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Structure and properties of nonstoichiometric mixed perovskites A3B′1+xB″2−xO9−δ
Solid State Ionics 154 – 155 (2002) 659 – 667 www.elsevier.com/locate/ssi
Structure and properties of nonstoichiometric mixed perovskites A3BV 1+xBW 2xO9d Shanwen Tao, John T.S. Irvine * School of Chemistry, University of St. Andrews, St. Andrews, Fife KY16 9ST, Scotland, UK Accepted 14 March 2002
Abstract Nonstoichiometric mixed perovskites A3BV 1 + xBW 2 xO9 d, e.g. Ba3Ca1.18Nb1.82O9 d, exhibit high proton and oxygen-ion conductivity. It is expected that mixed ionic and electronic conductors may be found in these compounds if the B-sites are partially substituted by a first row transition element. These mixed conductors may be potential anode materials for fuel cell applications. The structure of single phase SrCu0.4Nb0.6O2.9 was studied by both X-ray and neutron diffraction. It is tetragonal ˚ , c = 3.9757(2) A ˚ , V = 62.37(2) A ˚ 3 according to neutron diffraction. Rietveld with space group P4/mmm (123), a = 3.9608(4) A refinement indicates that the oxygen vacancy tends to stay at O1 (1c) site with O2 (2e) fully occupied. AC impedance measurements indicate that electronic conduction is probably dominant in air. The DC conductivity of SrCu0.4Nb0.6O2.9 at pO2 in the range of 10 22 – 10 12 atm exhibits a p(O2) 1/4 dependence consistent with n-type electronic conduction. The material was unstable in 5% H2 at elevated temperatures but stable in argon at 900 jC. Using manganese instead of copper, a phase that is redox stable was prepared. SrMn0.4Nb0.6O3 d exhibits an orthorhombic structure with space group Pbnm (62), ˚ , b = 5.6589(2) A ˚ , c = 7.9729(2) A ˚ , V = 254.69(7) A ˚ 3 according to X-ray diffraction. Such a unit cell indicates a = 5.6451(3) A that it is a double perovskite and therefore the formula is better written as Sr2Mn0.8Nb1.2O6 d. The material maintains perovskite structure in 5% H2 although thermal expansion was observed on reduction. The conductivity of Sr2Mn0.8Nb1.2O6 is 0.36 S/cm in air at 900 jC. Conductivity decreases in 5% H2 indicates p-type conduction at low pO2. D 2002 Elsevier Science B.V. All rights reserved. Keywords: Electrical conductivity; Perovskite structure; Neutron diffraction; Stability; Fuel cell; Defect
1. Introduction High temperature proton conductors have been of great interest for their potential applications in solid oxide fuel cells, sensors and steam electrolyzers since their discovery, primarily by Iwahara et al. [1] and Uchida et al. [2]. These rare earth elements-doped ABO3 perovskite oxides with A = Ca, Sr, Ba, B = Ce, *
Corresponding author. Tel.: +44-1334-463817; fax: +441334-463808. E-mail address: [email protected] (J.T.S. Irvine).
Zr and Ti, etc. have been studied in detail [3 –6]. An interesting discovery was that some nonstoichiometric mixed perovskites with formula A3BV 1 + xBW 2 xO9 d where A is an alkali earth element Ca, Sr, Ba, etc., BVis an element with valence + 2 or + 3 and BW is an element with valence + 5, such as Nb and Ta exhibit quite high proton and oxygen-ion conductivity [7 –12]. It is expected that mixed ionic and electronic conductors may be found in these compounds if the B-site is partially substituted by appropriate first row transition elements. Such mixed conductors may provide potential anode materials for fuel cell applications.
0167-2738/02/$ - see front matter D 2002 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 7 - 2 7 3 8 ( 0 2 ) 0 0 5 1 6 - 7
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In a reducing atmosphere, a transition elementcontaining perovskite oxide loses some lattice oxygen due to the reduction of transition elements coupled with the constraint of electroneutrality, SS OxO ¼ 1=2O2 þ VO þ 2eV ð1Þ The released electrons may result in the increase of ntype conductivity or decrease of p-type conductivity. The oxygen vacancies created may be used for oxygen transport leading to enhanced oxygen-ion conductivity. They may also combine with water vapour resulting in the formation of proton defects. SS S VO þ H2 OðgÞ þ OxO ¼ 2OHO
ð2Þ
Therefore, the first row transition element substituted A3BV 1 + xBW 2 xO9 d perovskites may be potential oxygen-ion (proton) and electronic mixed conductors. Perovskite-type oxides with general formula A3 BV 1 + xBW 2 xO9 d with A = Ca, Sr, Ba, etc., BV= Ca, Sr, Cr, Mn, Cu, Fe, Co, Ni, etc., BW = Nb, Ta have been synthesised and widely reported [13 – 22]. A wide range of structures from high cubic symmetry to the lowest triclinic symmetry has been observed. For example, Sr2CrNbO6 is a double perovskite with alternate B-site ordering. Ca3CaNb2O9 prepared under different conditions gives 1:1, 1:2 and 1:3 orderings, respectively [23]. The ordering of B-site cations and tilting of BO6 octahedral make the real structure of these perovskite oxides very flexible. Splitting of the B-site is commonly observed because two (or more) different elements BVand BW are involved. For example, in A3BVBBW 2O9, 1:1 ordering often occur if the Bsite splits with one position fully occupied by BW and the other one shared by 2/3BV and 1/3BW. However, sometimes this ordering is not complete. Partial ordering can occur if the two B-sites are still occupied by both BVand BW but with different cation ratios. Similarly, the B-site ordering may be 1:1, 1:2 or 1:3 depending on the elements involved and the synthesis conditions. The formulae of complex perovskite oxides with 1:1 and 1:2 B-site ordering are normally written as A2BVBWO6 for double perovskites and A3BVBW 2O9 for triple perovskites. Most compositions of the studied oxides have the formula of A3BVBW 2O 9 where theoretically, no oxygen deficiency exists according to electroneutrality. Defect concentrations and thus defect chemistry in perovskites can be
controlled and tailored significantly by doping. Departure from the composition of a crystallographically perfect solid by the substitution of small concentrations of similar sized hetero-valent ions has an immense impact on properties. By the introduction of di- or trivalent transition elements onto the BV-sites, materials possessing mixed ionic and electronic conductivity may be obtained. It was recently reported that the cermet Cu-ceria may be used as a SOFC anode material for direct oxidation of hydrocarbons without carbon deposition, indicating that copper-containing materials might be suitable catalysts for the complete oxidation of hydrocarbon in fuel cell operation [24]. It would be interesting if copper can be introduced into a complex perovskite oxide retaining both high ionic and electronic conductivity. The stability of a Cu-containing oxide in a reducing atmosphere with pO2 as low as 10 15 – 10 20 atm for SOFCs is a potential problem. However, copper metal may be homogeneously dispersed in the mixture if obtained directly from the reduction of Cu-containing complex perovskite oxides provided that sintering of the copper can be avoided. From this point of view, our first target was to synthesise the nonstoichiometric oxide Sr3Cu1 + x Nb2 xO9 d. The structure, conductivity and stability of SrCu0.4Nb0.6O2.9 were studied by X-ray diffraction, neutron diffraction, AC impedance, TGA, etc. It was found that the chemical stability of SrCu0.4Nb0.6 O2.9 at low pO2 is not good. The stability is improved when copper is replaced by the early first row transition element manganese. The structure and electrical properties of Sr2Mn0.8Nb1.2O6 are also reported.
2. Experimental 2.1. Powder synthesis The nonstoichiometric mixed perovskites were synthesised by solid state reaction using SrCO3, Cr 2 O 3 , MnO 2 , Fe 2 O 3 , CoCO 3 , NiO, CuO and Nb2O5 as starting materials. The carbonates and MnO2 were dried at 100 – 300 jC, oxides except MnO2 at 300 –800 jC overnight to remove absorbed H2O and CO2. Stoichiometric amounts were mixed and ball milled in zirconia containers with zirconia balls for two 15-min periods. The final firing temper-
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atures are between 1000 and 1500 jC according to the chemical composition. The mixture for Sr3Cu1 + x Nb2 xO9 d was calcined at 1000 – 1300 jC. 2.2. Sample characterisation XRD analysis of powders reacted at different temperatures was carried out on a Stoe Stadi-P diffractometer to determine phase purity and measure crystal parameters. Electron diffraction was carried out on a JEOL 2011 HRTEM operating at 200 keV. Powder neutron diffraction data were collected on the diffractometer HRPD, ISIS, Rutherford Appleton Laboratory. Structure refinement was performed by the Rietveld method using the program General Structure Analysis System (GSAS) [25]. Thermal analysis of SrCu0.4Nb0.6O2.9 was carried out on a Rheometric Scientific TG 1000M+ TA Instruments SDT2960 from room temperature to 950 jC (5 jC/ min), holding at 950 jC for 30 min then cooling down to 30 jC (10 jC/min) under flowing 5% H2 in argon at a rate of 35 ml/min. 2.3. Conductivity measurements For AC impedance measurements in air, a Schlumberger Solartron 1260 Frequency Response Analyser coupled with a 1287 Electrochemical Interface controlled by Zplot electrochemical impedance software was used over the frequency range 1– 100 mHz. AC impedance measurements were made in 50 jC steps in air between 600 and 900 jC. The DC conductivity was measured by a conventional four-terminal method using a Keithley 220 Programmable Current Source to control current and a Schlumberger Solartron 7150 Digital Multimeter to measure the voltage. The SrCu0.4Nb0.6O2.9 samples were mounted with four Pt wire electrodes to measure the DC conductivity dependence on pO2, which was monitored by a zirconia oxygen sensor.
3. Results and discussion 3.1. Structure of SrCu0.4Nb0.6O2.9 For the SrM0.5Nb0.5O3 y system with M = Cr, Mn, Fe, Co. Ni, Cu, a single phase was achieved for all but
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M = Ni, Cu. The second phase for nominal formula Sr3Cu1.5Nb1.5O9 d is SrCuO2 (JCPDS No. 39-1492). In order to get a single perovskite, a series of Sr3 Cu1 + xNb2 xO9 d (x V 0.5) was prepared by solid state reaction as described above. After calcining the ball-milled mixture of SrCO3, CuO and Nb2O5 at 1000 jC, the main phase for Sr3CuNb2O 9 is a perovskite, but not well crystallised. It also contains some un-indexed peaks, which may belong to second phases. Further heating the oxide at 1300 jC, the main phase is well crystallised with a weak impurity peak at 2h c 30j, which may be due to traces of the phase Sr5Nb4O15 (JCPDS Card No. 28-1248). The main phase is a copper-rich strontium niobate with 1:1 Bsite ordering along the [111]c* zone according to electron diffraction. The main phase a ffiffitetragonal pffiffiffi is p ffi perovskite with lattice parameter 2ap 2ap 2ap. Fesenko et al [26] reported that a second tetragonal phase was detected when calcining the starting materials Sr2CuO3 and SrNb2O6 at high temperatures (1270 K, 12 h) to form the perovskite phase Sr3CuNb2O9. Ba3CuNb2O9 was not obtained either when prepared by solid state reaction using BaCO3, CuO and Nb2O5 as precursors [18]. Phase purity was only found with x = 0.2 in the Sr 3 Cu 1 + x Nb 2 x O 9 d (0 V x V 1) system. With x < 0.2, second phases were evident, e.g. Sr5Nb4O15 and second perovskite phases. When x>0.2, the second phase SrCuO2 was identified. No super structure was observed by electron diffraction indicating a simple perovskite. The accurate formula for Sr3Cu1.2Nb1.8O8.7 is SrCu0.4Nb0.6O2.9. Fig. 1 shows the electron diffraction of a cubic primitive cell indicating no super structure along [100] c* and [010]c* zones. The lattice parameters for a and b appear to be equal. No symmetrical absences were observed in the XRD pattern indicating a primitive cubic or tetragonal structure with slight distortion. Space group Pm3¯m and P4/mmm were chosen for Rietveld refinement of X-ray and neutron diffraction data. From the refinement of XRD data, the cubic Pm3¯m model gave smaller v2 value and R-factors indicating it might be a cubic perovskite; however, there was visible extra broadening on peaks that would be expected to split under tetragonal distortion. Clearly, this preference for the cubic space group was an artefact due to the limited resolution of the Xdiffractometer, even though a high resolution labora-
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Fig. 1. Electron diffraction pattern of SrCu0.4Nb0.6O2.9. The electron beam is parallel to [001].
tory system was used. Therefore, the real symmetry for SrCu0.4Nb0.6O2.9 should be lower than cubic. In addition, X-ray diffraction is not so sensitive to the light oxygen atoms making it difficult to determine the exact oxygen positions. The problem may be solved by neutron diffraction because the neutron
scattering length of oxygen is slightly lower than strontium, copper and niobium. It is believed that neutron diffraction may precisely determine the oxygen atom positions; however, it should be noted that neutron diffraction does not discriminate so well between the cations. The final refined occupancies of Sr, Cu and Nb from XRD are very close to the ratio in the starting formula although the oxygen occupancy cannot be refined from the XRD data. The occupancies for metal elements from XRD data were introduced and held constant in the refinement of the neutron data. The tetragonal model P4/mmm gives a smaller v2 value (1.590) and R-factor than those for the cubic model Pm3¯m (v2 = 1.953). The refined oxygen stoichiometry is consistent with the formula SrCu0.4Nb0.6O2.9. Fig. 2 shows the neutron patterns of SrCu0.4Nb0.6O2.9. The final refined structure data are given in Table 1. During the refinement, it was found that O1 tends to be partially occupied indicating vacancies may preferentially locate on this site as shown in Fig. 3. There is a degree of ordering of the oxygen vacancies at room temperature. Such ordering may limit the diffusion of charge carriers SS VO resulting in lower oxygen-ion conductivity. However, at high temperatures, this situation may change. The bond distance between Cu/Nb and O1 is also slightly larger than that to O2. The Jahn – Teller effect
Fig. 2. Powder neutron diffraction profiles for SrCu0.4Nb0.6O2.9.
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Table 1 Structure parameters of SrCu0.4Nb0.6O2.9 at room temperature from neutron diffraction data ˚ 2) Atom Site Occupancy x Y z Uiso (A Sr Cu Nb O(1) O(2)
1a 1d 1d 1c 2e
1.00573 0.39772 0.60229 0.91274 0.99645
0 0.5 0.5 0.5 0.5
0 0.5 0.5 0.5 0
0 0.02133 0.5 0.01027 0.5 0.01027 0 0.03784 0.5 0.03548 ˚ , c = 3.9757(2) Note: Space group P4/mmm (123); a = 3.9608(4) A ˚ , V = 62.37(2) A ˚ 3. Rwp = 9.67%, Rp = 7.54%, vred2 = 1.590. A
may be responsible for this distortion of the BO6 octahedra. Fig. 4. The total conductivity of SrCu0.4Nb0.6O2.9 in air.
3.2. Conductivity 3.3. Chemical stability in reducing atmosphere The conductivity of SrCu0.4Nb0.6O2.9 in air was studied by AC impedance spectroscopy. It was found that the conduction is electronically dominated. The total conductivity is about 1 10 2 S/cm at 900 jC as shown in Fig. 4. The conduction activation energy is 0.60 F 0.01 eV between 600 and 800 jC, and 0.73 F 0.02 eV between 800 and 900 jC. The slight activation energy change observed around 800 jC may be due to the change of activation energy for carrier creation, as the number of carriers may increase at high temperatures. It is expected that the conductivity may be further improved by modifying the chemical composition. However, the chemical stability is more important, which is studied by TGA.
Fig. 3. Crystal structure of SrCu0.4Nb0.6O2.9 at room temperature with space group P4/mmm.
Fig. 5 shows the TGA analyses of SrCu 0.4 Nb0.6O2.9 (a) and CuO (b) performed in dry 5% H2. For comparison, the mass loss per Cu atom in formula unit is plotted against temperature. Both correspond to reduction of CuII to Cu0 although CuII is stable to higher temperatures in the perovskite. The total weight losses for CuO and SrCu0.4Nb0.6O2.9 were 20.69% and 3.07%, respectively, close to theoretical, e.g. the mass loss is 20.4% if CuO is completely reduced to metallic copper. The slight weight gain for SrCu0.4Nb0.6O2.9 during cooling may be attributed to the reoxidisation of the sample, although it could easily be drifted due to buoyancy effects. As anticipated from TGA analysis of pure Nb2O5, no signifi-
Fig. 5. The TGA analysis of SrCu0.4Nb0.6O2.9 (a) and CuO (b) in 5% H2 in Ar.
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The expression of n-type electronic conductivity rn is, rn ¼ neln
Fig. 6. Isothermal conductivity vs. pO2 diagram for SrCu0.4Nb0.6 O2.9 in air (a), 5% H2 (b) and Ar (c) at 900 jC.
cant mass change due to the reduction of niobium was observed. Scanning electron microscopy (SEM) observation indicates that the agglomeration of yielded copper is insignificant. It is supposed that copper is homogeneously dispersed in the mixture. 3.4. Conductivity change with pO2 The conductivity change in low pO2 was also studied by slowly introducing 5% H2 into a sealed chamber at 900 jC. The pO2 was monitored by a zirconia oxygen sensor. At low pO2, SrCu0.4Nb0.6O2.9 may be reduced to metallic copper and new perovskite strontium niobates. The copper formed is not enough to create a percolation path for electron conduction. As shown in Fig. 6, the conductivity of newly formed strontium niobate is n-type at pO2 lower than 10 12 atm with a slope 1/4 which was also observed by Lecomte et al. [27] for Sr3SrNb2O9. From stoichiometry considerations, it is likely that the reduction of SrCu0.4Nb0.6O2.9 results in Cu metal and a phase similar to Sr3SrNb2O9 with a small amount of a niobium-rich impurity. XRD analysis indeed shows the presence of Cu and a Sr3SrNb2O9-like phase. At lower pO 2 , some lattice oxygen may be removed and extra oxygen vacancies formed according to Eq. (1). The corresponding mass action equation is, SS 1=2 KR ðT Þ ¼ ½VO n2 pO2
ð3Þ
ð4Þ
where n is free electron concentration, e is electronic elementary charge, ln is electron mobility. In the low pO2 region, the n-type conductivity contribution from intrinsic Schottky disorder is quite small and the majority of carriers are due to the reduction of niobium, although the concentration is insignificant in TGA analysis. Combination of Eqs. (3) and (4) gives, 1 KR ðT Þ 2 14 rn ¼ eln pO2 ð5Þ SS ½VO Considering the whole system, the concentration of SS oxygen vacancies [VO] in Eq. (5) is composed of the contribution of intrinsic Schottky disorder and the release of lattice oxygen at low pO2 according to Eq. (1). The total oxygen vacancy concentration might be constant as long as the additional oxygen vacancies in equilibrium with the ambient gas is very small compared to those of intrinsic oxygen vacancies. Therefore, the logr vs. logpO2 relation of n-type electronic conductivity at low pO2 should give a slope 1/4 which is observed in our experiment. Dry argon was introduced into the system after measurement under 5% H2 in order to determine the stability of SrCu0.4Nb0.6O2.9 in less reducing atmosphere. The results shown in Fig. 6 indicate that SrCu0.4Nb0.6O2.9 is stable until at least 10 5 atm at 900 jC. The slightly lower conductivity in argon than in air is likely due to the change of relative density during the prior redox process although a slight structure change is another possibility. The sample
Table 2 Structure parameters of Sr2Mn0.8Nb1.2O6 at room temperature from X-ray diffraction data ˚ 2) Atom Site Occupancy x y z Uiso (A Sr Mn Nb O(1) O(2)
4c 4b 4b 4c 8d
1 0.4 0.6 1 1
0.00363 0.5 0.5 0.05009 0.74837
0.00616 0 0 0.50646 0.26891
0.25 0.02174 0 0.01455 0 0.01455 0.25 0.01376 0.00435 0.01646 ˚ , b = 5.6589(2) A ˚, Note: Space group Pbnm (62); a = 5.6451(3) A ˚ , V = 254.69(7) A ˚ 3. Rwp = 3.64%, Rp = 3.26%, c = 7.9729(2) A vred2 = 2.323.
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Fig. 7. Powder XRD diffraction profiles for Sr2Mn0.8Nb1.2O6.
was then cooled down in argon to room temperature to examine the phase composition. XRD analysis indicates it contained tetragonal phases with a c ap, c c ap. 3.5. Sr2Mn0.8Nb1.2O6 system To improve the stability, Sr3Mn1.2Nb1.8O9 d was prepared using manganese instead of copper with the same BVto BW ratio as in SrCu0.4Nb0.6O2.9. It was synthesised by firing the mixture of SrCO3, MnO2 and Nb2O5 at 1400 jC for 12 h. A single pffiffiffi perovskite pffiffiffi phase was found with lattice parameter 2ap 2ap 2ap according to XRD analysis. It is a double perovskite and therefore the formula is better written as Sr2Mn0.8Nb1.2O6 d. The material exhibits an orthorhombic structure Pbnm (62) with a = 5.6451(3) ˚ , b = 5.6589(2) A ˚ , c = 7.9729(2) A ˚ , V = 254.69(7) A ˚3 A as listed in Table 2 and shown in Fig. 7. After heating the as-prepared Sr2Mn0.8Nb1.2O6 d in 5% H2 at 950 jC for half an hour and cooling down in 5% H2, the perovskite structure remained (Fig. 8(b)) and gave a ˚, orthorhombic arrangement with a = 5.6693(3) A ˚ ˚ ˚ b = 5.6828(6) A, c = 8.0379(9) A, V = 258.96(8) A3. The cell volume expands 1.7% during the reduction process, which may be attributed to the partial reduction of manganese. This is confirmed by TGA analysis
as shown in Fig. 9. The as-prepared Sr 2 Mn 0.8 Nb1.2O6 d started to lose mass at around 450 jC. The total weight loss is 0.71% between 450 and 950 jC. If it is supposed that manganese is completely reduced to MnII as suggested by Kruth et al. [20], then the manganese valence in the sample prepared in air can be deduced as + 2.47 indicating d c 0 in Sr2Mn0.8 Nb1.2O6 d (Fig. 9). It must also assumed that niobium maintains its + 5 valence. In a reducing atmosphere,
Fig. 8. XRD patterns of Sr2Mn0.8Nb1.2O6 d obtained at 1400 jC (a) and after TGA analysis till 950 jC/0.5 h and cooling down in 5% H2 (b).
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Fig. 9. The TGA analysis of Sr2Mn0.8Nb1.2O6 in 5% H2 in Ar.
the material loses lattice oxygen and provides oxygen vacancies for possible oxygen migration. The total conductivity of Sr2Mn0.8Nb1.2O6 was measured as 0.36 S/cm by a four-terminal DC method at 900 jC in air. The apparent activation energy for conduction 0.47 eV between 600 and 900 jC indicates the compound may be a semiconductor (Fig. 10). No low frequency arc was observed in the impedance spectra that might indicate diffusion or charge transfer limitation, demonstrating that electronic conduction is more important than ionic. The conductivity behaviour with pO2 at 900 jC in wet and dry 5% H2 is shown in Fig. 11. The data indicate that conductivity is pO2 independent under oxidising conditions, and that it remains stable until fairly reducing conditions. The decrease of conductivity at pO2 lower than 10 13 atm, which gives a slope close to 1/4
Fig. 11. Isothermal conductivity vs. pO2 diagram for Sr2Mn0.8 Nb1.2O6 in dry 5% H2 (5) and wet 5% H2 (o) at 900 jC.
indicating that the material is a p-type conductor at low pO2. A contributing factor for the decrease in conductivity may also be increased lattice parameters that makes the electron hopping more difficult. The comparable conductivity of this material at the same pO2 under wet and dry atmosphere within experimental derivation indicates that proton conduction, if it exists, is negligible. This is expected because it has been shown previously that the proton conduction of A3BV 1 + xBW 2 xO9 d perovskites becomes insignificant at elevated temperatures [11]. In summary, the Mn-doped strontium niobates are stable in a reducing atmosphere. Unlike copper-containing perovskites, the manganese does not reduce completely to the metal and thus the perovskite structure maintained.
4. Conclusion
Fig. 10. The DC conductivity of Sr2Mn0.8Nb1.2O6 d in air.
First row transition element-doped perovskites A3BV 1 + xBW 2 xO9 d might be potential candidates for SOFC anode materials. Single phases were obtained in Sr3M1.5Nb1.5O9 d with M = Cr, Mn, Fe, Co. SrCu0.4Nb0.6O2.9 was prepared by solid state reaction at 1000 jC. It is tetragonal with space group P4/mmm ˚ , c = 3.9757(2) A ˚ , V = 62.37(2) (123), a = 3.9608(4) A 3 ˚ A according to neutron diffraction. Rietveld refinement indicates that the oxygen vacancies tend to locate at O1 (1c) site with O2 (2e) fully occupied, i.e. with ordering along the [001] direction. The distortion of BO6 octahedra may be attributed to the
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Jahn – Teller effect because of the existence of d9 Cu2 + ions. AC impedance measurements indicate that electronic conduction is dominant in air. This material is reduced to metallic copper and new perovskite phases at elevated temperatures in 5% H2; however, it is stable in argon at 900 jC. It was observed using SEM that the copper metal formed is almost homogeneously distributed in the mixture. The DC conductivity of SrCu0.4Nb0.6O2.9 at pO2 in the range of 10 22 – 10 12 atm gives a slope 1/4. This suggests n-type electronic conduction and may be explained by a simple defect model. The stability problem of the first row transition elements-doped A3BV 1 + xBW2 x O9 d may be solved by using early elements, such as manganese instead of copper. The primary results show that Sr2Mn0.8Nb1.2O6 d is orthorhombic in air and in 5% H2 although thermal expansion was observed during the reduction. At 900 jC, the total conductivity in air is 0.36 S/cm in air; however, it exhibits p-type behaviour at low pO2.
Acknowledgements The authors would like to thank EPSRC and NEDO for the funding. Thanks also to Dr. Richard Ibberson at HRPD station, ISIS, Rutherford and Dr Angela Kruth at St. Andrews for the assistance with the collection of neutron diffraction data. One of the authors (Tao) thanks Dr. Tom McColm at St. Andrews for a critical reading of the manuscript.
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[3] N. Bonanos, B. Ellis, M.N. Mahmood, Solid State Ionics 44 (1991) 305. [4] F.L. Chen, O.T. Sørensen, G.Y. Meng, D.K. Peng, J. Mater. Chem. 7 (1997) 481. [5] H.G. Bohn, T. Schober, J. Am. Ceram. Soc. 83 (2000) 768. [6] K.D. Kreuer, Solid State Ionics 125 (1999) 285. [7] K.C. Liang, Y. Du, A.S. Nowick, Solid State Ionics 69 (1994) 117. [8] A.S. Nowick, Y. Du, Solid State Ionics 77 (1995) 137. [9] H.G. Bohn, T. Schober, T. Mono, W. Schilling, Solid State Ionics 117 (1999) 219. [10] A.S. Nowick, Y. Du, K.C. Liang, Solid State Ionics 125 (1999) 303. [11] R. Glockner, A. Neiman, Y. Larring, T. Norby, Solid State Ionics 125 (1999) 369. [12] A.S. Nowick, K.C. Liang, Solid State Ionics 129 (2000) 201. [13] F. Galasso, L. Katz, R. Ward, J. Am. Chem. Soc. 81 (1959) 820. [14] F. Galasso, J. Pyle, J. Phys. Chem. 67 (1963) 1561. [15] U. Treiber, S. Kemmler-Sack, J. Solid State Chem. 43 (1982) 51. [16] J.A. Chavez, T.C. Gibb, A.R. West, J. Mater. Chem. 6 (1996) 1957. [17] J.H. Choy, S.T. Hong, K.S. Choi, J. Chem. Soc., Faraday Trans. 92 (1996) 1051. [18] H. Langbein, M. Bremer, I. Krabbes, Solid State Ionics 101 – 103 (1997) 579. [19] A.R. Chakhmouradian, R.H. Mitchell, J. Solid State Chem. 138 (1998) 272. [20] A. Kruth, U. Guth, A.R. West, J. Mater. Chem. 9 (1999) 1579. [21] J.H. Choy, J.H. Park, S.T. Hong, D.K. Kim, J. Solid State Chem. 111 (1994) 370. [22] I. Levin, L.A. Bendersky, J.P. Cline, R.S. Roth, T.A. Vanderah, J. Solid State Chem. 150 (2000) 43. [23] I. Levin, J.Y. Chan, R.G. Geyer, J.E. Maslar, T.A. Vanderah, J. Solid State Chem. 156 (2001) 122. [24] S. Park, J.M. Vohs, R.J. Gorte, Nature 404 (2000) 265. [25] A.C. Larson, R.B. Von Dreele, GSAS—Generalised Crystal Structure Analysis System, Los Alamos National Laboratory Report No. LA-UR-86-748 (1994). [26] E.G. Fesenko, V.G. Smotrakov, V.V. Eremkin, L.A. Shilkina, Inorg. Mater. 28 (1992) 1751. [27] J. Lecomte, J.P. Loup, M. Hervieu, B. Raveau, Phys. Status Solidi, A 65 (1981) 743.
Structure and properties of nonstoichiometric mixed perovskites A3BV 1+xBW 2xO9d Shanwen Tao, John T.S. Irvine * School of Chemistry, University of St. Andrews, St. Andrews, Fife KY16 9ST, Scotland, UK Accepted 14 March 2002
Abstract Nonstoichiometric mixed perovskites A3BV 1 + xBW 2 xO9 d, e.g. Ba3Ca1.18Nb1.82O9 d, exhibit high proton and oxygen-ion conductivity. It is expected that mixed ionic and electronic conductors may be found in these compounds if the B-sites are partially substituted by a first row transition element. These mixed conductors may be potential anode materials for fuel cell applications. The structure of single phase SrCu0.4Nb0.6O2.9 was studied by both X-ray and neutron diffraction. It is tetragonal ˚ , c = 3.9757(2) A ˚ , V = 62.37(2) A ˚ 3 according to neutron diffraction. Rietveld with space group P4/mmm (123), a = 3.9608(4) A refinement indicates that the oxygen vacancy tends to stay at O1 (1c) site with O2 (2e) fully occupied. AC impedance measurements indicate that electronic conduction is probably dominant in air. The DC conductivity of SrCu0.4Nb0.6O2.9 at pO2 in the range of 10 22 – 10 12 atm exhibits a p(O2) 1/4 dependence consistent with n-type electronic conduction. The material was unstable in 5% H2 at elevated temperatures but stable in argon at 900 jC. Using manganese instead of copper, a phase that is redox stable was prepared. SrMn0.4Nb0.6O3 d exhibits an orthorhombic structure with space group Pbnm (62), ˚ , b = 5.6589(2) A ˚ , c = 7.9729(2) A ˚ , V = 254.69(7) A ˚ 3 according to X-ray diffraction. Such a unit cell indicates a = 5.6451(3) A that it is a double perovskite and therefore the formula is better written as Sr2Mn0.8Nb1.2O6 d. The material maintains perovskite structure in 5% H2 although thermal expansion was observed on reduction. The conductivity of Sr2Mn0.8Nb1.2O6 is 0.36 S/cm in air at 900 jC. Conductivity decreases in 5% H2 indicates p-type conduction at low pO2. D 2002 Elsevier Science B.V. All rights reserved. Keywords: Electrical conductivity; Perovskite structure; Neutron diffraction; Stability; Fuel cell; Defect
1. Introduction High temperature proton conductors have been of great interest for their potential applications in solid oxide fuel cells, sensors and steam electrolyzers since their discovery, primarily by Iwahara et al. [1] and Uchida et al. [2]. These rare earth elements-doped ABO3 perovskite oxides with A = Ca, Sr, Ba, B = Ce, *
Corresponding author. Tel.: +44-1334-463817; fax: +441334-463808. E-mail address: [email protected] (J.T.S. Irvine).
Zr and Ti, etc. have been studied in detail [3 –6]. An interesting discovery was that some nonstoichiometric mixed perovskites with formula A3BV 1 + xBW 2 xO9 d where A is an alkali earth element Ca, Sr, Ba, etc., BVis an element with valence + 2 or + 3 and BW is an element with valence + 5, such as Nb and Ta exhibit quite high proton and oxygen-ion conductivity [7 –12]. It is expected that mixed ionic and electronic conductors may be found in these compounds if the B-site is partially substituted by appropriate first row transition elements. Such mixed conductors may provide potential anode materials for fuel cell applications.
0167-2738/02/$ - see front matter D 2002 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 7 - 2 7 3 8 ( 0 2 ) 0 0 5 1 6 - 7
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In a reducing atmosphere, a transition elementcontaining perovskite oxide loses some lattice oxygen due to the reduction of transition elements coupled with the constraint of electroneutrality, SS OxO ¼ 1=2O2 þ VO þ 2eV ð1Þ The released electrons may result in the increase of ntype conductivity or decrease of p-type conductivity. The oxygen vacancies created may be used for oxygen transport leading to enhanced oxygen-ion conductivity. They may also combine with water vapour resulting in the formation of proton defects. SS S VO þ H2 OðgÞ þ OxO ¼ 2OHO
ð2Þ
Therefore, the first row transition element substituted A3BV 1 + xBW 2 xO9 d perovskites may be potential oxygen-ion (proton) and electronic mixed conductors. Perovskite-type oxides with general formula A3 BV 1 + xBW 2 xO9 d with A = Ca, Sr, Ba, etc., BV= Ca, Sr, Cr, Mn, Cu, Fe, Co, Ni, etc., BW = Nb, Ta have been synthesised and widely reported [13 – 22]. A wide range of structures from high cubic symmetry to the lowest triclinic symmetry has been observed. For example, Sr2CrNbO6 is a double perovskite with alternate B-site ordering. Ca3CaNb2O9 prepared under different conditions gives 1:1, 1:2 and 1:3 orderings, respectively [23]. The ordering of B-site cations and tilting of BO6 octahedral make the real structure of these perovskite oxides very flexible. Splitting of the B-site is commonly observed because two (or more) different elements BVand BW are involved. For example, in A3BVBBW 2O9, 1:1 ordering often occur if the Bsite splits with one position fully occupied by BW and the other one shared by 2/3BV and 1/3BW. However, sometimes this ordering is not complete. Partial ordering can occur if the two B-sites are still occupied by both BVand BW but with different cation ratios. Similarly, the B-site ordering may be 1:1, 1:2 or 1:3 depending on the elements involved and the synthesis conditions. The formulae of complex perovskite oxides with 1:1 and 1:2 B-site ordering are normally written as A2BVBWO6 for double perovskites and A3BVBW 2O9 for triple perovskites. Most compositions of the studied oxides have the formula of A3BVBW 2O 9 where theoretically, no oxygen deficiency exists according to electroneutrality. Defect concentrations and thus defect chemistry in perovskites can be
controlled and tailored significantly by doping. Departure from the composition of a crystallographically perfect solid by the substitution of small concentrations of similar sized hetero-valent ions has an immense impact on properties. By the introduction of di- or trivalent transition elements onto the BV-sites, materials possessing mixed ionic and electronic conductivity may be obtained. It was recently reported that the cermet Cu-ceria may be used as a SOFC anode material for direct oxidation of hydrocarbons without carbon deposition, indicating that copper-containing materials might be suitable catalysts for the complete oxidation of hydrocarbon in fuel cell operation [24]. It would be interesting if copper can be introduced into a complex perovskite oxide retaining both high ionic and electronic conductivity. The stability of a Cu-containing oxide in a reducing atmosphere with pO2 as low as 10 15 – 10 20 atm for SOFCs is a potential problem. However, copper metal may be homogeneously dispersed in the mixture if obtained directly from the reduction of Cu-containing complex perovskite oxides provided that sintering of the copper can be avoided. From this point of view, our first target was to synthesise the nonstoichiometric oxide Sr3Cu1 + x Nb2 xO9 d. The structure, conductivity and stability of SrCu0.4Nb0.6O2.9 were studied by X-ray diffraction, neutron diffraction, AC impedance, TGA, etc. It was found that the chemical stability of SrCu0.4Nb0.6 O2.9 at low pO2 is not good. The stability is improved when copper is replaced by the early first row transition element manganese. The structure and electrical properties of Sr2Mn0.8Nb1.2O6 are also reported.
2. Experimental 2.1. Powder synthesis The nonstoichiometric mixed perovskites were synthesised by solid state reaction using SrCO3, Cr 2 O 3 , MnO 2 , Fe 2 O 3 , CoCO 3 , NiO, CuO and Nb2O5 as starting materials. The carbonates and MnO2 were dried at 100 – 300 jC, oxides except MnO2 at 300 –800 jC overnight to remove absorbed H2O and CO2. Stoichiometric amounts were mixed and ball milled in zirconia containers with zirconia balls for two 15-min periods. The final firing temper-
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atures are between 1000 and 1500 jC according to the chemical composition. The mixture for Sr3Cu1 + x Nb2 xO9 d was calcined at 1000 – 1300 jC. 2.2. Sample characterisation XRD analysis of powders reacted at different temperatures was carried out on a Stoe Stadi-P diffractometer to determine phase purity and measure crystal parameters. Electron diffraction was carried out on a JEOL 2011 HRTEM operating at 200 keV. Powder neutron diffraction data were collected on the diffractometer HRPD, ISIS, Rutherford Appleton Laboratory. Structure refinement was performed by the Rietveld method using the program General Structure Analysis System (GSAS) [25]. Thermal analysis of SrCu0.4Nb0.6O2.9 was carried out on a Rheometric Scientific TG 1000M+ TA Instruments SDT2960 from room temperature to 950 jC (5 jC/ min), holding at 950 jC for 30 min then cooling down to 30 jC (10 jC/min) under flowing 5% H2 in argon at a rate of 35 ml/min. 2.3. Conductivity measurements For AC impedance measurements in air, a Schlumberger Solartron 1260 Frequency Response Analyser coupled with a 1287 Electrochemical Interface controlled by Zplot electrochemical impedance software was used over the frequency range 1– 100 mHz. AC impedance measurements were made in 50 jC steps in air between 600 and 900 jC. The DC conductivity was measured by a conventional four-terminal method using a Keithley 220 Programmable Current Source to control current and a Schlumberger Solartron 7150 Digital Multimeter to measure the voltage. The SrCu0.4Nb0.6O2.9 samples were mounted with four Pt wire electrodes to measure the DC conductivity dependence on pO2, which was monitored by a zirconia oxygen sensor.
3. Results and discussion 3.1. Structure of SrCu0.4Nb0.6O2.9 For the SrM0.5Nb0.5O3 y system with M = Cr, Mn, Fe, Co. Ni, Cu, a single phase was achieved for all but
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M = Ni, Cu. The second phase for nominal formula Sr3Cu1.5Nb1.5O9 d is SrCuO2 (JCPDS No. 39-1492). In order to get a single perovskite, a series of Sr3 Cu1 + xNb2 xO9 d (x V 0.5) was prepared by solid state reaction as described above. After calcining the ball-milled mixture of SrCO3, CuO and Nb2O5 at 1000 jC, the main phase for Sr3CuNb2O 9 is a perovskite, but not well crystallised. It also contains some un-indexed peaks, which may belong to second phases. Further heating the oxide at 1300 jC, the main phase is well crystallised with a weak impurity peak at 2h c 30j, which may be due to traces of the phase Sr5Nb4O15 (JCPDS Card No. 28-1248). The main phase is a copper-rich strontium niobate with 1:1 Bsite ordering along the [111]c* zone according to electron diffraction. The main phase a ffiffitetragonal pffiffiffi is p ffi perovskite with lattice parameter 2ap 2ap 2ap. Fesenko et al [26] reported that a second tetragonal phase was detected when calcining the starting materials Sr2CuO3 and SrNb2O6 at high temperatures (1270 K, 12 h) to form the perovskite phase Sr3CuNb2O9. Ba3CuNb2O9 was not obtained either when prepared by solid state reaction using BaCO3, CuO and Nb2O5 as precursors [18]. Phase purity was only found with x = 0.2 in the Sr 3 Cu 1 + x Nb 2 x O 9 d (0 V x V 1) system. With x < 0.2, second phases were evident, e.g. Sr5Nb4O15 and second perovskite phases. When x>0.2, the second phase SrCuO2 was identified. No super structure was observed by electron diffraction indicating a simple perovskite. The accurate formula for Sr3Cu1.2Nb1.8O8.7 is SrCu0.4Nb0.6O2.9. Fig. 1 shows the electron diffraction of a cubic primitive cell indicating no super structure along [100] c* and [010]c* zones. The lattice parameters for a and b appear to be equal. No symmetrical absences were observed in the XRD pattern indicating a primitive cubic or tetragonal structure with slight distortion. Space group Pm3¯m and P4/mmm were chosen for Rietveld refinement of X-ray and neutron diffraction data. From the refinement of XRD data, the cubic Pm3¯m model gave smaller v2 value and R-factors indicating it might be a cubic perovskite; however, there was visible extra broadening on peaks that would be expected to split under tetragonal distortion. Clearly, this preference for the cubic space group was an artefact due to the limited resolution of the Xdiffractometer, even though a high resolution labora-
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Fig. 1. Electron diffraction pattern of SrCu0.4Nb0.6O2.9. The electron beam is parallel to [001].
tory system was used. Therefore, the real symmetry for SrCu0.4Nb0.6O2.9 should be lower than cubic. In addition, X-ray diffraction is not so sensitive to the light oxygen atoms making it difficult to determine the exact oxygen positions. The problem may be solved by neutron diffraction because the neutron
scattering length of oxygen is slightly lower than strontium, copper and niobium. It is believed that neutron diffraction may precisely determine the oxygen atom positions; however, it should be noted that neutron diffraction does not discriminate so well between the cations. The final refined occupancies of Sr, Cu and Nb from XRD are very close to the ratio in the starting formula although the oxygen occupancy cannot be refined from the XRD data. The occupancies for metal elements from XRD data were introduced and held constant in the refinement of the neutron data. The tetragonal model P4/mmm gives a smaller v2 value (1.590) and R-factor than those for the cubic model Pm3¯m (v2 = 1.953). The refined oxygen stoichiometry is consistent with the formula SrCu0.4Nb0.6O2.9. Fig. 2 shows the neutron patterns of SrCu0.4Nb0.6O2.9. The final refined structure data are given in Table 1. During the refinement, it was found that O1 tends to be partially occupied indicating vacancies may preferentially locate on this site as shown in Fig. 3. There is a degree of ordering of the oxygen vacancies at room temperature. Such ordering may limit the diffusion of charge carriers SS VO resulting in lower oxygen-ion conductivity. However, at high temperatures, this situation may change. The bond distance between Cu/Nb and O1 is also slightly larger than that to O2. The Jahn – Teller effect
Fig. 2. Powder neutron diffraction profiles for SrCu0.4Nb0.6O2.9.
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Table 1 Structure parameters of SrCu0.4Nb0.6O2.9 at room temperature from neutron diffraction data ˚ 2) Atom Site Occupancy x Y z Uiso (A Sr Cu Nb O(1) O(2)
1a 1d 1d 1c 2e
1.00573 0.39772 0.60229 0.91274 0.99645
0 0.5 0.5 0.5 0.5
0 0.5 0.5 0.5 0
0 0.02133 0.5 0.01027 0.5 0.01027 0 0.03784 0.5 0.03548 ˚ , c = 3.9757(2) Note: Space group P4/mmm (123); a = 3.9608(4) A ˚ , V = 62.37(2) A ˚ 3. Rwp = 9.67%, Rp = 7.54%, vred2 = 1.590. A
may be responsible for this distortion of the BO6 octahedra. Fig. 4. The total conductivity of SrCu0.4Nb0.6O2.9 in air.
3.2. Conductivity 3.3. Chemical stability in reducing atmosphere The conductivity of SrCu0.4Nb0.6O2.9 in air was studied by AC impedance spectroscopy. It was found that the conduction is electronically dominated. The total conductivity is about 1 10 2 S/cm at 900 jC as shown in Fig. 4. The conduction activation energy is 0.60 F 0.01 eV between 600 and 800 jC, and 0.73 F 0.02 eV between 800 and 900 jC. The slight activation energy change observed around 800 jC may be due to the change of activation energy for carrier creation, as the number of carriers may increase at high temperatures. It is expected that the conductivity may be further improved by modifying the chemical composition. However, the chemical stability is more important, which is studied by TGA.
Fig. 3. Crystal structure of SrCu0.4Nb0.6O2.9 at room temperature with space group P4/mmm.
Fig. 5 shows the TGA analyses of SrCu 0.4 Nb0.6O2.9 (a) and CuO (b) performed in dry 5% H2. For comparison, the mass loss per Cu atom in formula unit is plotted against temperature. Both correspond to reduction of CuII to Cu0 although CuII is stable to higher temperatures in the perovskite. The total weight losses for CuO and SrCu0.4Nb0.6O2.9 were 20.69% and 3.07%, respectively, close to theoretical, e.g. the mass loss is 20.4% if CuO is completely reduced to metallic copper. The slight weight gain for SrCu0.4Nb0.6O2.9 during cooling may be attributed to the reoxidisation of the sample, although it could easily be drifted due to buoyancy effects. As anticipated from TGA analysis of pure Nb2O5, no signifi-
Fig. 5. The TGA analysis of SrCu0.4Nb0.6O2.9 (a) and CuO (b) in 5% H2 in Ar.
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The expression of n-type electronic conductivity rn is, rn ¼ neln
Fig. 6. Isothermal conductivity vs. pO2 diagram for SrCu0.4Nb0.6 O2.9 in air (a), 5% H2 (b) and Ar (c) at 900 jC.
cant mass change due to the reduction of niobium was observed. Scanning electron microscopy (SEM) observation indicates that the agglomeration of yielded copper is insignificant. It is supposed that copper is homogeneously dispersed in the mixture. 3.4. Conductivity change with pO2 The conductivity change in low pO2 was also studied by slowly introducing 5% H2 into a sealed chamber at 900 jC. The pO2 was monitored by a zirconia oxygen sensor. At low pO2, SrCu0.4Nb0.6O2.9 may be reduced to metallic copper and new perovskite strontium niobates. The copper formed is not enough to create a percolation path for electron conduction. As shown in Fig. 6, the conductivity of newly formed strontium niobate is n-type at pO2 lower than 10 12 atm with a slope 1/4 which was also observed by Lecomte et al. [27] for Sr3SrNb2O9. From stoichiometry considerations, it is likely that the reduction of SrCu0.4Nb0.6O2.9 results in Cu metal and a phase similar to Sr3SrNb2O9 with a small amount of a niobium-rich impurity. XRD analysis indeed shows the presence of Cu and a Sr3SrNb2O9-like phase. At lower pO 2 , some lattice oxygen may be removed and extra oxygen vacancies formed according to Eq. (1). The corresponding mass action equation is, SS 1=2 KR ðT Þ ¼ ½VO n2 pO2
ð3Þ
ð4Þ
where n is free electron concentration, e is electronic elementary charge, ln is electron mobility. In the low pO2 region, the n-type conductivity contribution from intrinsic Schottky disorder is quite small and the majority of carriers are due to the reduction of niobium, although the concentration is insignificant in TGA analysis. Combination of Eqs. (3) and (4) gives, 1 KR ðT Þ 2 14 rn ¼ eln pO2 ð5Þ SS ½VO Considering the whole system, the concentration of SS oxygen vacancies [VO] in Eq. (5) is composed of the contribution of intrinsic Schottky disorder and the release of lattice oxygen at low pO2 according to Eq. (1). The total oxygen vacancy concentration might be constant as long as the additional oxygen vacancies in equilibrium with the ambient gas is very small compared to those of intrinsic oxygen vacancies. Therefore, the logr vs. logpO2 relation of n-type electronic conductivity at low pO2 should give a slope 1/4 which is observed in our experiment. Dry argon was introduced into the system after measurement under 5% H2 in order to determine the stability of SrCu0.4Nb0.6O2.9 in less reducing atmosphere. The results shown in Fig. 6 indicate that SrCu0.4Nb0.6O2.9 is stable until at least 10 5 atm at 900 jC. The slightly lower conductivity in argon than in air is likely due to the change of relative density during the prior redox process although a slight structure change is another possibility. The sample
Table 2 Structure parameters of Sr2Mn0.8Nb1.2O6 at room temperature from X-ray diffraction data ˚ 2) Atom Site Occupancy x y z Uiso (A Sr Mn Nb O(1) O(2)
4c 4b 4b 4c 8d
1 0.4 0.6 1 1
0.00363 0.5 0.5 0.05009 0.74837
0.00616 0 0 0.50646 0.26891
0.25 0.02174 0 0.01455 0 0.01455 0.25 0.01376 0.00435 0.01646 ˚ , b = 5.6589(2) A ˚, Note: Space group Pbnm (62); a = 5.6451(3) A ˚ , V = 254.69(7) A ˚ 3. Rwp = 3.64%, Rp = 3.26%, c = 7.9729(2) A vred2 = 2.323.
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Fig. 7. Powder XRD diffraction profiles for Sr2Mn0.8Nb1.2O6.
was then cooled down in argon to room temperature to examine the phase composition. XRD analysis indicates it contained tetragonal phases with a c ap, c c ap. 3.5. Sr2Mn0.8Nb1.2O6 system To improve the stability, Sr3Mn1.2Nb1.8O9 d was prepared using manganese instead of copper with the same BVto BW ratio as in SrCu0.4Nb0.6O2.9. It was synthesised by firing the mixture of SrCO3, MnO2 and Nb2O5 at 1400 jC for 12 h. A single pffiffiffi perovskite pffiffiffi phase was found with lattice parameter 2ap 2ap 2ap according to XRD analysis. It is a double perovskite and therefore the formula is better written as Sr2Mn0.8Nb1.2O6 d. The material exhibits an orthorhombic structure Pbnm (62) with a = 5.6451(3) ˚ , b = 5.6589(2) A ˚ , c = 7.9729(2) A ˚ , V = 254.69(7) A ˚3 A as listed in Table 2 and shown in Fig. 7. After heating the as-prepared Sr2Mn0.8Nb1.2O6 d in 5% H2 at 950 jC for half an hour and cooling down in 5% H2, the perovskite structure remained (Fig. 8(b)) and gave a ˚, orthorhombic arrangement with a = 5.6693(3) A ˚ ˚ ˚ b = 5.6828(6) A, c = 8.0379(9) A, V = 258.96(8) A3. The cell volume expands 1.7% during the reduction process, which may be attributed to the partial reduction of manganese. This is confirmed by TGA analysis
as shown in Fig. 9. The as-prepared Sr 2 Mn 0.8 Nb1.2O6 d started to lose mass at around 450 jC. The total weight loss is 0.71% between 450 and 950 jC. If it is supposed that manganese is completely reduced to MnII as suggested by Kruth et al. [20], then the manganese valence in the sample prepared in air can be deduced as + 2.47 indicating d c 0 in Sr2Mn0.8 Nb1.2O6 d (Fig. 9). It must also assumed that niobium maintains its + 5 valence. In a reducing atmosphere,
Fig. 8. XRD patterns of Sr2Mn0.8Nb1.2O6 d obtained at 1400 jC (a) and after TGA analysis till 950 jC/0.5 h and cooling down in 5% H2 (b).
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Fig. 9. The TGA analysis of Sr2Mn0.8Nb1.2O6 in 5% H2 in Ar.
the material loses lattice oxygen and provides oxygen vacancies for possible oxygen migration. The total conductivity of Sr2Mn0.8Nb1.2O6 was measured as 0.36 S/cm by a four-terminal DC method at 900 jC in air. The apparent activation energy for conduction 0.47 eV between 600 and 900 jC indicates the compound may be a semiconductor (Fig. 10). No low frequency arc was observed in the impedance spectra that might indicate diffusion or charge transfer limitation, demonstrating that electronic conduction is more important than ionic. The conductivity behaviour with pO2 at 900 jC in wet and dry 5% H2 is shown in Fig. 11. The data indicate that conductivity is pO2 independent under oxidising conditions, and that it remains stable until fairly reducing conditions. The decrease of conductivity at pO2 lower than 10 13 atm, which gives a slope close to 1/4
Fig. 11. Isothermal conductivity vs. pO2 diagram for Sr2Mn0.8 Nb1.2O6 in dry 5% H2 (5) and wet 5% H2 (o) at 900 jC.
indicating that the material is a p-type conductor at low pO2. A contributing factor for the decrease in conductivity may also be increased lattice parameters that makes the electron hopping more difficult. The comparable conductivity of this material at the same pO2 under wet and dry atmosphere within experimental derivation indicates that proton conduction, if it exists, is negligible. This is expected because it has been shown previously that the proton conduction of A3BV 1 + xBW 2 xO9 d perovskites becomes insignificant at elevated temperatures [11]. In summary, the Mn-doped strontium niobates are stable in a reducing atmosphere. Unlike copper-containing perovskites, the manganese does not reduce completely to the metal and thus the perovskite structure maintained.
4. Conclusion
Fig. 10. The DC conductivity of Sr2Mn0.8Nb1.2O6 d in air.
First row transition element-doped perovskites A3BV 1 + xBW 2 xO9 d might be potential candidates for SOFC anode materials. Single phases were obtained in Sr3M1.5Nb1.5O9 d with M = Cr, Mn, Fe, Co. SrCu0.4Nb0.6O2.9 was prepared by solid state reaction at 1000 jC. It is tetragonal with space group P4/mmm ˚ , c = 3.9757(2) A ˚ , V = 62.37(2) (123), a = 3.9608(4) A 3 ˚ A according to neutron diffraction. Rietveld refinement indicates that the oxygen vacancies tend to locate at O1 (1c) site with O2 (2e) fully occupied, i.e. with ordering along the [001] direction. The distortion of BO6 octahedra may be attributed to the
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Jahn – Teller effect because of the existence of d9 Cu2 + ions. AC impedance measurements indicate that electronic conduction is dominant in air. This material is reduced to metallic copper and new perovskite phases at elevated temperatures in 5% H2; however, it is stable in argon at 900 jC. It was observed using SEM that the copper metal formed is almost homogeneously distributed in the mixture. The DC conductivity of SrCu0.4Nb0.6O2.9 at pO2 in the range of 10 22 – 10 12 atm gives a slope 1/4. This suggests n-type electronic conduction and may be explained by a simple defect model. The stability problem of the first row transition elements-doped A3BV 1 + xBW2 x O9 d may be solved by using early elements, such as manganese instead of copper. The primary results show that Sr2Mn0.8Nb1.2O6 d is orthorhombic in air and in 5% H2 although thermal expansion was observed during the reduction. At 900 jC, the total conductivity in air is 0.36 S/cm in air; however, it exhibits p-type behaviour at low pO2.
Acknowledgements The authors would like to thank EPSRC and NEDO for the funding. Thanks also to Dr. Richard Ibberson at HRPD station, ISIS, Rutherford and Dr Angela Kruth at St. Andrews for the assistance with the collection of neutron diffraction data. One of the authors (Tao) thanks Dr. Tom McColm at St. Andrews for a critical reading of the manuscript.
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