Mixed electronic and ionic conductivity of LaCo(M)O3 (M=Ga, Cr, Fe or Ni)

Solid State Ionics 104 (1997) 67–78

Mixed electronic and ionic conductivity of LaCo(M)O 3 (M 5 Ga, Cr, Fe or Ni) I. Oxygen transport in perovskites LaCoO 3 –LaGaO 3 V.V. Kharton*, A.P. Viskup, E.N. Naumovich, N.M. Lapchuk Institute of Physico-Chemical Problems, Belarus State University, 14 Leningradskaya Str., 220080, Minsk, Republic of Belarus Received 29 April 1997; accepted 9 July 1997

Abstract The formation of a continuous series of solid solutions with a rhombohedrally-distorted perovskite type structure has been found in the pseudobinary oxide system LaCoO 3 –LaGaO 3 . Substitution of gallium with cobalt in the lanthanum gallate results in blocking oxygen ionic transport and increasing electronic conductivity. The activation energy of the electrical conductivity of the LaGa 12x Co x O 3 ceramics (x 5 0.2–0.6) is in the range 57–65 kJ mol 21 . The thermal expansion coefficients increase regularly with the cobalt content and lie in the range from 11.2 3 10 26 to 22.3 3 10 26 K 21 . Doping LaCoO 3 with gallium has been ascertained to lead to lower electronic conductivity and oxygen permeability. The results of electrical conductivity and electron paramagnetic resonance (EPR) studies suggest that introduction of gallium into the cobalt sublattice leads to insulating cobalt ions. Keywords: Perovskite; Mixed conductor; Ionic conductivity; Oxygen permeability; Lanthanum cobaltite–gallate

1. Introduction Dense mixed ionic-electronic conducting oxide membranes are of considerable interest for industrial processes such as oxygen separation, waste reduction and recovery, coal gasification and selective oxidation of hydrocarbons [1–9]. The advantage of mixed conductive membranes lies in their infinite theoretical separation factor [1–5]. For oxidation processes, a promising feature is that the oxygen flux may alter *Corresponding author. Tel.: 1375 17 2207681; fax: 1357 17 2265567; e-mail: [email protected] or [email protected].

the relative presence of different active oxygen species on the membrane surface, thereby providing higher selectivity for partial oxidation reactions [6– 9]. Moreover, it is possible to omit the preliminary step of purification of oxygen from nitrogen at the oxidation [2]. One of the most promising groups of mixed conductors are the perovskite-like oxides having high ionic conductivity and a prevailing electronic conductivity [10–12]. So, ceramic membranes of perovskite-related oxide phases Sr(Co 12x Fe x )a Og (a 5 1– 20, x 5 0–1.0) and La(Sr)Co(Fe)O 32d , were tested successfully under various external conditions [3,6,11–14].

0167-2738 / 97 / $17.00  1997 Elsevier Science B.V. All rights reserved. PII S0167-2738( 97 )00397-4

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V.V. Kharton et al. / Solid State Ionics 104 (1997) 67 – 78

For industrial application of the mixed-conducting membranes, one must avoid a decrease in oxygen permeation flux with time, as has been mentioned in a number of works [14–17]. This degradation in behaviour was attributed to an ‘order-disorder’ phase transition in the oxygen sublattice [16,17]. The transition from a high-temperature structure in which the oxygen vacancies are disordered to a low temperature phase in which the vacancies are highly ordered, leads to reducing oxygen ion mobility. Such a transition, confirmed by detailed structural investigations [17,18], is characteristic of most perovskites derived from the strontium cobaltite SrCoO 32d [16]. Another important reason for membrane degradation is interaction of the oxide materials containing alkaline-earth elements with gas species such as CO 2 and SO 2 [3,12,19–21]. Owing to this reaction, the applicability of the perovskites containing alkalineearth metals is essentially limited. Moreover, a high oxygen chemical potential gradient in the Sr(Co, Fe)O x -type membranes results in a crystal lattice mismatch inside the ceramics, leading to fracture [21]. Thus, the disadvantages of SrCoO 32d -based ceramics create a necessity to develop new membrane materials for electrocatalytical oxidation reactors. Preferably, such materials should not contain alkaline earth elements and possess a higher stability in reducing atmospheres as compared to the strontium cobaltites-ferrites. Note that the stability range enlargement is associated, as a rule, with increasing energy of crystal lattice bonds [22]. So, an increase of the oxide phase stability may result in decreasing ionic conductivity and oxygen permeability of the mixed conductors. As LaCoO 3 exhibits significant ionic conductivity with prevailing electronic conductivity [23–25], a study of trends in the oxygen transport properties of oxides derived from lanthanum cobaltite is of interest in order to search for new membrane materials. Perovskite phases LaGaO 3 and LaCoO 3 are closely related structurally [26–28], which permits us to expect a high solid solubility of gallium in the cobalt sublattice of lanthanum cobaltite. Substitution of cobalt by gallium is assumed to increase the ionic character of the B–O bonds in the ABO 3 perovskite lattice. This can increase the stability of LaCoO 3 in reducing atmospheres but will decrease its electronic conductivity.

On the other hand, perovskite-type La(Sr)Ga(Mg)O 3 solid solutions are promising solid electrolytes which have a high oxygen ionic conductivity (so . 10 22 S cm 21 at 870 K) with an ionic transference number to t O ¯ 1 over the oxygen 220 partial pressure range 10 , pO 2 ,0.4 atm [27– 29]. Their advantages were reported to be a relatively low thermal expansion coefficient, close to that of yttria stabilized zirconia, and sufficient chemical stability [28]. Lanthanum–strontium cobaltite cathodes were tested successfully in solid oxide fuel cells (SOFCs) with La(Sr)Ga(Mg)O 3 electrolytes [30,31]. Obviously, the formation of LaGa(Co)O 3 solid solutions may occur in the diffusion layer between cobaltite electrodes and LaGaO 3 based solid electrolytes. The present work was aimed at studying perovskite-type oxides of the LaCoO 3 –LaGaO 3 pseudobinary system.

2. Experimental A standard ceramic synthesis route using highpurity La(NO 3 ) 3 ?6H 2 O, Co(NO 3 ) 2 ?6H 2 O and metallic gallium as starting materials was used for the preparation of LaGa 12x Co x O 3 (x50.2, 0.4, 0.6, 0.8 and 1.0) powders. The starting mixtures were initially dissolved in nitric acid and then dried. After thermal decomposition of the nitrates, solid-state reactions were conducted in air at temperatures of 1570 to 1770 K for 35–40 h with multiple repeated intermediate grindings. Ceramic specimens were pressed (200–600 MPa) into the shape of bars (43 4330 mm 3 ) and disks of various thickness (diameter 12 or 15 mm). Gas-tight ceramics were sintered in air at 1720–1840 K for 15–30 h. The formation of the perovskite phases in the ceramics prepared by both methods was verified by X-ray diffraction (XRD) technique. X-ray fluorescence analysis (XFA) and atomic emission spectroscopy (AES) were used to verify the cation composition of the powders and ceramics. The experimental procedures for XRD, AES and XFA studies, testing gas tightness, investigating electrical conductivity and thermal expansion were described in detail earlier [23,24,32–34]. The technique to determine oxygen ion transference numbers by the e.m.f. method was described in Ref. [32]. Only

V.V. Kharton et al. / Solid State Ionics 104 (1997) 67 – 78

specimens which had been verified as being gas tight, were used for measurements of electrical conductivity, oxygen permeability and ion transference numbers. The electron paramagnetic resonance (EPR)-spectra were recorded at room temperature using modulation-type spectrometers RadioPAN and Varian-E112 (frequency 9.3 GHz). As the VarianE112 spectrometer uses a reference ultra-high frequency and provides a linear signal within a large power range, it was used to perform the continuous saturation studies. The signal of Cr 31 -ions of a ruby standard sample located permanently in the cavity together with the specimen measured was used for quality factor control, for magnetic field phase fine tuning and for calibration of the H1 -component. Oxygen permeation fluxes through the oxide ceramics were determined using an electrochemical solid electrolyte cell equipped with an oxygen pump and a sensor (Fig. 1). The experimental technique for the permeation measurements was described previously [4,23,34]. The oxygen permeability was determined under the condition of equality of oxygen

69

flows leaving the cell through the pump and entering it through the hermetically sealed sample by ionic transport. The ceramics Zr 0.9 Y 0.1 O 1.95 , used in the measurement cells were manufactured by the Ukrainian Research Institute of Refractories (Kharkov, Ukraine). Only the perovskite ceramic samples which were gas-tight were used in the experiment. In order to avoid diffusional limitations which may appear when the inner volume of the cell contains a mixture of oxygen and nitrogen, the cell shown in Fig. 1 was blown through with oxygen before the sample was sealed. This was achieved by applying an electrical potential difference over the electrodes of the oxygen pump which resulted in oxygen entering the cell. The cell was heated up and the sample was sealed after the sensor e.m.f. had reached values corresponding to values of an oxygen pressure of 1 atm within the cell. As the thermal expansion coefficient (TEC) of zirconia differs significantly from that of the oxides studied, the measurements were carried out after slowly coolingdown the cell, a process which was started immediately after the melted sealant had bound the sample onto the cell. The oxygen permeability was measured at temperatures from 960 to 1200 K with a difference of the oxygen partial pressure in the internal and external spaces of the measuring cell between 2310 4 and 1310 3 Pa. Some values of the flow through the sample corresponding to different values of the oxygen chemical potential gradient were determined isothermally during each measurement. The time of attainment of steady state in the cell (when the sensor e.m.f. is independent of time) was 20 to 60 h. For the analysis of the oxygen permeation processes, we use quantities of oxygen permeation flux density j chem (mol s 21 cm 22 ) and specific oxygen permeability J(O 2 ) (mol s 21 cm 21 ). The values of J(O 2 ) were calculated by following formula [35]: RTd j chem J(O 2 ) 5 ]] ? ]] 4F E

(1)

where d is the thickness of the sample, E is the e.m.f. of the electrochemical oxygen sensor: Fig. 1. Schematic drawing of the electrochemical cell for the oxygen permeability measurements: (1)-YSZ solid electrolyte; (2)-electrodes of the oxygen sensor; (3)-electrodes of the oxygen pump: (4)-membrane measured; (5)-high-temperature glass; (6)high-porosity ceramic insertions; (7)-thermocouple.

S D

p2 RT E 5 ] ? ln ] 4F p1

(2)

with p2 and p1 the oxygen partial pressures at the membrane feed and permeate sides ( p2 $p1 ), respec-

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V.V. Kharton et al. / Solid State Ionics 104 (1997) 67 – 78

tively. For all results presented in this paper, the value of p2 was maintained at 0.21310 5 Pa (atmospheric air). The ceramic membrane thickness (d) was varied from 0.40 to 1.40 mm. The dependence of the oxygen permeation flux on the membrane thickness is discussed below. The physical meaning of J(O 2 ) is the factor of proportionality between the oxygen flux density passing through the membrane and the difference of the chemical potentials of oxygen at opposite sides of the specimen. ln case the ionic and electronic conductivity of the oxide does not depend on the chemical potential of oxygen and there are no surface limitations of the oxygen flux, the value of the oxygen permeability may be represented as follows [35]: RT J(O 2 ) 5 ]]2 ? t O (1 2 t O ) ? s 16F

(3)

where t O is the oxygen ion transference number, and s is the specific electrical conductivity of the oxide. Oxygen transport through lanthanum cobaltite is limited mainly by the bulk ionic conductivity which depends on the oxygen partial pressure [23,25]. So, experimental values of j chem and J(O 2 ) are given hereafter in combination with p1 or E. In order to investigate the oxygen ion transference processes, we also used Tubandt’s method which is based on a comparison between the total electrical current passing through an ionic conductor and a quantity of the substance transferred through a galvanic cell [36]. We modified this method to apply it for the oxygen ion-conducting oxide studies. A schematic drawing of the measuring cell is presented in Fig. 2. The cell consists of a zirconia solid electrolyte oxygen pump, a sensor, and the oxide membrane under investigation. Platinum electrodes with a loading density of 20–25 mg cm 22 were applied to both membrane surfaces. Before the investigations, the cell was blown through with oxygen analogously to the case described above. In the course of the measurement, an electrical current (Iin ) was passed through the membrane to pump oxygen into the cell, and the solid electrolyte pump removed oxygen from the cell. The measurements were performed at steady state which corresponds to the condition:

Fig. 2. Schematic drawing of the electrochemical cell for the oxygen ion transference number measurements: (1)-YSZ solid electrolyte; (2)-electrodes of the oxygen sensor; (3)-electrodes of the oxygen pump; (4)-ceramic membrane measured; (5)-hightemperature glass; (6)-high-porosity ceramic insertions; (7)-thermocouple, (8)-electrodes of the membrane.

4FS ? ( j el 1 j chem ) 5 Iout

(4)

where S is the membrane surface area, Iout is the electrical current passed through the solid electrolyte oxygen pump, j el and j chem are the densities of oxygen fluxes entering the cell through the membrane due to the electrical potential difference and due to the difference of the oxygen chemical potentials at the membrane electrodes, respectively. The oxygen flux density j el is defined as follows [37]: t O ? Iin j el 5 ]]. 4FS

(5)

In case the difference of the oxygen chemical

V.V. Kharton et al. / Solid State Ionics 104 (1997) 67 – 78

potentials at the membrane electrodes is insignificant, i.e., E ¯ 0,

temperature for 3–6 months. Thus, the perovskites LaCoO 3 and LaGaO 3 were found to form solid solutions. The electrical conductivity of the La(Ga,Co)O 3 solid solutions increases steadily with cobalt content (Fig. 3). For the ceramics with x50.2–0.6, the temperature dependence of the conductivity, expressed in the Arrhenius coordinates, is close to linear. In order to calculate activation energy for electrical conductivity, we used two standard models (for instance, see Refs. [22,36,38]):

(6)

and, correspondingly, j chem ¯ 0.

(7)

the oxygen ion transference number can be measured directly: Iout t O 5 ]. Iin

71

(8)

The case in which the oxygen chemical potentials at the electrodes differ considerably from one another is analyzed below. The time of steady state attainment (when the sensor e.m.f. does not depend on time) was 20 to 60 h.

3. Results and discussion XRD studies of LaGa 12x Co x O 3 showed the formation of a continuous series of solid solutions with a rhombohedrally distorted perovskite structure which is characteristic of lanthanum cobaltite [26]. All powders and ceramics of LaGa 12x Co x O 3 were single phase in the range 0.2#x#1.0. The parameters of the perovskite lattice with rhombohedral distortion (a and a ) are presented in Table 1. The data on LaCoO 3 have been reported earlier [23]. The rhombohedral distortion of the La(Ga,Co)O 3 solid solution unit cells decreases with increasing cobalt concentration. No significant changes in the XRD patterns were observed after storage in air at room

Fig. 3. Temperature dependence of the electrical conductivity of the LaGa 12x Co x O 3 ceramics in air: (1) x50.2; (2) x50.4; (3) x50.6; (4) x50.8.

Table 1 Properties a of the LaGa 12x Co x O 3 ceramics x

rexp , kg m 23

Parameters of the unit cell

Mean values of thermal expansion coefficient

a, nm

a, 8

T, K

a¯ ?10 6 , K 21

370–650 650–1070 300–1150 430–680 680–1100 300–1130

11.260.3 15.860.4 17.760.5 18.760.5 22.360.3 22.260.3

0.2

6380

0.5383

60.75

0.4 0.6

6670 6780

0.5402 0.5423

60.70 60.64

0.8

6940

0.5447

60.62

a

a and a are the parameters of the perovskite-type lattice with a rhombohedral distortion. rexp is the density of the ceramics. a¯ is the thermal expansion coefficient calculated from dilatometric data and averaged in the given temperature range.

V.V. Kharton et al. / Solid State Ionics 104 (1997) 67 – 78

72

F G A E s 5 ] ? exp F 2 ] G , T RT s 5 A 0 ? exp

Ea 2] , RT

0

(9)

a

(10)

where s is the specific electrical conductivity, A 0 is the pre-exponential factor, and Ea is the activation energy. Eq. (9) and Eq. (10) are commonly used and may be applied to numerous thermally-activated transport processes such as ionic transport or electron conduction via a small polaron mechanism [22,36,38]. Regression parameters obtained by fitting of the experimental data are presented in Table 2. The adequacy of both models is approximately the same. For LaGa 0.2 Co 0.8 O 3 , the error of the calculations is larger in comparison with other materials due to the non-linear dependence of the logarithm of the conductivity versus the reciprocal temperature (Fig. 3). The activation energy values for the cobaltites–gallates at x50.2–0.6 are close to each other and lie in the range 57–65 kJ mol 21 . Closely related values of Ea suggest that the conductivity mechanism for La(Ga,Co)O 3 at high gallium content is the same, independent of composition. The electrical properties of the perovskites leads to the conclusion that gallium cations block electronic conduction. The electronic conductivity of LaCoO 3 was reported to occur by a transfer of charge carriers via Co–O–Co bonds [26,38,39]. Excitation of electrons from a narrow valence band to localized t 2b states at high-spin cobalt sites introduces mobile small-polaron holes and trapped electrons at station-

ary Co 21 ions in the temperature range 110 K,T , 350 K [38]. At higher temperatures, a metallic phase containing high-spin Co 31 and intermediate-spin Co(III) ions is stabilized, and the partially filled s a* band is responsible for the p-type conductivity above 650 K [38]. As gallium ions have only one oxidation state, the introduction of gallium into the cobalt sublattice results in a decreasing electronic charge carrier concentration and insulating cobalt ions which provide the electronic conduction. Thus, the most credible mechanism of electronic conduction at high gallium content (x#0.6) is electron hole hopping between cobalt cations where the holes form due to a disproportionation of the trivalent cobalt ions (Co 31 and Co(III)) [25,38]. At x.0.6, the nature of the electronic transfer is closely related to that of LaCoO 3 described in Refs. [26,38]. The conductivity of La(Ga,Co)O 3 was shown to decrease with reducing oxygen partial pressure (Fig. 4). Such a decrease, typical for LaCoO 3 [23], is caused by the release of oxygen from the crystal lattice, which leads to a decrease in the electron hole concentration due to the displacement of the cobalt ion charge disproportionation equilibrium [25]. The EPR spectra of the La(Ga,Co)O 3 solid solutions are presented in Fig. 5. For LaG 0.2 Co 0.8 O 3 and LaGa 0.4 Co 0.6 C 3 , the high electrical conductivity causes a poor quality factor-the standard signal magnitude is relatively low. Only a weak peak with g51.956760.0002 is observed in the EPR spectrum of LaGa 0.2 Co 0.8 O 3 . For the oxide with x50.6, one can detect a number of low-intensity signals which correspond to values of the g-factor being approximately equal to 6.0, 4.3, 2.8 and 1.95 (Table 3). The

Table 2 Regression parameters a of the temperature dependence of the electrical conductivity of LaGa 12x Co x O 3 ceramics in air x

T, K

Model

Ea , kJ mol 21

ln(A 0 ), S cm 21

r

0.2

570–1250

0.4

470–1260

0.6

580–1150

0.8

500–1200

Eq. (9) Eq. (10) Eq. (9) Eq. (10) Eq. (9) Eq. (10) Eq. (9) Eq. (10)

5764 6464 5761 6364 5965 6565 3964 4564

9.060.6 16.860.6 11.560.8 19.160.2 14.260.8 21.960.7 15.060.7 22.760.7

0.993 0.994 0.9995 0.9993 0.995 0.996 0.98 0.990

a

Ea is the activation energy for electrical conductivity. r is the correlation coefficient.

V.V. Kharton et al. / Solid State Ionics 104 (1997) 67 – 78

73

Table 3 Selected EPR signals of the La(Ga,Co)O 3 , solid solutions x

g-factor 60.0002

Intensity A?DH 2 /m, relative units

Signal width DH, mT

0.2

5.9930 4.3145 1.9558 5.9731 4.2956 2.1771 6.0022 4.3475 2.1906 1.9567

51.64 3.34 274 29.31 6.06 52.63 4.22 1.76 32.39 14.05

30.10 6.02 36.15 28.52 8.44 35.0 20.49 13.26 31.33 26.5

0.4

0.6

0.8

A is the signal amplitude DH is the line width. m is the specimen weight.

Fig. 4. Dependence of the electrical conductivity of the LaGa 12x Co x O 3 ceramics on the oxygen partial pressure at 115363 K: (A) x50.2; (B) x50.4; (C) x50.6; (D) x50.8.

Fig. 5. EPR spectra of LaGa 12x Co x O 3 ceramics at room temperature: (1), x50.2; (2), x50.4; (3), x50.6; (4), x50.8.

magnitude of these lines is comparable to the spectrometer sensitivity (10 12 spin cm 23 ). Such signals are also characteristic for the solid solutions with x50.4 and 0.2, except for the one with a g-factor of 2.8 (LaGa 0.8 Co 0.2 O 3 ). The resolution of the EPR spectra increases with increasing gallium content due to the decrease in conductivity. A signal with an anomalously high value of g¯11 is observed in the spectra of LaGa 0.6 C 0.4 O 3 and LaGa 0.8 Co 0.2 O 3 . The position of the zero line in the EPR spectra of La(Ga,Co)O 3 indicates the existence of a very wide line typical for ferromagnetic or antiferromagnetic systems. This is also corroborated by the relatively large linewidth of the EPR signals. An increase of the gallium concentration leads to increasing intensity of the EPR signal with gi 5 1.957660.0002 and g' 51.956160.0002. Therewith, the width of this line increases from 26.50 to 36.1560.02 mT (Table 3). The signal is assigned to Ga 31 cations which have a 3d 10 configuration ( 1 So state). So, the g-factor of ions in the S-state is close to that of the free electron spin ( ge ). The deviations from the free electron spin g-factor value (Dg5ge 2 gobs ) are caused by excited state impurities. The Dg quantity depends on the crystal field quantity and can be both positive and negative [40]. In case of the LaGa 12x Co x O 3 ceramics, the values of Dg are 10.0456, 20.1883, 20.1748 and 10.0465 for x5 0.8, 0.6, 0.4 and 0.2, respectively. According to the literature data [41–43], the EPR signals with g54.3 and 6.0 are attributed to trivalent

74

V.V. Kharton et al. / Solid State Ionics 104 (1997) 67 – 78

and bivalent cobalt ions respectively. The ratio between the total cobalt concentration and the gallium concentration ([Co] / [Ga]) was estimated from the intensities of the EPR lines to be 0.67 for LaGa 0.6 Co 0.4 O 3 and 0.20 for LaGa 0.8 Co 0.2 O 3 (Table 3). Taking into consideration a standard error in the concentration measurement by EPR which is about 30%, one can conclude that the assignment of the signals is adequate. The trivalent cobalt ion signal width decreases with x in the EPR spectra (Table 3). One should also note that the width of this line is low in comparison with those of other paramagnetic centres. Such behaviour attests to a relative insulated state of the cobalt ions in respect to the interaction with the surroundings. Fig. 6 shows the EPR signal amplitude as a function of the microwave field power for the LaGa 0.8 Co 0.2 O 3 ceramics. No saturation was observed when the field power increases. Such behaviour is characteristic of the spin–lattice relaxation of broken covalent bonds in the amorphous semiconductors such as a-Si and a-Ge [44]. The calculated spin-lattice relaxation time values are 2310 210 to 1310 29 s. Notice that all the EPR signals which

relate to the different paramagnetic centres demonstrate the same behaviour at changing magnetic field power. This fact indicates an identical spin system for the all-type centres. The spin–spin interaction results in additional widening of the lines. The mean thermal expansion coefficients (TECs) calculated from the dilatometric data on the LaGa 12x Co x O 3 ceramics increase regularly with increasing cobalt concentration (Table 1). The values of TEC are in the range of 11.2310 26 to 22.3310 26 K 21 . Temperature dependencies of the relative elongation are presented in Fig. 7. For ceramics with x50.2 and 0.6, the dilatometric curves can be approximated by two straight segments with a break at 650–680 K. At high temperatures, the thermal expansion of these solid solutions features a noticeable increase. The TEC of LaGa 0.2 Co 0.8 O 3 is close to that of lanthanum cobaltite which was reported to be equal to (22.960.3)310 26 K 21 at temperatures of 300 to 1100 K [23]. The results of the oxygen permeability measurements are listed in Table 4. The quantity of J(O 2 ) which is proportional to j 3d (see Eq. (1)), should be independent of thickness if there are no surface limitations to the oxygen permeation flux [3,12]. In

Fig. 6. Dependence of the EPR signal intensity on the microwave field power for LaGa 0.8 Co 0.2 O 3 : (1) the signal with g-factor of 6.0; (2) the signal with g-factor of 4.3; (3) the signal with g-factor of 1.95. H10 is the value of the H1 -component of the field at 0 dB.

V.V. Kharton et al. / Solid State Ionics 104 (1997) 67 – 78

75

Fig. 7. Temperature dependence of the relative elongation of LaGa 12x Co x O 3 ceramics in air: (1) x50.2; (2) x50.4; (3) x50.6; (4) x50.8.

Table 4 Parameters of oxygen permeation through LaGa 12x Co x O 3 membranes at 115564 K ( p2 50.21310 5 Pa) x

d, mm

E, mV

j chem 310 9 , mol s 21 cm 22

J(O 2 ) 10 11 , mol s 21 cm 21

0.8

0.98

0.6

0.45

170 730 30.1 53.6 102 147 190 151 145 189 731 227 590 717

4.6 9.7 5.1 6.9 9.3 9.7 9.9 4.7 2.9 3.1 6.9 1.2 3.0 3.3

6.7 3.3 18.9 14.3 10.2 7.4 5.8 7.7 7.1 5.7 3.3 1.8 1.8 1.6

1.00 1.40

0.4

1.40

cases where interphase oxygen exchange limitations are considerable, J(O 2 ) should increase with increasing d due to the reducing role of the surface limitation at a given oxygen chemical potential gradient. For the LaGa 12x Co x O 3 ceramic membranes, the thickness dependence of the oxygen permeability is insignificant in the range d50.4–1.4

mm (Table 4). The variation of the J(O 2 ) values for membranes of various thickness are within the limits of experimental error. This behaviour suggests an absence of surface exchange limitations which is in excellent agreement with the permeation data on LaCoO 3 [25]. The oxygen flux through LaGa 0.2 Co 0.8 O 3 is close to that of LaGa 0.4 Co 0.6 O 3 . A further decrease of the cobalt content leads to a decreasing oxygen permeation flux. One should note that the oxygen permeability of all the La(Ga,Co)O 3 solid solutions studied is lower than that of lanthanum cobaltite [23]. The oxygen ion transference number measurements using the cell shown in Fig. 2 were performed under the condition E¯0, and also at significant gradients of the oxygen chemical potential. In the first case, Eq. (8) was used to calculate the ion transference number. In the cases where E .0, the steady oxygen flux pumped from the cell ( j out ) can be described using Eq. (4): Iout t O ? Iin j out 5 ]] 5 ]] 1 j chem . 4FS 4FS

(11)

As the surface exchange limitations of the permeation process are negligible, one can express the permeation flux density by the equation [5,12]:

V.V. Kharton et al. / Solid State Ionics 104 (1997) 67 – 78

76 p2

E

sO ? se RT ]]] j chem 5 ]] ≠(ln p O ) 2 ? s 2 16F d p O 2 se

(12)

Table 5 Oxygen ion transference numbers a of LaGa 12x Co x O 3 ceramics at 115563 K ( p2 50.21310 5 Pa)

1

where se is the electronic conductivity. In order to simplify Eq. (12), we used the standard assumption that electrical conductivity, as well as the transference numbers, are independent of the oxygen chemical potential [35]: RTs ? t O (1 2 t O ) p2 j chem 5 ]]]]] ? ln ]. p1 16F 2 d

(13)

Such an assumption is valid only in a very narrow range of oxygen partial pressures [36]. Substituting Eqs. (3), (5) and (13) into Eq. (4), one can obtain

S

D

E E ] ? t 2O 2 ] 1 Iin ? t O 1 Iout 5 0 R R

(14)

and, correspondingly, ]]]]]]] E 1 Iin ? R 2 E ? Iout 1 ] (E 1 Iin ? R) 2 R ? ]]] 2 ]] 2 2R R t O 5 ]]]]]]]]]]]]] E

œS

D

(15) where R is the resistance of the membrane studied. For each oxygen partial pressure difference, the value of R was measured independently using AC (frequency 10 and 20 kHz) and calculated from the relation: U R5] Iin

(16)

where U is the potential difference between the membrane electrodes. The difference of the R values obtained using DC and AC was less than 5%. Table 5 presents values of the oxygen ion transference number for the LaGa 0.6 Co 0.4 O 3 and LaGa 0.8 Co 0.2 O 3 ceramics. Calculation using Eq. (15) permits the obtainment of quantities which are sufficiently close to those measured at zero oxygen chemical potential gradient. The difference between

x

Formula used for the calculation

E, mV

tO

so , S cm 21

0.2

Eq. Eq. Eq. Eq. Eq. Eq. Eq.

,2 17.3 40.3 ,2 19.7 40.8 75.9

0.021 0.020 0.011 0.0012 0.0018 0.0037 0.0063

0.0049

0.4

(8) (15) (15) (8) (15) (15) (15)

0.0034

a

Measurements were performed using the cell shown in Fig. 2. so is the oxygen ionic conductivity in air.

these values may be explained by the dependence of the oxygen ionic and electronic conductivities on the oxygen pressure [23,25]. The ionic conductivity of LaCoO 3 has been reported to increase with decreasing oxygen partial pressure [25]. Both La(Ga,Co)O 3 solid solutions studied are predominantly electronic conductors (Table 5). This fact is confirmed also by the e.m.f. measurements. Studies by the e.m.f. method were performed using the oxygen concentration cell [32]. Fluxes of atmospheric air ( pO 2 52.1310 4 Pa), nitrogen ( pO 2 ¯20 Pa) or oxygen (1310 5 Pa) were supplied continuously to the electrodes. For the LaGa 0.6 Co 0.4 O 3 ceramics, the e.m.f. did not exceed 1 mV for all oxygen partial pressure differences at 800–1200 K. For LaGa 0.8 Co 0.2 O 3 , the e.m.f. maximum was approximately 3 mV in the cell with a nitrogen / air gradient. The e.m.f. values were excessively low which does not allow the estimation of transference numbers quantitatively. One can deduce, however, that the oxygen ion transference numbers are about 0.02 for LaGa 0.8 Co 0.2 O 3 and are much less than experimental error of the e.m.f. method for LaGa 0.6 Co 0.4 O 3 . The values of the oxygen ionic conductivity of LaGa 0.6 Co 0.4 O 3 and LaGa 0.8 Co 0.2 O 3 are close to one another (Table 5). The ionic conductivity of the La(Ga,Co)O 3 solid solutions is substantially lower than that of LaGaO 3 [27]. Thus, substitution of gallium with cobalt results in a decreasing oxygen ionic conductivity and an increasing electronic conduction.

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4. Conclusion The perovskite-type phases LaCoO 3 , and LaGaO 3 have been found to be single-phase solid solutions. The crystal structure of the LaGa 12x Co x O 3 solid solutions at 0.2#x#1.0 is characterized as rhombohedrally-distorted perovskite. Doping LaCoO 3 with gallium leads to decreasing electrical conductivity and oxygen permeability. Considering the results of the electrical conductivity and EPR studies, one can conclude that introduction of gallium into the cobalt sublattice of LaCoO 3 leads to insulating cobalt ions. The averaged thermal expansion coefficients increase regularly with increasing cobalt concentration and lie in the range of 11.2310 26 to 22.3310 26 K 21 . The activation energy for the electrical conductivity of the LaGa 12x Co x O 3 ceramics (x50.2–0.6) in air falls in the range 57–65 kJ mol 21 . Substitution of gallium by cobalt in the lanthanum gallate has been ascertained to result in blocking ionic conduction and increasing electronic conductivity.

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