On phase relations, transport properties and defect structure in mixed conducting SrFe1.5−xCoxOz

Solid State Ionics 129 (2000) 285–297 www.elsevier.com / locate / ssi

On phase relations, transport properties and defect structure in mixed conducting SrFe 1.52x Co x O z a, b Rune Bredesen *, Truls Norby a

b

SINTEF Materials Technology, P.O. Box 124, Blindern, N-0314 Oslo, Norway ´ 21, N-0349 Oslo, Norway Department of Chemistry, University of Oslo, Centre for Materials Science, Gaustadalleen Received 1 April 1999; accepted 27 November 1999

Abstract The stability of the Sr 4 Fe 6 O 13 phase, and the phase changes taking place in the presence of Co, are discussed. Reported transport parameters are analysed and discussed in relation to defects and phase changes in SrFe 1.52x Cox O z . A defect model is suggested that is in qualitative agreement with the observed transport parameters. It is suggested that defect ordering due to Co 21 ions is decisive for the transport properties.  2000 Elsevier Science B.V. All rights reserved. Keywords: Phase relations; Electronic and ionic conductivity; Chemical diffusion; Defect models of pure and Co-containing Sr 4 Fe 6 O 13

1. Introduction Transport properties of metal oxides are of significant importance in many technological applications. Well known phenomena like creep, sintering, and oxidation of metals are usually controlled by simple diffusion of matter. High diffusion rates represent in many cases an instability problem, limiting the range of operation and technological utilisation. However, there are also applications where high diffusion rates are desirable, such as, for instance, in membrane materials. Dense oxide membranes may conveniently be divided into two main classes of materials: electrolytes, where only ionic transport is important, and mixed conductors, where both ionic and electronic transport are important. *Corresponding author. Tel.: 147-2-206-7522; fax: 147-2206-7350. E-mail address: [email protected] (R. Bredesen)

High temperature oxide electrolyte materials for membrane applications are typically oxygen ion or proton conductors. During operation in chemical or / and electrical gradients, ions diffuse from one side of the membrane to the other, creating an electrical current. Since the electronic resistance is very high the electrons are transported in an outer loop closing the circuit. Solid oxide electrolyte membranes can therefore be used to generate electrical energy in so-called solid oxide fuel cells (SOFC), illustrated in Fig. 1. Such systems may also operate in the opposite direction, i.e. by applying an electrical potential across the electrolyte the ions can be forced to move in the membrane. This is the principle of electrolysers and pumps, which can be used to produce, enrich, and purify gases such as hydrogen and oxygen. The currently high operating temperatures of oxygen conducting solid oxide electrolyte systems, typically 9008C, place strong demands on the stabili-

0167-2738 / 00 / $ – see front matter  2000 Elsevier Science B.V. All rights reserved. PII: S0167-2738( 99 )00333-1

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Fig. 1. Illustration of working principle of solid oxide electrolyte (left) and mixed conducting (right) membrane.

ty and compatibility of the different cell materials. A significantly simpler system for oxygen separation can be envisaged when the membrane can transport both ions and electrons (see Fig. 1). In this case there is no need for an external electrical circuit or electrode / interconnect materials. However, these membranes may only operate under an external oxygen partial pressure gradient, i.e. oxygen pump-

ing by applying an electrical potential is not applicable due to the high electronic conductivity. In oxygen separation with mixed conducting membranes the operational cost is related to pressurising and heating air to the working temperature (800– 9008C). Combined integrated processing (oxygen production 1 chemical conversion) where heat is generated, as in partial oxidation or combustion, is therefore attractive. A typical example of combined processing is synthesis gas production from natural gas where a membrane reactor may replace a cryogenic oxygen production plant [1]. Fig. 2 gives an illustration of the integrated production of oxygen in the secondary reactor of a combined (steam reforming 1 partial oxidation) synthesis gas production process. Estimations have shown that the cost of syngas production may be reduced by up to 50% [2,3] if membranes replace cryogenic systems. Cost reductions appear possible already at flux levels of approximately 10 cm 3 O 2 / cm 2 min, currently achieved in lab-scale experiments [2]. Other interesting possibilities are to integrate oxygen selective membranes in gas turbine electrical power generation and coal gasification [3]. In general, integrated production and use of oxygen instead of air in chemical industry and power production could re-

Fig. 2. Simplified flow sheet of synthesis gas production by combined reforming using an oxygen selective membrane.

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duce plant dimension, separation costs, and facilitate CO 2 recovery from exhaust gases. Several types of mixed conducting materials have been investigated for possible use in oxygen membrane separation [4]. These include oxide–metal composites, and various types of mixed conducting oxide materials. While fluorites (e.g., CeO 2 ), pyrochlores (e.g., Gd 2 Ti 2 O 7 ) and traditional perovskites (e.g., LaCoO 3 -based compositions) have been in focus for many years, very high ionic and mixed conductivity were recently reported in novel materials with the composition SrFeCo 0.5 O z [5]. Recently, a total conductivity of approximately 20 S cm 21 with an ionic component close to 10 S cm 21 has been reported in air at 9008C [6]. These promising findings have created a lot of interest in mixed conducting oxide membranes, and, furthermore, opened up new insights which increases the possibility of identifying novel candidates. Before discussing properties further a brief review of these recent materials will be given for completion.

287

Fig. 3. Illustration of the perovskite structure SrFeO 32d [10].

2. Brief survey of mixed conductor SrFe 1.52x Co x O z materials

2.1. Phase relations Soon after the initial discovery of the properties of SrFeCo 0.5 O z it was found that the material tends to consist of two major phases, a layered perovskite related structure Sr 4 Fe 62y Co y O 136d and a normal perovskite SrFe 12y Co y O 32d [7,8]. The layered perovskite related structure is isostructural with Sr 4 Fe 6 O 13 [9]. The structure Sr 4 Fe 62y Co y O 136d can be visualised by layers parallel to the a–c plane of a perovskite related layer, with FeO 6 octahedral coordination, and a double layer with (Fe,Co)O 4 and (Fe,Co)O 5 coordination polyhedra [8]. The presence of Co only in the double layer where the number of coordinating oxygen is lower was discovered in neutron diffraction studies. Figs. 3–5 illustrate the structures of the perovskite, brownmillerite and layered perovskite related phases. The brownmillerite structure is an ordered oxygen-deficient perovskite related structure, which at high temperatures may transform into a disordered state similar to a perovskite.

Fig. 4. Illustration of the brownmillerite structure Sr 2 Fe 2 O 5 [11].

Below we will use the notation Sr 4 Fe 62y Co y O 136d when we refer specifically to the layered perovskite related phase, and SrFe 1.52x Co x O z when we refer only to an overall composition. The latter notation is used both for single phase and multiphase materials. Several studies have recently demonstrated that

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288

Fig. 5. Illustration of the layered perovskite related structure Sr 4 Fe 62y Co y O 136d [8].

the stability of the layered perovskite related phase is dependent on temperature, oxygen pressure and the cobalt content, x (SrFe 1.52x Co x O z ) [6,7,12,13]. A single layered perovskite related phase can be obtained for 0 # x # 0.375 in samples prepared at 1150–12008C in air. For 0.425 # x # 0.525 it coexists with the perovskite, while for x . 0.6 only the perovskite phase is formed along with CoO [7]. Lowering the temperature appears to stabilise the Sr 4 Fe 62y Co y O 136d phase relative to the normal perovskite phase [14]. However, phase transitions are sluggish and may go on for weeks at 9008C. In studies where samples with varying Co contents were quenched from reducing atmospheres ( pO 2 5

10 210 atm) at 11008C, it was found that the amounts of the layered perovskite related phase increased with decreasing Co content [13]. While the Co-free sample (x 5 0) contained only minor amounts of second phases (SrFeO 32d ), a sample with x 5 0.375 contained no layered perovskite related phase at all. Furthermore, in reducing atmospheres, furnacecooled samples with x 5 0.5 show two major phases, a brownmillerite and cobalt oxide [6,12]. Evidence of much higher stability of Co-free Sr 4 Fe 6 O 13 has also been reported from conductivity measurements between 950 and 10508C in the oxygen partial pressure range ¯1–10 212 atm [15]. On the basis of the behaviour of the conductivity, no clear evidence of phase changes was revealed in this oxygen pressure–temperature window. All in all, current knowledge clearly indicates that Co-containing Sr 4 Fe 6 O 13 has only a limited stability regime compared to the Co-free material.

2.2. Transport properties It has been shown that, in air, the total conductivity of SrFe 1.52x Co x O z increases with Co content from x 5 0 to 0.5 [6,7] (see Table 1). The electrical conductivity of the Co-free compound, Sr 4 Fe 6 O 13 , displays a relatively shallow minimum at around pO 2 5 10 23 atm, which is interpreted as ntype conduction at low and p-type conduction at high oxygen activities and with a pO 2 -independent ionic component [15]. The n-type contribution is roughly proportional to ( pO 2 )21 / 4 down to around 10 28 atm. Below this oxygen activity, it approaches a ( pO 2 )21 / 6

Table 1 Transport parameters for SrFe 1.52x Co x O z at 9008C Composition SrFe 1.5 O z SrFe 1.5 O z SrFe 1.5 O z SrFe 1.45 Co 0.05 O z SrFe 1.35 Co 0.15 O z SrFe 1.25 Co 0.25 O z SrFe 1.20 Co 0.30 O z SrFe 1.20 Co 0.30 O z SrFe Co 0.5 O z SrFe Co 0.5 O z SrFe Co 0.5 O z

stot (air) (S / cm)

sion (air) (S / cm)

sel (air) (S / cm)

stot (Ar) (S / cm)

D?10 6 (cm 2 / s)

0.5 1.0 8.9 1.9 2 1.2 9.9 10.0 23.5

7.9

2.0

9.3

14.2

0.6 6 1.2 0.9

Ref. [6] [7,15] [16] [14] [7] [14] [6] [7] [6] [14] [17]

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dependency, but it is uncertain whether this reflects a partial decomposition, or a real change in defect structure. However, at higher pO 2 there is little doubt that the defect structure may be well described by a simple defect model with a dominating concentration of ionic point defects [16], which is largely independent of pO 2 , and with the electronic carriers (electrons and holes) as minority defects. Studies of the electrical conductivity in SrFeCo 0.5 O z as a function of the oxygen partial pressure show that stot ~( pO 2 )1 / 6 2 ( pO 2 )1 / 4 from 1 to 10 212 atm O 2 [18,19]. At lower oxygen partial pressures the conductivity may pass through a minimum and increase on further reduction in partial pressure [19]. The ionic conductivity has been reported to fall approximately one order of magnitude when the partial pressure of oxygen is reduced from air to 10 23 atm O 2 , whereafter it levels out and becomes independent of the oxygen pressure. In air between 950 and 6508C, the ionic transference number decreases gradually from 0.4 to 0.2 [19]. The activation energy of the ionic conductivity has been reported to be as low as 0.37 eV in the same temperature range [20]. Studies of the oxygen permeation in SrFeCo 0.5 O z membranes show that bulk diffusion appears to be rate limiting [6,21,22]. At 9008C the chemical diffusion coefficient, D, determined by the conductivity relaxation method, is reported to be ¯9310 27 cm 2 / s [17] with an activation energy of 0.92 eV. These values are in good agreement with values obtained by transient thermogravimetric measurements of SrFe 1.52x Co x O z (0 $ x $ 0.5) [14,16]. These latter results also suggest that D decreases with Co content up to x 5 0.25 (see Fig. 6). Chemical diffusion coefficients reported in steady state permeation measurements [22] are approximately 1.5 orders of magnitude lower at 9008C. All diffusion measurements described above were performed in the same oxygen partial pressure range of 1–10 25 atm. Up to now, the observed transport properties of pure Sr 4 Fe 6 O 13 have been fit into a common defect model [15,16]. Similar attempts to rationalise the behaviour of the SrFeCo 0.5 O z material [19] appear more speculative since later findings clearly indicate that this composition is outside the stability range of the Sr 4 Fe 62y Co y O 136d single phase region. In this presentation these aspects are re-examined in the

289

Fig. 6. Values of chemical diffusion coefficients as a function of reciprocal temperature obtained by transient thermogravimetric measurements [14,16], transient electrical conductivity measurement [17] and steady state permeation measurements [22].

light of recent results on phase relations and transport properties of Co-free and Co-containing Sr 4 Fe 6 O 13 . Our aim is to present a defect model for Co-containing Sr 4 Fe 6 O 13 which is qualitatively consistent with the observed transport properties.

3. Materials and methods For experimental details related to the results discussed in this presentation, the readers are referred to the experimental descriptions in the cited references, or references cited within these.

4. Discussion

4.1. Phase relations The common way of preparing SrFeCo 0.5 O z is by solid state reaction of precursor powders of metal salts. Room temperature X-ray diffraction patterns of SrFeCo 0.5 O z at different stages in the synthesis show that the regular perovskite, SrFe 12y Co y O 32d , is formed first (see Fig. 7). Higher temperatures and long term annealing promotes formation of the layered perovskite related Sr 4 Fe 62y Coy O 136d phase. These observations suggest that the formation of the

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In Eq. (1) for the sake of illustration it has been assumed that oxygen uptake occurs and the formed SrFe 12 O 19 is stoichiometric. However, such phase changes may involve both oxygen uptake and evolution depending on the nonstoichiometry of the different phases. Since these reactions to some extent involve transport of metal ions as well as oxygen, one may expect these to be slower than changes in nonstoichiometry only depending on oxygen diffusion. It can be seen from Eq. (1) that formation of stoichiometric brownmillerite from stoichiometric Sr 4 Fe 6 O 13 could occur without exchange of oxygen with the atmosphere (i.e. n 5 d 5 0) Fig. 7. Room temperature XRD pattern of SrFeCo 0.5 O z powders prepared by solid state reaction of precursor powders in air. (a) Calcined at 8008C for 2 h; (b) 9508C for 2 h; (c) calcined at 9508C for 16 h; (d) calcined at 9508C for 16 h and then for 1 h at 12008C.

layered perovskite related phase is slow and possibly proceeds first via formation of the regular perovskite phase. Studies using other synthesis routes show similar results for SrFeCo 0.5 O z [23]. Further evidence of slow formation of the Sr 4 Fe 62y Co y O 136d phase is given in Fig. 8, where surfaces of membranes used in transient thermogravimetric measurements at high oxygen partial pressures are shown before and after the measurements between 1000 and 7508C. The characteristic plate-shaped grains belonging to the Sr 4 Fe 62y Co y O 136d phase have increased in number during the approximate 1 week duration of the measurements. A reason for this slow formation of the layered perovskite related phase could be a low thermodynamical driving force. An indication of this feature is given by observing the effect of adding dopants to Sr 4 Fe 6 O 13 which stabilise the perovskite phase [15]. When only 10% of the Sr is substituted by La in Sr 4 Fe 6 O 13 , the material is nearly entirely converted to the perovskite (Sr,La)FeO 32d , where most of the La is found, and to SrFe 12 O 19 . The presence of SrFe 12 O 19 compensates for the excess of iron in mixed phase materials through the overall reaction (for simplicity omitting La) n Sr 4 Fe 6 O 136d 1 ]O 2 2 42 2 5 ]SrFeO [2.51(11n / 84)6( 11d / 42)] 1 ]SrFe 12 O 19 . 11 11 (1)

42 2 Sr 4 Fe 6 O 13 5 ]SrFeO 2.5 1 ]SrFe 12 O 19 . 11 11

(2)

Formation of ordered brownmillerite SrFeO 2.5 is promoted by lower temperatures and oxygen partial pressures [12,24]. Thus, it may be questioned if the layered perovskite related structure might decompose into a brownmillerite at lower oxygen partial pressures and / or temperatures and a perovskite at higher partial pressures and temperatures by coformation of SrFe 12 O 19 . When high contents of Co are present in SrFe 1.52x Co x O z , however, formation of cobalt oxide and Co–Fe-spinel appears to dominate over SrFe 12 O 19 formation [7,12,14]. Another important feature observed in La-doped Sr 4 Fe 6 O 13 is the existence of a nanoscaled intergrowth of the perovskite and layered perovskite related phases [15]. Such intergrowth, consisting of regions of approximately 10 nm in size, could only be detected by high resolution TEM investigation (not by XRD or SEM), an analysis that has not yet been carried out on other compositions. Due to the structural similarity these findings are not surprising, however it seems relevant to ask if intergrowth also exists in Co-containing materials. Further studies should be carried out to clarify this, since it can be expected that this feature will affect both the mechanical and the transport properties.

4.2. Defect structure and transport properties 4.2.1. Defect structure model of Sr4 Fe6 O13 One may ask what models may describe the defects likely to be dominant in compounds with the Sr 4 Fe 6 O 13 structure. Based on the conclusion from the dependencies of the conductivity of Sr 4 Fe 6 O 13

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291

Fig. 8. (a) Membrane surface of SrFeCo 0.5 O z as-prepared (bar 50 mm). (b) Membrane surface of SrFeCo 0.5 O z after thermogravimetric measurements at 750–10008C for approximately 1 week (bar 50 mm).

on pO 2 we take as a starting point that the defect structure is dominated by a relatively constant concentration of point defects such as oxygen vacancies with electronic carriers as minority defects. This is in general accordance with the high ionic transport rates

and modest electronic conductivities. At least three types of point defect models may be envisaged: (i) the anion-Frenkel disorder, (ii) the acceptor-doped system, and (iii) the partial occupancy disordered system. They may, as will be seen, more or less

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describe the same physical system and give similar results in terms of predictions of pO 2 dependencies for the defects.

( i) Dominating anion-Frenkel defect pairs In the rather loosely packed double layer, formation of anion-Frenkel type of defect pairs appears possible, i.e. O xO 5V ??O 1 O i99 ,

(3)

KAF [O xO ] 5 [V ??O ] ? [O i99 ].

(4)

If oxygen vacancies and interstitials are dominating defects, we may approximate the electroneutrality condition by 1/2 [V ??O ] 5 [O 99 i ] ¯ K AF .

(5)

In this model, the ionic conductivity may be dominated by transport of the vacancies or the interstitials (or both) depending on their mobilities. Furthermore, it will be independent of pO 2 but probably increase considerably with temperature due to the increase in defect concentration as well as defect mobility.

( ii) Dominating lower valent cation ‘ acceptor’ and oxygen vacancy point defects In the acceptor-doped model, we take the structure, rather than the valence of the constituent cations, as reference state. In analogy with the related perovskite (SrFeO 32d ) we will assume that the layered perovskite related material is oxygen deficient compared to its reference structure. This model implies that iron should have a higher average valence than 31 in the reference structure. Let us assume that the reference structure has one more oxygen in the double layer per unit formula, i.e. written Sr 4 Fe 6 O 14 . The reference structure then has two irons out of six oxidised to Fe 41 . One could, for instance, imagine that the two iron ions located in the perovskite layer have reference valence 41 as in the normal perovskite SrFeO 32d . However, under the applied thermodynamic conditions most Fe 41 sites are reduced to Fe 31 , which may be regarded as lower valent cation point defects. The electroneutrality condition may then be approximated by 2[V ??O ] 5 [Fe 9Fe 41 ] 5 constant

(6)

and the behaviour of the transport properties would be much the same as in the anion-Frenkel case, except that now we clearly view only vacancies as mobile point defects, and we consider their concentration as roughly temperature independent. Such a model would be in accordance with the low activation energy reported [20] for the ionic conductivity. This kind of model works nicely to explain features of the nonstoichiometry and transport properties of other oxygen-deficient perovskites, such as the ‘YBCO’ high-T c superconductors [25] and the SrFeO 32d perovskite.

( iii) Partial site occupancy model In the third model, we consider again a number of oxygen sites to be empty, but in order to avoid the concept of fully substituted Fe sites, we instead regard the structure as having a number of oxygen distributed over a larger number of equivalent sites. This is equivalent to having the three oxygen ions over the four sites in d-Bi 2 O 3 , or to the five oxygen ions over the six sites in the disordered SrFeO 2.5 or Ba 2 In 2 O 5 phases [26]. These kinds of disordered structures have seemingly not fitted well into the ¨ Kroger–Vink terminology, as there is now no perfect lattice to define the defects in relation to. In order to overcome this, we may introduce the concept of a partially occupied site as a perfect constituent of the compound. In the Sr 4 Fe 6 O 13 structure the transition metal has lower coordination in the double layer (FeO 4 or FeO 5 ) than in the perovskite layer (FeO 6 ). In a completely disordered state (as in disordered SrFeO 2.5 ) we will assume that the coordination is FeO 6 also in the double layer. The ‘lacking’ oxygen ions or structural vacancies when compared to the ordered Sr 4 Fe 6 O 13 structure would then be distributed randomly on all six oxygen coordination sites of the transition metal. For the sake of illustration, let us look at a transition metal site with five oxygen coordination in the ordered Sr 4 Fe 6 O 13 structure. Then in the disordered state the five oxygen ions occupy on average all six oxygen sites equally. In the disordered state the reference oxygen sites are only 5 / 6 occupied sites, and the average valence of each reference site is 5 3 (22) / 6 5 2 5 / 3. Thus, when a site is occupied by an oxygen ion with a real charge 22, the effective charge will be 21 / 3 (22 compared to the reference 25 / 3). Likewise, the

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293

empty site will have zero real charge, and therefore an effective charge of 15 / 3. The order–disorder reaction may with real charges be written

2≠mO =mO 2 5 ]]=c O and JO 5 2JO 2 , ≠c O

22 5O 22 O 1 Vi (ordered structure)⇔5O 5 / 6O 1

and by comparing with Fick’s first law for flow of oxygen in the oxide

V5 / 6O (disordered structure),

(7)

where Vi is a vacant interstitial site in the ordered structure, and the subscript 5 / 6O means that six equal sites share five oxygen ions. Both the occupied and the vacant site then become defects, namely (in 3? 39 ¨ Kroger–Vink notation) V 55 // 6O and O 51 // 6O . Both defects would contribute to net transport in an electrical field, as expected from a disordered system. The electroneutrality may now be expressed by 3? 5 / 3[V 55 // 6O ] 5 1 / 3[O 15 // 39 6O ].

(8)

Oxidation and reduction may be expressed by exchange of oxygen gas and electronic defects between these two species, e.g. 5 / 3? O 15 // 39 6O 5 1 / 2O 2 (g) 1 V 5 / 6O 1 2e9

(9)

and the pO 2 -dependencies under the above prevailing electroneutrality become just like for the other models (i) and (ii); a pO 2 -independent ionic conductivity and ( pO 2 )21 / 4 and ( pO 2 )1 / 4 dependencies of the electrons and holes, respectively.

4.2.2. Chemical diffusion coefficient and defect model In order to further discuss the parameters determining bulk transport we first correlate the chemical diffusion coefficient to the defects in the oxide. In general treatment of bulk transport of oxygen ions and electrons in oxides, an expression for the flow of oxygen through the membrane can be derived, which may be written [4] 1 sel ? sion JO 2 5 2 ]]2 ]]]=mO 2 , 16F sel 1 sion

(10)

where sion 5 sV ??O 1 sO0i is the sum of the conductivity of oxygen vacancies and interstitial oxygen, and sel 5 sn 1 sp is the sum of the conductivity of electrons and electron holes. =mO 2 is the change in chemical potential of oxygen gas in the direction of flow. One may write

JO 5 2 D=c O

(11)

it is seen that D may be expressed by 1 sel ? sion ≠mO D 5 2 ]]2 ]]] ]], 4F sel 1 sion ≠c O

(12)

where ≠mO / ≠c O is called the thermodynamical factor. Previously we applied the simple approach of a dominating anion-Frenkel point defect (model (i)) in pure Sr 4 Fe 6 O 13 [16], i.e. the electroneutrality condition is given by [V ??O ] 5 [O 99 i ].

(13)

As a first approach for discussion of the observed changes in D, let us assume that this defect model applies also to Co-containing Sr 4 Fe 6 O 13 . In this case it can be shown that D may be written [27] RT (sn 1 sp ) ? (sV ??O 1 sO0i ) 1 D 5 ]]2 ]]]]]]] ]]]] ?? 4F (sn 1 sp ) 1 (sV O?? 1 sO0i ) [V O ] 1 [O 99 i ]

S

D

4 1 ]] , n1p

(14)

where the term

S

1 4 ]]]] 1 ]] ?? n1p [V O ] 1 [O 99 i ]

D

is equal to 1 ≠mO ] ]] RT ≠c O and represents the thermodynamical factor. Strictly, this expression is correct only for dilute defect concentrations without defect–defect interactions, i.e. when ≠ ln a V ??O ≠ ln a O0i ≠ ln a p ≠ ln a n ]]] ?? ¯ ]]] ¯ ]] ¯ ]] ¯ 1. ≠ ln p ≠ ln n ≠ ln[V O ] ≠ ln[O 99 i ]

(15)

4.2.3. Electronic transport in SrFe1.52 x Co x Oz Previous studies suggest that the compositions Sr 4 Fe 6 O 13 , SrFe 1.45 Co 0.05 O z , and SrFe 1.25 Co 0.25 O z

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may consist mainly of the layered perovskite related phase [14,15]. Therefore, in our approach we neglect possible effects of other phases and relate the transport properties in these materials to the layered perovskite related phase only. As summarised in Table 1, the total electrical conductivity in air increases with Co content in SrFe 1.52x Co x O z . Furthermore, the limited data also suggest that the ionic transference number decreases with increasing Co content, i.e. the main change in total conductivity is caused by increasing electronic conductivity. At high oxygen pressures the electronic conductivity is, according to our defect model, determined by transport of electron holes. It should be noted that if the mobility of electron holes is much higher than the mobility of oxygen point defects, the defect model implies that [V ??O ] ¯ [O 99 i ] 4 p 4 n in air. In this case the expression for the chemical diffusion coefficient, D, may be simplified to RT sel ? sion 1 D 5 ]2 ]]] ]. stot p F

(16)

This expression consists of two parts: a conductivity and a purely defect concentration dependent term. Evaluated only on the basis of the conductivity term, one would expect D to increase with Co content (see Table 1). However, judged from the reported values of D the opposite appears to be the case. In light of the model and assumption given, this suggests that the thermodynamical factor term decreases faster with increasing Co content than the increase in the conductivity term. If the thermodynamical factor term has a dependency ~(4 /p) this implies that the concentration of mobile electron holes increases with increasing Co content in Sr 4 Fe 62y Co y O 136d. From perovskite systems of (La,Sr)(Fe,Co)O 32d , it has been observed that the electron hole conductivity increases when Fe is substituted with Co [28,29]. The same feature appears to be found in Sr 4 Fe 62y Co y O 136d , and one may speculate if thermal charge disproportionation of Co 31 to Co 21 and Co 41 , which has been suggested in the perovskite systems, could increase the concentration of electronic defects. Since iron ions may stabilise the Co 31 state [28], the particular Sr 4 Fe 62y Co y O 136d structure having Co ions only occupying sites in the double layer could perhaps facilitate Co 31 charge dispropor-

tionation. Formation of different valence states in the double layer could also facilitate electron transfer between sites. Further studies are needed to separate the electronic and ionic contributions before a closer analysis of the Co concentration dependency is possible. As previously emphasised, the composition SrFeCo 0.5 O z contains two major phases so in this case the transport properties cannot be interpreted as the layered perovskite related phase only. In Table 1 it can be seen that the electronic conductivity increases significantly when the Co content increases from x 5 0.3 to 0.5, i.e. when the stability range of the layered perovskite related phase is exceeded. This feature is most probably related to the presence of the perovskite phase, SrFeO 32d , which has higher electronic conductivity [15]. Furthermore, there is practically no further change in D from x 5 0.25 to 0.5. This could also be explained by the presence of the perovskite phase since diffusion coefficients have been reported to be higher in related perovskites [13,30,31]. In light of these findings one may speculate if the overall D may go through a minimum in SrFe 1.52x Cox Oz materials for the composition having maximum solubility of Co in Sr 4 Fe 62y Co y O 136d. Fig. 6 shows that the ratio of diffusion coefficients obtained in chemical relaxation and steady state permeation measurements, respectively, is approximately 40 near 9008C. In the latter measurements fixed oxygen partial pressures where applied on the two sides of the membrane during permeation, and the chemical diffusion coefficient, Da , was defined as [31] RT sel sion ]]. Da 5 ]] 4F 2 c O stot

(17)

Compared to the chemical diffusion coefficient from chemical relaxation measurements [14,16,17] derived from Fick’s law (Eq. (11)), the ratio D/Da (disregarding for simplicity the complication of a twophase mixture) is equal to

S

D

1 4 c O ]]]] 1 ]] , n 1 p [V ??O ] 1 [O 99 ] i according to the applied model. Using c O 5 0.074

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295

mol / cm 3 for Sr 4 Fe 62y Co y O 136d [8] this estimate gives

could explain a higher oxygen vacancy concentration according to

S

2Co xB 1 O xO 5 ]12 O 2 1 V ??O 1 2Co 9B ,

D

1 4 ]]]] 1 ]] 5 540 cm 3 / mol. ?? 99 n 1 p [V O ] 1 [O i ]

If one very roughly assumes that this is approximately equal to (4 /p) at high oxygen pressures, then p ¯ 0.007 mol / cm 3 . With electronic p-type conductivity of 1–10 S / cm (see Table 1) the electron hole mobility m p 5 sp /e ? p is in the range 10 23 –10 22 cm 2 / V s. This order of magnitude is consistent with a small polaron mobility [32]. All in all, based on very rough estimations, it seems reasonable to assume that the concentration of electron holes is rather high, and that their mobility is low. Furthermore, in light of the high ionic conductivity, and assuming a reasonable concentration of point defects, the mobility ratio of electron holes to oxygen ions appears unusually low. A similar feature has been observed in CeO 22d , where a mobility ratio of electrons to oxygen vacancies of only ¯10 results in a mixed conducting material even at low oxygen partial pressures [32]. Low electron hole mobility may originate from Fe sites having deeper polaron energy levels, and thereby trapping the polaron for longer times than Co sites [28]. The layered perovskite related structure with Co only in the double layer may in this case have higher two-dimensional electronic conductivity in the double layer, and lower electronic conductivity normal to the double layer and in the Fe-containing perovskite layer.

4.2.4. Ionic transport in SrFe1.52 x Co x Oz As previously discussed ionic conductivity in Sr 4 Fe 62y Co y O 136d originates from transport of oxygen ion vacancies and / or interstitial oxygen. A reduction in oxygen content and average oxidation state of the transition metal ion (B n1 ) has been reported with increasing Co content in Sr 4 Fe 62y Co y O 136d [7]. For y 5 1.2, d was found by iodometric titration to be equal to 20.01 and n 5 3.00, while for y 5 0, d was reported to be equal to 0.41 and n 5 3.14. These figures may be different at high temperatures, probably displaced towards lower d and n values. Preferential thermal reduction of Co

(18)

where B is a transition metal ion site. It is expected that (Co 41 , Co 31 ) will be reduced easier than (Fe 41 , Fe 31 ). If easier reduction of Co in Sr 4 Fe 62y Co y O 136d results in the observed higher ionic conductivity this suggests that diffusion involving vacancies contributes significantly to the total ionic conductivity. Furthermore, this implies that point defect models (ii) or (iii) mentioned earlier appear to be the more probable. Formation of oxygen vacancies also results in a simultaneous reduction in interstitial oxygen ions since the concentrations of these species are linked through the anion-Frenkel defect equilibrium. During oxygen diffusion the oxygen coordination of the transition metal ion could probably vary between 3 and 6 [7,8]. The possibility of changing coordination in a flexible manner may therefore be the key to understanding the high oxygen diffusivity in Sr 4 Fe 62y Co y O 136d . Previous reports claimed that the transport is by an interstitial mechanism [7,19], mainly because it has been observed that the ionic conductivity decreases when lowering the ambient oxygen partial pressure. Such reduction has been observed both in SrFeCo 0.5 O z and SrFe 1.2 Co 0.3 O z when the ambient gas is changed from air to Ar (see Table 1). The suggested thermal reduction and simultaneous formation of oxygen vacancies in Sr 4 Fe 62y Co y O 136d to explain the increased ionic conductivity with increasing Co content seems therefore to contradict the reduction in ionic conductivity with decreasing oxygen partial pressures (which is expected to promote vacancy formation). A tentative qualitative explanation may, however, be as follows: (i) Addition of Co to Sr 4 Fe 62y Co y O 136d results in a higher capacity of thermal reduction and therefore a higher concentration of vacancies (Eq. (18)). A higher vacancy concentration agrees with the reported reduction of d with increasing Co content [7]. This explains the higher ionic conductivity with increasing total content of Co. (ii) Addition of Co to Sr 4 Fe 62y Coy O 136d increases the p-conductivity due to thermal dis-

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proportionation of Co 31 and formation of polarons with higher mobility. When reducing the partial pressure of oxygen, the lower valence states becomes more stable and the electron hole concentration is reduced. This explains the decreasing p-conductivity with decreasing oxygen partial pressure. (iii) In the double layer of the Sr 4 Fe 62y Co y O 136d structure we suggest that Co 21 has a strong tendency of forming clusters with oxygen vacancies by electrostatic interaction. For instance, ¨ clusters may be formed by (in Kroger–Vink notation): Co 9B 1 V ??O 5 (Co B VO )? 2Co B9 1 V ??O 5 (2Co B VO ) Co 9B 1 2V O?? 5 (Co B 2VO )??? 2Co B9 1 2V O?? 5 (2Co B 2VO )??

[1:1] cluster, [2:1] cluster, [1:2] cluster, [2:2] cluster,

etc. When reducing the oxygen partial pressure the concentration of Co 21 increases, i.e. more clusters are formed. Under such conditions the defects become rapidly ordered in a similar manner as in transitions from perovskites to brownmillerites, which is also facilitated by reduced oxygen partial pressure [24]. Thus, one may interpret the ordered structure with lower oxygen coordination (referring to defect model (iii)) as an ordering of electrostatically stabilised cluster defects in the double layer. The consequence of these clusters, or ordering, on the oxygen ion diffusion is to reduce the concentration of mobile vacancies and their mobility. This explains the reduction in ionic conductivity with decreasing oxygen partial pressure. A strong ordering tendency may, furthermore, also explain the unusually rapid reduction in the ionic transference number with decreasing oxygen partial pressure [19]. The latter is explained by the two possible effects obtained: a reduction in the number of mobile vacancies and higher activation energy of diffusion due to the electrostatic interaction. As final comments, one may ask if the lower stability of Co-containing Sr 4 Fe 6 O 13 under reducing conditions and high temperature is a consequence of

long range ordering promoted by electrostatic interaction where Co 21 plays a decisive role. Furthermore, it should be mentioned that extended defect clusters have for a long time been recognised in the ¨ highly defective wustite phase [33], and similar defect ordering in other highly defective structures would not be unexpected. Finally, a valuable test of the effect of Co on the oxygen diffusion would be to measure the ionic transport number under reducing conditions in Co-free Sr 4 Fe 6 O 13 . One may, from the above discussion, expect less lowering of the ionic conductivity in such a material.

5. Concluding remarks The stability range of the layered perovskite related phase Sr 4 Fe 6 O 13 is highly dependent on additions of other elements. A relatively small addition of lanthanum results in formation of the perovskite and a nanoscale intergrowth of the perovskite and the layered perovskite related structure. The intergrowth demonstrates the close structural relationship between the two phases. With addition of Co, the stability of Sr 4 Fe 6 O 13 is also lowered, and this feature could be related to strong ordering of Co and oxygen vacancies in the double layer. An electrostatic interaction between Co 21 and V O?? has been put forward as a possible explanation for the ordering tendency. It has been suggested that such an ordering strongly affects the ionic diffusion, and explains the strong oxygen partial pressure dependency of the ionic conductivity. Based on a defect model dominated by an oxygen pressure independent oxygen point defect concentration, and with electronic minority defects, it has been suggested that the electron hole mobility follows a small polaron mechanism and that the ratio of the mobility of oxygen point defects to electron holes is unusually high.

Acknowledgements The authors are grateful for the preparation of Figs. 3–5 by Dr. Arve Holt with the program CrystalDesigner 6.0.3.

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