high-temperature NMR is an important tool to study the local structure and oxygen dynamics in the highly defective perovskites, which certainly merits its further use. In the case of La0.6Sr0.4Co0.sCu0.2034, Adler et al. [238] assumed that oxygen nonstoichiometry is accommodated by local ordering of MOs square pyramids in a manner similar to that in YBaCuFeO5 [239] and YBaCo2_xCuxO544 [240]. A similar type of structure was hypothesized for Lal_xSrxCoO34i (x = 0.50 and 0.70) by Van Doorn et al. [241]. As seen in Fig. 10.16, two different sites for oxygen are present in the proposed structure for Lal_~SrxCoO3_s. Four oxygens are in the basal plane of the square pyramid and one at the apex. Electron diffraction and HRTEM of powders, annealed about 15 h in air and furnace-cooled, indicated that the presence of the tetragonal super-


i i i

�9 Co (~ La/Sr @ 0 I-I O-vacancy

Fig. 10.16. Ideal structure proposed for Lal-,SrxCoO34 (x = 0.5 and x = 0.7). After van Doom et al. [241].

structure, corresponding to a doubling of the pseudocubic perovskite unit cell, gives rise to a texture consisting of microdomains with a mean size estimated at ca. (500 ~ ) 3 .

10.6.7 Observations from Permeability Measurements

This section discusses selected observations from oxygen permeation experiments. General trends have been summarized already in Section 10.6.1.

10.6.7.1 SrCoo.sFeo.203.6

The highest oxygen flux at 850~ in the series Lal_~SrxCol_yFeyO3_~ studied by Teraoka et al. [37] was found in SrCoo.sFe0.2034 (see Table 10.1) and this composition has been reexamined by a number of investigators [13,41,153,154]. In a study of oxygen permeation through SrCo0.8B0.2034 (B = Cr, Fe, Co, Cu), Kruid- hof et al. [154] showed for the first time that in SrCo0.8Fe0.2034 a change in permeation mechanism, corresponding with an order-disorder transition, occurs at about Tt = 790~ This change has been observed by others too [13,153] but was not noted in the previous study by Teraoka et al. [37]. Results from thermal analysis and X-ray powder diffraction under controlled oxygen partial pressures indicate that the observed phenomena can be attributed to the transition of SrCo0.sFe0.203_6 from vacancy-ordered brownmillerite to defective perovskite, occurring at reduced oxygen partial pressure. Both the structures of the cubic p e r o v s k i t e and o r t h o r h o m b i c b r o w n m i l l e r i t e - t y p e phases of SrCo0.8Fe0.2034 have been refined using powder neutron diffraction data by


Harrison et al. [242]. A phase diagram was presented by Qiu et al. [153], showing that only at relatively high oxygen partial pressure (>0.1 atm) and high temperature the perovskite phase is thermodynamically stable. At relatively low oxygen partial pressure and low temperature a perovskite-brownmillerite two-phase region is found. The brownmillerite phase has only a small homoge- neity region around 3 - 5 = 2.5. Below Tt, the situation during flux measurements therefore becomes very complicated, considering the fact that the Po2- gradient across the membrane also may cross the two-phase region provided, of course, that such a gradient is imposed during experiment. The studies report slow kinetics of transformation between the brownmillerite and perovskite phases in view of the long times for the oxygen flux to reach steady-state conditions at these modest temperatures. Kruidhof et al. [154] attributed these to a progressive growth of microdomains of the ordered structure in a disordered perovskite matrix.

Based on experiments, in which the membrane thickness was varied in the range 5.5-1.0 mm, Qiu et al. [153] arrived at the conclusion that the surface oxygen exchange process is the rate limiting step in the overall oxygen permeation mechanism. Further experimental evidence that the oxygen fluxes through SrCo0.8Fe0.2034 are limited by the surface exchange kinetics was given by the present authors [41]. Fitting the oxygen permeation fluxes obtained from measurements at 750~ under various oxygen partial pressure gradients to Eq. (10.18) yielded a positive slope of n = +0.5, where a value between 0 and -0.5 is expected from the experimentally observed In 5--ln PO2 relationship" However, these results merit further investigation as the flux data were taken at a temperature just below the order--disorder transition in this material.

It is already known for some time that SrCoO3_ 6 transforms reversibly from a brownmillerite-like structure to defective perovskite at about Tt = 900~ in air. Kruidhof et al. [154] observed that the transition temperature is not, or only slightly, affected if SrCoO3_ a is substituted with either 20 mol% Cr or Cu at the Co-sites. Interesting to note is that the oxygen flux for the undoped and doped specimens is very small below T t, as expected for an ordered arrangement of oxygen vacancies, but is found to increase sharply (between 5-6 orders of magnitude) at the onset of the phase transition to defective perovskite, up to values between 0.3-3 x 10 -7 mol c m -2 s -1. In view of these results, the perovskite phase in SrCoO3_ a s e e m s to be stabilized by the partial substitution of Co with Fe, but not with Cu or Cr, thereby suppressing the brownmillerite-perovskite two phase region to lower oxygen partial pressures.

10.6.7.2 Experimental Difficulties

In a number of studies, the oxygen fluxes through, e.g., SrCo0.sFe0.203_a have been reported to be significantly lower than claimed by Teraoka et al. [37]. Such


conflicting results reflect the difficulties in measuring the oxygen fluxes at high temperatures and may, at least partly, be due to specific conditions, including (1) edge-effects associated with the required sealing of sample discs to avoid gas bypassing, giving rise to non-axial contributions to the oxygen flux, (2) possible interfacial reactions when a glass is used for sealing, (3) undesired spreading of the glass seal (when its softening temperature is too low) over the oxide disc surface, and (4) the precise value of the Po2-gradient across the membrane. With regard to the first point, it is frequently the cross-sectional area of the disc that is used in the calculation of the oxygen flux. In the usual experimental arrangements however an appreciable portion of the membrane is 'clamped' between impermeable annular plates or glass rings. This edge effect means that the usual assumption of one-dimensional diffusion is not strictly correct. Another contribution to non-axial transport is that of flow of oxygen through the side walls of disc specimens, if left uncovered. Appreciable errors creep in if these edge effects are neglected as shown, for example, on the basis of a solution of Fick's second diffusion equation (with a constant diffusion coefficient) by Barrer et al. [243]. This is further demonstrated in Fig. 10.17, showing the effect of sealing edges on the departure from one-dimensional diffusion. These results were obtained from a numerical procedure to solve the steady-state diffusion equation in cylindrical coordinates [112]. Neglecting edge effects corrupts analysis of experiments in which the membrane thickness is varied, and may lead to erroneous conclusions when one tries to infer from the acquired data the influence of the surface exchange kinetics on overall oxygen permeation. Finally, it cannot be excluded that the observed oxygen fluxes are specific for the particular sample under investigation and may be affected, for instance, by microstructural effects, a point to which we return in Section 10.6.7.5.

The gas flow rate of, in particular, the inert gas used to sweep the oxygen-lean side of the membrane affects the Po~ -gradient across the membrane. Under ideal gas mixing conditions, the Po2 at the oxygen-lean side of the membrane is determined by the amount of oxygen permeating through the membrane. If the flow rate is not adjusted to obtain a constant Po2 at this side of the membrane, but a constant gas flow rate is used, the Po2-gradient gets smaller with increasing oxygen flux. This may give rise to an apparent activation energy for overall permeation, which may depart significantly from the one derived if a constant Po2 were maintained at this side of the membrane [148,149,153]. The adjustable range of the sweeping gas flow rate (to a constant Po2 at the outlet of the reactor) may be limited during experiment, being determined by the requirement that the reactor behaviour remains close to that of a CSTR (continuous stirred tank reactor).

Using a constant value of Po2 at the oxygen-lean side of 2 mm thick disc membranes of Lal_~SrxCoO3, s (x = 0.2, 0.3, 0.4, 0.5 and 0.6), Van Doorn et al. [148,149] showed that the activation energy Eac t for oxygen permeation in the


(a) 2b

membrane

r seal

2a

(b) 1 . 4 0

C9 t _

0 d k . J

o r

o . I

t _ .

(1) E 0 ID

1.30 ~ 3.75

1.20

7.5

1 . 1 0

1 . 0 0 , , , , ,

0 . 7 0 0 . 8 0 0 . 9 0 1 . 0 0

a/b Fig. 10.17. (a) Schematic cross-section of a disk membrm~e. Dashed parts indicate insulating boundaries. (b) Influence of sealing edge-effects oil the departure from one-dimensional diffusion. A geometric factor G is used for correction of the flux (normalized to surface area with diameter 2a).

Relevant parameters are defined in Fig. 10.17a.

range 900-1100~ decreases from 164 kJ mole -1 for x = 0.2 to 81 kJ mole -1 for x = 0.6. Opposed to these results, Eac t decreased from 121 kJ mole -1 to 58 kJ mole -1 w h e n a constant gas flow rate of the hel ium was used. Besides an

improved fit to the Arrhenius equat ion in the former case, Eac t can be correlated wi th the sum of the enthalpies for migra t ion and that for the format ion of oxide ion vacancies for each of the invest igated composit ions. Such a correlation is expected if oxygen t ranspor t is dr iven by the gradient in oxygen nonstoichio-

me t ry across the membrane due to the imposed Podgradient . It suggests that oxygen vacancies are free and non-interactive in Lat_xSrxCoO34 under the con-

5 0 6 10 ~ DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

ditions covered by experiment. Oxygen permeation fluxes for strontium-doping levels above x = 0.6 were found to be partially controlled by the surface exchange kinetics, as already mentioned in Section 10.6.4.

In contradiction to the observed behaviour at high temperatures, results from thermal analysis and oxygen permeation measurements indicated that a phase transition, with a small first order component, probably related with order-disorder of oxygen vacancies, occurs in selected compositions Lal_xSrxCoO3,s in the range 750-775~ [148,149]. Long times extending to over 30 h were needed for equilibration towards steady-state oxygen pemeation at these modest temperatures. Such a behaviour is reminiscent of that observed for SrCo0.8Fe0.203@ where this can be attributed to the slow kinetics of the transformation between the brownmillerite and perovskite phases at modest temperatures. In the case of La1_vSrxCoO34 (x = 0.50 and 0.70), microdomains were observed in electron diffraction and HRTEM, corresponding to ordered arrangements of oxygen vacancies in these compounds at room temperature, as mentioned in the previous section.

Another factor that is considered to be responsible for a reduced oxygen flux is the surface modification of the perovskite oxide membrane by reaction with impurities in the gas phase, as emphasized by Qiu et al. [153]. Referring to the surface degradation by reaction with minor amounts of CO2 and corresponding deterioration of the properties observed for YBa2CuO6+ x superconducting thin films [244], a similar modification effect could occur when, e.g. ambient air is used as the source of oxygen at the membrane feed side. With the help of N2 and 02 admixed to feed side pressure Po2' = 0.21 atm, Qiu et al. found the oxygen fluxes through SrCo0.8Fe0.2034 in the range 620-920~ to be larger by a factor of about 6 than when ambient air was used as feed gas, but still a factor of about 5 smaller than measured by Teraoka et al. Similar experiments were conducted in our study on SrCo0.8Fe0.203, s [41,154], where this effect was not noted in the temperature range 700-950~ so that we are inclined to believe that other factors must account for the disagreements in oxygen fluxes. This interpretation is supported by experimental evidence disclosed in a number of patents: that the oxygen fluxes through perovskite membranes remain stable as long as these are operated above certain critical temperatures, the precise value depending on the type of alkaline-earth dopant applied. Below these temperatures, a loss in oxygen flux may be observed over a period of about 100 h by as much as 30-40% when a membrane is exposed to CO2 and H20 impurities in the feed gas. This is further exemplified in Section 10.7.

10.6.7.3 Surface Exchange Kinetics

Attention has already been drawn to the importance of the surface exchange kinetics in determining the rate of oxygen permeation through mixed-conduct-


ing oxides in Section 10.3.2.2. Though for the perovskites a value of 100 ~tm is often quoted for the characteristic membrane thickness Lc, at which the change over from bulk to surface control occurs, in a number of cases much higher values are found, up to about 3000 ~tm (Table 10.2). As was emphasized earlier, the parameter L c is not an intrinsic material property and, hence, may be specific to the sample under investigation and experimental conditions. The basic assumptions made in the derivation, notably that of small Po -gradients across the membrane, may restrict its use in practical situations, where these gradients can be substantial. Experimental evidence that the oxygen fluxes are limited by the surface exchange kinetics has been found in a number of cases, as discussed elsewhere in this text.

10.6.7.4 Behaviour in Large Po2-Gradients

The mixed-conducting perovskite oxides have attracted particular interest for use as dense ceramic membrane to control partial oxidation of methane to C2-products or syngas. Such a process bypasses the use of costly oxygen since air can be used as oxidant on the oxygen-rich of the membrane.

Using SrCo0.8Fe0.2034 tubular membranes fabricated by an extrusion method, Pei et al. [13] observed two types of fracture of the tubes during the process for generating syngas. The first fracture, occurring short (within 1 h) after initiation of the reaction at 800~ resulted from the Po -gradient across the membrane and the accompanying strain due to lattice mismatch and the brownmillerite-perovskite phase transition. The second type offracture, occurring after prolonged exposure to the reducing environment, resulted from chemical decomposition towards SrCO3, and elemental Co and Fe. Similar observations have been reported for tubes made of La0.2Sr0.8Co0.4Fe0.603-~ [14], and in that study an optimized composition was also claimed, but not given, showing stable performance for up to 500 h. Using a rhodium-based reforming catalyst inside the tubes, methane conversions over 99% were achievable.

Ten Elshof et al. [10] studied the oxidative coupling of methane using a disc reactor with La0.6Sr0.4Co0.sFe0.2034 as the catalyst membrane for the supply of oxygen to the methane feed stream. Examination of the oxygen fluxes measured under various Po2-gradients in the range of thickness 0.55-0.98 mm suggested that the surface exchange reaction limits the rate of oxygen permeation. The oxygen flux was found to increase only slightly when methane was admixed with the helium used as the carrier gas. The methane was converted to ethane and ethene with selectivities up to 70%, albeit with a low conversion, typically in the range 1-3% at operating temperatures 1073-1173 K. The selectivity observed at a given oxygen flux and temperature was about twice as low if the same amount of molecular oxygen was co-fed with the methane feed stream in a single chamber reactor design, suggesting that the membrane-mode of operation

508 1 0 - DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

is conceptually more attractive for generating C2-products. Decomposition of the oxide surface did not occur as long as molecular oxygen could be traced at the reactor outlet, which emphasizes the importance of surface-controlled oxygen flux for membrane-driven methane coupling. That is, for a bulk diffusion-controlled oxygen flux the surface would become reduced by the methane, until the depth of reduction has progressed up to a point where the oxygen flux counterbalances the consumption of oxygen by methane. On the one hand the slow surface exchange kinetics observed on La0.6Sr0.4Co0.aFe0.2034 limits the magnitude of the oxygen fluxes, on the other hand its existence prevents the oxide surface from reduction, i.e. as long as the rate of oxygen supply across the membrane exceeds the rate of (partial) oxidation of methane. Noteworthy is that segregation of strontium occurred on both sides of the membrane, as confirmed by depth-profiling Auger analysis. The extent of segregation appeared to be influenced by the imposed Podgradient across the membrane, and was also found if a pure helium stream was passed along the oxygen-lean side of the membrane.

Van Hassel et al. [150] studied oxygen permeation through Lal.~SrxFeO3,s (x = 0.1, 0.2) membranes in a disc reactor using CO-CO2 based gas mixtures to control the Po2 at the oxygen-lean side. Ambient air was used as the oxygen source at the opposite side of the membrane. At 800-1100~ the oxygen flux was found to increase linearly with the partial pressure of CO. Deposition of a 50 nm thin porous Pt layer on this side of the membrane increased the oxidation rate and likewise the oxygen flux, by a factor of about 1.8. In a separate study [245], the oxygen flux was found to be invariant with the thickness of the membrane in the range 0.5-2.0 ram, while no effect was observed upon varying the Po~ at the oxygen-rich side. It was concluded that the oxygen flux is fully limited by the carbon monoxide oxidation rate. The experimentally determined rate constants scale with Sr-content in the extended range of composition 0.1 < x < 0.4. The latter can be accounted for, in view of the fact that the oxygen deficiency of the ferrites is fixed by the dopant concentration in a wide range of oxygen partial pressure, by assuming that oxygen vacancies act as active sites in the oxidation reaction of CO on the perovskite surface following either an Eley-Rideal or a Langmuir-Hinselwood type of mechanism.

10.6.7.5 Grain Boundary Diffusivity

Besides the possibility of surface exchange limitations, oxygen transport through dense ceramics is necessarily influenced by the presence of high diffusivity paths along internal surfaces such as grain boundaries. A systematic study investigating to which extent these preferred diffusion paths contribute to the diffusivity in the perovskite oxides is however still lacking. Both impurity and solute segregation take place at grain boundaries (and the external surface) or in their close proximities (less than 3 or 4 atomic distances) during sintering


and subsequent heat treatments. An obvious consideration is that, in general, these significantly alter the magnitude of ionic transport along and across the grain boundaries. In many cases the ceramics invariably contain impurities present in the starting powder or added as a sintering aid to lower the sintering temperature and/or to achieve high density. It therefore can not be excluded that disagreements in the literature regarding the magnitude of the oxygen fluxes can be explained on the basis of different ceramic processing techniques used by various authors.

In general, the presence of high diffusivity paths is important in ceramics where lattice diffusion is slow. Analyzing 180 depth profiles using secondary ion mass spectroscopy, Yasuda et al. [246] noted a significant contribution of the grain boundary diffusion to the diffusivity in the interconnect material La0.?Ca0.35CrO34, where the tracer diffusivity is of the order of ~10 -13 cm 2 s -1 at 900~ Erroneous results were obtained when isotopic exchange was performed by gas phase analysis, which resulted in apparent tracer diffusion coefficients that were almost 2 orders in magnitude higher. More recently, Kawada et al. [184] confirmed the existence of high diffusivity paths along grain boundaries in La0.~Ca0.3CrO34 using depth profiling and imaging SIMS of lso-160 ex- changed specimens. But, oxygen permeation measurements suggested negligible contribution of grain boundary diffusion to the steady-state oxygen flux. These data were obtained at 1000~ for a sample of thickness 0.75 mm. An oxygen pump and sensor were used to control the permeate side Po2. The results are well-described by the Wagner equation assuming a random point defect scheme for La0.~a0.3CrO34, as discussed in Section 10.6.2.2.

For fast ionic conductors grain boundary diffusion will have little influence, or indeed may become blocking to the diffusion from one grain to the next as is recognized in the interpretation of impedance spectra from ionic conductivity of zirconia and ceria-based solid electrolytes. In these ceramics silicon is the most common impurity detected along with enhanced yttrium segregation. Various models to account for the effects of segregation at grain boundaries and how these affect the electrical properties have been discussed by Badwal et al. [247]. Although there is no unique model describing the ceramic microstructure, the most widely adopted model for doped zirconia and doped ceria is the brick-layer model. In this model bricks present the grains and mortar the grain boundary region, i.e. assuming the grain boundary phase to completely wet the grains [248,249]. The grain boundaries in series with the grains, along the direction of charge flow, mainly contribute to the grain boundary resistivity. For doped zirconia and ceria the grain boundary resistivity can be of similar order of magnitude or higher than the bulk resistivity. Pores at the grain boundaries can have a positive effect on the oxygen transport. It is evident that more detailed studies are needed to aid in the interpretation of oxygen transport through the mixed-conducting perovskite oxides, where similar blocking effects can be expected.

5 1 0 10 - - DENSE CERAMIC MEMBRANES FOR OXYGEN S E P A R A T I O N

1 0 . 7 F I N A L R E M A R K S

The considerations in this chapter were mainly prompted by the potential application of mixed-conducting perovskite-type oxides to be used as dense ceramic membranes for oxygen delivery applications, and lead to the following general criteria for the selection of materials

- high electronic and ionic conductivity, - high catalytic activity towards oxygen reduction and reoxidafion, - ability to be formed into dense thin films, free of micro-cracks and con-

nected-through porosity, - chemical and structural integrity (i.e. no destructive phase transition)

within appropriate ranges of temperature and oxygen partial pressure, - low volatility at operating temperatures, - thermal and chemical compatibility with other cell components, - low cost of material and fabrication.

The precise perovskite composition may be tailored for a specific application. To obtain a high performance membrane, however, many technical and material problems remain to be solved. This final section will focus on several issues, which are not yet well understood, but are thought to be of importance for further development of the membrane devices.

In the first place our understanding of factors that control and limit the interfacial kinetics is still rudimentary, and therefore should be a fruitful area for further investigation. The apparent correlation between the surface oxygen exchange coefficient ks and the tracer diffusion coefficient D* for different classes of oxides, the fluorite-related and the perovskite-related oxides, as noted by Kilner et al. [73], clearly indicate the potential of isotopic 180--160 exchange. However, a problem remains how to relate the observations (at equilibrium) from isotopic exchange to the conditions met during membrane operation. In chemical relaxation experiments, the oxide is studied after perturbation of the equilibrium state. These methods are thus complementary and probably their combined application, whenever possible together with spectroscopic techniques, such as FT-IR, UV and EPR, has a great capacity to elucidate the kinetics of surface oxygen exchange.

Though, at first glance, the limited exchange capability of the perovskites, relative to diffusion, puts limits on attempts to improve the oxygen fluxes or to lower the operating temperatures by making thinner membranes, it is expected that the surface exchange kinetics can be significantly improved by surface modification. One approach is coating with a porous surface layer which will effectively enlarge the surface area available to exchange, as discussed in Section 10.3.2.3. Improvements can also be expected by finely dispersing pre- cious metals or other exchange active second phases on the oxide surface. It is clear that further investigations are required to evaluate these innovative approaches.


As yet, more work is also required to gain insight in the role of the ceramic microstructure in the performance values of membranes, and to evaluate different processing routes for the fabrication of perovskite thin films.

Besides the technological challenge of fabrication of dense and crack-free thin perovskite films, which need to be supported if its thickness is less than about 150 ~tm, a number of other problems relate to the long-term stability of perovskite membranes, including segregation, a low volatility of lattice components, etc. Some of these problems are linked to the imposed oxygen pressure gradient across the membrane. Aside from the lattice expansion mismatch of opposite sides of the membrane, attention is drawn to the potential problem of demixing, which arises in almost all situations where a multicomponent oxide is brought into a gradient of oxygen chemical potential. The available theories predict that, if the mobilities of the cations are different and non-negligible at high temperatures, concentration gradients appear in the oxide in such a way that the high oxygen pressure side of the membrane tends to be enriched with the faster moving cation species. De- pending on the phase diagram, the spatially inhomogeneous oxide may eventually decompose. The latter may cause surprise, if the (homogeneous) oxide is stable in the range of oxygen partial pressures covered by experiment. This is why these processes have been termed kinetic demixing and kinetic decomposition by Schmalzried et al. [250,251] who were the first to study them.

Degradation phenomena have been shown to occur in, for example, Col_xMgxO, Fe2SiO4 and NiTiO3. Internal oxidation or reduction processes sometimes lead to precipitation of a second phase in the matrix of the parent phase. Another possible consequence of the demixing process is the morphological instability of the (moving) low pressure interface due to formation of pores, which may eventually penetrate throughout the ceramic. The above phenomena have been the subject of a number of theoretical and experimental studies in the last decade [252-256], to give only a brief number. A review up to 1986 has been written by Schmalzried [257]. To our knowledge, no report has been made up to now of demixing phenomena in mixed oxide ion-electronic conductors. Since they cannot be excluded to occur on the basis of theoretical arguments, this is also why the phenomena deserve (more) attention in order to be able to control deterioration of membrane materials.

Intergrowth structures in which perovskite-type blocks or layers are held apart by non-perovskite ones could offer a new strategy for identifying new materials, as was suggested earlier by Goodenough et al. [212]. In such structures, vacancy transport is confined to two-dimensional layers or to sites which link up to form channels extending throughout the crystal. An interesting variation to the BIMEVOX compounds, already discussed in Section 10.4.3.1, is found in deriva- tives of Sr4Fe6013. Its orthorhombic structure can be described as built of perovskite layers alternating with sesquioxide Fe203 layers perpendicular to the b-axis. The discovery of high levels of oxygen permeation through mixed

512 10 m DENSE CERAMIC M E M B R A N E S FOR OXYGEN S E P A R A T I O N

metal oxide compositions obtained by partial substitution of iron for cobalt, for instance, SrCo0.sFeOx recently translated into a patent for this class of materials [258]. Tubes made from the given composition showed oxygen fluxes similar to those through known state-of-the-art materials having a perovskite structure, but did not fracture in the process for preparing syngas as was found for some of the perovskite materials.

As noted before, the membrane performance could be affected by the presence of H20, CO2 or other volatile hydrocarbons in the gas phase of both compartments. As laid down in patent literature [1-3], the oxygen fluxes through Mg-, Ca-, Sr-, and Ba-doped perovskites deteriorated over time, roughly 30-50% over a time period of about 100 h, if the air used as feed gas contained several percent of H20 and amounts of CO2 on a hundreds of ppm level. It was claimed, that either no deterioration is found or the fluxes can be restored to their initial values if the temperature is raised above certain critical values, 500~ for magnesium, 600~ for calcium, 700~ for strontium and 810~ for barium. Though no explanation was given, it is possible that carbonate formation took place. One may further note that the tendency for carbonate formation increases at lower temperatures.

A surprising observation was recently made in the author's laboratory in a study of oxygen permeation through Lal_xSrxFeO34 (0.1 < x < 0.4) [151]. Long times to reach steady-state oxygen permeation at 1000~ extending over hundreds of hours were observed, yet could be avoided by exposing the permeate side surface of the membrane for a 1-2 h to 1:1 CO/CO2 gas mixture. A clear explanation cannot yet be given for this observation, which is still under investigation, though a reconstruction of the surface by the reducing ambient cannot be excluded. The oxygen permeability measured if helium was used again as the sweeping gas on this side of the membrane, was found to be limited by diffusional transport of oxygen across the membrane [151]. A similar type of observation was made by Miura et al. [152], who noticed the oxygen flux through slib-casted membranes of La0.6Sr0.4Co0.sFe0.203,sto be greatly improved if these were freed from surface impurities, like SrO, following an acid treatment.

One final point to note is the ability of acceptor-doped perovskite oxides to incorporate water, and some contribution of proton conduction therefore cannot be excluded. If water insertion occurs at low temperature, this might lead to residual stresses in the ceramics. Besides water may play an active role in the surface oxygen exchange. For example, on Bi2MoO6, which has an intergrowth structure consisting of Bi2 O2+ blocks alternating with MO42- layers of corner- shared MO 6 octahedra, exchange with 1802-enriched oxygen could not be observed experimentally [259]. On the other hand, Novokova and Jiru [260] demonstrated that exchange of water with lattice oxygen on an industrial bismuth molybdate catalyst proceeds rapidly at 200~ and is even measurable at room temperature.

1 0 - DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION 513

Acknowledgements

The authors are indebted to colleagues H. Kruidhof, R.H.E. van Doorn, J.E. ten Elshof, M.H.R. Lankhorst and B.A. van Hassel for many useful discussions and for providing experimental data. Paul Gellings and Henk Verweij are gratefully acknowledged for valuable comments and careful reading of the manuscript. The Commission of the European Communities and the Nether- lands Foundation for Chemical Research (SON) are thanked for financial support.

List of Abbreviations and Symbols

Abbreviations: BE25 25 mole% erbia-stabilized bismuth oxide BICUVOX Bi4V2-yCUyOll BIMEVOX general acronym for materials derived from Bi4V2O11,

like BICUVOX BT40 BY25 CSZ ECVD EDS EPR emf FT-IR HRTEM MIEC NMR SEM SIMS SOFC TEM tpb TPR UV XANES XAS XRD YSZ

40 mole% terbia-stabilized bismuth oxide 25 mole% yttria-stabilized bismuth oxide calcia-stabilized zirconia electrochemical vapour deposition energy dispersive spectroscopy (of X-rays) electron proton resonance electromotive force fourier transform infrared spectroscopy high resolution transmission electron microscopy mixed ionic-electronic conductor nuclear magnetic resonance scanning electron microscopy secondary ion mass spectroscopy solid oxide fuel cell transmission electron microscopy three phase boundary temperature programmed reduction ultra-violet spectroscopy X-ray absorption near edge structure X-ray absorption spectroscopy X-ray diffraction yttria-stabilized zirconia

Symbols: Ci

N

D D*

mole fraction or concentration of species i chemical diffusion coefficient tracer diffusion coefficient

514 10- DENSE CERAMIC MEMBRANES FOR OXYGEN SEPARATION

Ds Dv dp e E Eeq F G

I

ji .0

Jex ks k K L Lc Ld Lp n

P Po2 PO2' Po2"

R Si

o si

S t tel ti tion T Tt

ui o

bli

Vm zi

self-diffusion coefficient vacancy diffusion coefficient pore diameter elementary charge emf emf at equilibrium Faraday constant geometric factor used to account for non-axial contributions to the oxygen flux Haven ratio electrical current flux of species i balanced surface exchange rate atequilibrium, mol 02 c m -2 s -1

surface exchange coefficient, cm s reaction rate constant equilibrium constant for a reaction membrane thickness characteristic thickness of membrane Debye-Hiickel screening length characteristic thickness (active width) of porous coating layer frequently used to designate the mole fraction of electrons, yet its use is multipurpose mole fraction of electron holes oxygen partial pressure oxygen partial pressure at feed side of the membrane oxygen partial pressure at permeate side of the membrane radius of species i gas constant entropy of species i entropy of species i at standard state surface area Goldschmidt factor electronic transference number transference number of species i ionic transference number temperature transition temperature electrical mobility of species i electrical mobility of species i in standard state molar volume charge number of species i (positive for cations and negative for anions)

10 m DENSE CERAMIC M E M B R A N E S FOR OXYGEN SEPARATION 515

Greek: (x

8

1~i 0

o

(~el (3" h

o

(~ion (~n (3"p

(~total Xs

surface exchange coefficient bulk diffusion coefficient reduct ion factor deviat ion from ideal oxygen stoichiometry enhancement factor overpotential electrochemical potential of species i porosi ty chemical potential of species i s t andard chemical potential of species i electronic conductivi ty polaron hopp ing conductivi ty electrical conductivi ty of species i conductivi ty of species i at s tandard state ionic conductivi ty n-type electronic conductivi ty p-type electronic conductivi ty total conductivi ty tortuosity electric potential of phase (Galvani potential) critical (percolation threshold) vo lume fraction

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