Materials Chemistry and Physics 111 (2008) 335–340

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Materials Chemistry and Physics

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Structural characterization, stability and electrical properties ofstrontium niobate ceramic

efu Roplied

de Sr2

6 stillloseseramfromce inn kin

Tian Xiaa,∗, Qin Lib, Jian Mengb, Xueqiang Caob

a School of Chemistry, Chemical Engineering and Materials, Heilongjiang University, Xub State Key Laboratory of Application of Rare Earth Resources, Changchun Institute of ApChangchun 130022, Jilin, People’s Republic of China

a r t i c l e i n f o

Article history:Received 26 October 2007Received in revised form 13 March 2008Accepted 11 April 2008

Keywords:Double perovskiteSr2CrNbO6

ConductivityOxygen partial pressure

a b s t r a c t

The double perovskite oxireducing in CO, Sr2CrNbOindicated that Sr2CrNbO6

and compositions of this catmospheres were variedpartial pressure dependento reflect slow equilibratio

1. Introduction

Mixed ionic–electronic conductors (MIECs) which exhibit bothionic and electronic conductivity have numerous technologi-

cal applications, such as water electrolysis, electrocatalytic, gassemipermeation membranes, electrolytes and electrodes for solidoxide fuel cell (SOFC), etc. [1–6]. For practical applications, thesematerials were required to have high ionic–electronic conductivityand sustainable structure stability under the harsh condition.

Perovskite-like oxides (ABO3) are among the most promisinggroups of mixed conductors due to their high oxygen ionic andelectronic conductivities at elevated temperatures in various atmo-spheres. One of the typical examples is La1−xSrxFeO3−ı, whose totalconductivity was 200–500 S cm−1 at intermediate temperaturesand ionic conductivity is one order of magnitude higher than thatof stabilized zirconia [7]. The simple perovskite structure may beappropriately modified by incorporating two types of B-site ionswith suitable different size and charge.

In recent years, some perovskites with nominal formulaA2B′B′′O6, AB′

0.5B′′0.5O3 and A3B′B′′

2O9 have attracted a great inter-ests due to the discover of Sr2FeMoO6 [8], where A is alkaline earthelement Sr, Ca and Ba, etc., B′ is a element with valence +2 or+3, and B′′ is a metal ion with valence +5 or +6, such as Nb, Ta,Mo and W (the mean charge on the B-sites is +4 when the com-

∗ Corresponding author. Tel.: +86 451 86608426; fax: +86 451 86608040.E-mail address: xiatian0621@yahoo.com.cn (T. Xia).

0254-0584/$ – see front matter © 2008 Elsevier B.V. All rights reserved.doi:10.1016/j.matchemphys.2008.04.021

ad, Heilongjiang Harbin 150080, People’s Republic of ChinaChemistry, Chinese Academy of Sciences,

CrNbO6 has a cubic structure according to powder X-ray diffraction. Afterexhibited a cubic structure refined by Rietveld technique. The TG analysis0.127 oxygen per formula unit from 400 to 700 ◦C in H2. The morphologyic did not significantly change on reduction. The conductivities in various0.015 to 0.039 S cm−1 at 800 ◦C. The conductivity exhibits a small oxygenoxygen partial pressure range of 1–10−16 atm, and this may be interpretedetics with decreasing oxygen activity.

© 2008 Elsevier B.V. All rights reserved.

pound is stoichiometric). The archetypal crystal structure of thisfamily is primitive cubic cell with space group Fm3m. The dou-ble perovskite structure is well known, and the framework can bedescribed as a build of chains of BO6 octahedrons running paral-lel to the c-axis, and A-site cations locate in the interstitial sitesbetween the octahedrons. The resulting unit cell may be viewedas doubled along the three axes, regarding the primitive cell of

ABO3. If the charge of B′ and B′′ is different, the oxygens are slightlyshifted toward the more charged cations in the ordered struc-ture although the octahedral symmetry of B cations is preserved[9]. Some mixed perovskites have quite high proton and oxygenion conductivity [10–13]. It is noted that mixed electronic andionic conductivity may be found in these materials if the A- orB-sites are partially substituted by other appropriate elements.Tao and co-workers reported some double perovskite materials,Sr2FeNbO6 [14], Sr2GaNbO6 and Sr2Mn0.5Nb0.5O3−ı [15,16], whichexhibits the electronic–ionic conductivity, strong chemical stabilityand the electronic conduction is predominant in these materi-als. It is noted that oxygen vacancies can be produced simply bychanging the stoichiometry of B-sites metal ions. An example isSr2(Sc1+xNb1−x)O6−ı (x = 0.05 and 0.1) [17]. The proton conductiv-ity of the simple perovskite Sr2(Sc1+xNb1−x)O6−ı is in the M3+ dopedSrCeO3 and BaCeO3 region. The ionic conducting behavior is domi-nant even in a severe reducing atmosphere because the compounddoes not contain Ce. Various changes such as aliovalent doping,temperature and the oxygen partial pressure in the ambient atmo-sphere may result in the formation of oxygen vacancies in oxideperovskites [18].

try and Physics 111 (2008) 335–340

due to the insensitivity of powder X-ray diffraction to small varia-tions in oxygen content in the presence of higher atomic numberelements although the refinement results were consistent with fulloccupation of these sites. The oxygen occupancy at the 24e sites wasfixed at 1 during the refinement as the initial stoichiometry seemsmost like to be Sr2CrNbO6. The Rietveld refinement converged to anRwp factor of 9.16%. The final atomic coordinates, occupancies andthermal parameters are listed in Table 1. The observed, calculatedand difference profiles are shown in Fig. 1.

Fig. 2 shows the thermogravimetric analyses (TGA) of Sr2CrNbO6performed in H2 (99.9%). The sample weight remains constantbelow 400 ◦C. There is a rapid weight loss occurring between 400and 700 ◦C in two steps, which may be attributed to the loss oflattice oxygen with the reduction of Nb5+ to Nb4+ or Cr3+ to Cr2+.The total weight loss in the range of 20–700 ◦C is only 0.49 wt%,suggesting that Sr2CrNbO6 is stable under the reducing condition.The formula of the reduced sample may therefore be written as

336 T. Xia et al. / Materials Chemis

The structure of Sr2CrNbO6 prepared by a solid-state reactionwas first reported as a primitive cubic cell with a = 7.8732(1) A [19].Therefore, in this paper, the electrical properties of Sr2CrNbO6 inoxidizing and reducing atmospheres were investigated. The elec-trochemical characterization shows that both ionic and electronicconductivities took place. On the other word, the chemical stabilityand crystal structure of Sr2CrNbO6 after the reduction were evalu-ated according to Rietveld refinement of XRD data, TG analysis andXPS measurements.

2. Experimental

Sr2CrNbO6 was synthesized by conventional solid-state reaction using SrCO3

(Beijing Chemical Company, China, >99%), CrO2 (Aldrich Chemical Company, US,>99%) and Nb2O5 (Shanghai Chemical Company, China, >99.5%) as starting materials.SrCO3 was dried at 150 ◦C, Nb2O5 at 500 ◦C overnight to remove absorbed H2O andCO2. Stoichiometric amounts were mixed in an agate mortar for 30 min. The finalfiring temperature was 1200 ◦C with intermediate grindings until single phase wasobtained. The pellets for conductivity were prepared by hand grinding the powderformed at 1200 ◦C, pressing at 150 MPa into pellets, and then sintering at 1200 ◦Cfor 8 h. The relative density of the as-prepared pellets is about 86% of theoreticaldensity. To study the phase stability of the material in reducing atmospheres, thesample obtained at 1200 ◦C was further heated in CO (99.5%) at 800 ◦C for 12 h andcooled slowly down to room temperature in CO.

XRD analyses (DMax 2500 diffractometer, Rigaku, Japan) of powder were car-ried out with Cu K� radiation (� = 1.5416 A) to determine phase purity and measurecell parameters. The 2� ranged from 10◦ to 120◦ with increments of 0.02◦ and acounting time of 1 s. Structure refinement was performed by the Rietveld methodusing the program General Structure Analysis System (GSAS) [19]. Thermal analy-sis of Sr2CrNbO6 was carried out in H2 (99.9%) on a instrument (SDT 2960, TA, US)from room temperature to 700 ◦C (10 ◦C min−1), holding at 700 ◦C for 30 min thencooling slowly down to room temperature under flowing H2 at a rate of 50 ml min−1.Field-emission scanning electron microscopy (FE-SEM) observations (XL-30, Philips,Holland) were conducted with energy-dispersive spectroscopy analysis (EDS). Thevalence states of the chromium and niobium ions for the samples were determinedby X-ray photoelectron spectroscopy (XPS). The XPS (ESCALAB 250 X-ray photoelec-tron spectrometer, Thermo Co., US) for powder samples was measured with Al K�radiation (hv = 1486.6 eV); the base pressure was 10−8 Pa. The binding energy (BE)for the samples was calibrated by setting the measured BE of C1s to 284.6 eV.

For a.c. impedance measurements in air, a Solartron 1255 Frequency ResponseAnalyser (Solartron Co., UK) coupled with a 1287 Electrochemical Interface con-trolled by Zplot electrochemical impedance software was used over the frequencyrange from 1 MHz to 0.1 Hz. Two electrodes of platinum paste were painted oneither side of the sample, and fired at 800 ◦C during 60 min to ensure adhesion.a.c. impedance measurements were made in 50 ◦C steps in air between 300 and800 ◦C. The d.c. conductivity was measured by a conventional four-probe methodusing a Keithley 2400 Programmable Current Source (Keithley Co., US) to controlcurrent and a Keithley 2000 Digital Multimeter to measure voltage. The Sr2CrNbO6

samples were mounted with four Pt wire electrodes to measure the d.c. conductivitydependence upon oxygen partial pressure in a slowly varying atmosphere, whichmonitored by a zirconia oxygen sensor. The oxygen partial pressure (PO2, with unit

“atm” in this work) in the ranges of log(PO2) = 1 to −4 and log(PO2) = −9 to −16 wascontrolled by the mixture of O2 + N2 and CO + CO2, respectively. The process con-sisted of flushing the system with O2, and then oxygen partial pressure decreasedslowly to low PO2.

3. Results and discussion

As stated above, the double perovskite compound, Sr2CrNbO6was obtained on reaction of the carbonate and oxide precursors.Fig. 1 shows the X-ray diffraction patterns of Sr2CrNbO6 obtained at1200 ◦C. A perovskite phase was obtained, and its pattern indexed inthe X-ray data base (JCPDS: 06–4365) looks like a typical primitiveperovskite oxide; and the appearance of a weak peak at 2� ∼37.9◦

may indicate that some extent of B-site ordering may happen in thisstructure. However, it is noticed that there are several small peaksmainly for 2� range between 20◦ and 30◦, but also for 2� of 50◦.These diffraction peaks did not correspond to the perovskite struc-ture, and were indexed by an impurity of SrNb6O16 and Sr2Nb2O7that was excluded in the refinement as the secondary phase. All thepeaks in the diffraction pattern for Sr2CrNbO6 at room temperaturecould be indexed by an ordered perovskite structure in cubic spacegroup Fm3m (No. 225) with a doubled unit cell parameter of a sim-

Fig. 1. RT X-ray diffraction patterns of Sr2CrNbO6 obtained at 1200 ◦C in air atmo-sphere, showing observed (continuous line), calculated (circles). The short verticallines indicate the angular position of the allowed Bragg reflections. At the bottomin each figure the difference plot, Iobs − Icalc, is shown.

ple perovskite lattice. The results are similar to those of early reportsobtained by Choy et al. [20]. The two B-sites are shared by bothchromium and niobium with Cr-rich 4a sites and Nb-rich 4b sites.The refinement of oxygen site occupancy was not very conclusive

Sr2CrNbO5.873 in H2 at 700 ◦C. The reduced sample regained littleweight on cooling from 700 ◦C and this may be explained by thereoxidization of the sample. In the overall reduction cycle, the totalweight loss is 0.42 wt%.

Table 1Structure and profile parameters for Sr2CrNbO6 at room temperature

Atom Site x y z Occupancy Uiso (A2)

O 24e 0.25 0 0 1 0.0485(3)Sr 8c 0.25 0.25 0.25 1 0.0601(2)Cr(1) 4a 0 0 0 0.76(4) 0.0108(6)Cr(2) 4b 0.5 0.5 0.5 0.24(4) 0.0032(4)Nb(1) 4a 0 0 0 0.24(4) 0.0108(6)Nb(2) 4b 0.5 0.5 0.5 0.76(4) 0.0032(4)

Symmetry, space group Cubic, Fm3mCell parameters a = 7.8761(0) A, V = 488.57(7) A3

2� (◦) 10–120 (0.02 steps)�the (g cm−3) 5.65Reflections 68Refined parameters 19R/Rw (%) 0.053/0.092, �2 = 3.516

T. Xia et al. / Materials Chemistry an

Fig. 2. TGA curves of Sr2CrNbO6 from room temperature to 700 ◦C in H2 (99.9%).The heating rate is 10 ◦C min−1, then holding at 700 ◦C for 30 min and cooling slowlydown to room temperature under flowing H2 at a rate of 50 ml min−1.

TGA itself cannot directly determine the phase stability of mate-rials at high temperature in the case of phase segregation. In orderto further investigate the crystal stability of Sr2CrNbO6 in a reduc-ing atmosphere, the as-prepared Sr2CrNbO6 obtained in air wasreduced in CO (99.95%, PO2 ≈ 10−20 atm) at 800 ◦C for 12 h andcooled naturally to room temperature in CO. The XRD pattern ofreduced Sr2CrNbO6 is shown in Fig. 3, which indicates that thereduced product is a single phase with a cubic structure and thereare no second phase (SrNb6O16 and Sr2Nb2O7) observed, suggest-ing that Sr2CrNbO6 is stable in reduced atmosphere at 800 ◦C.The similar profile fit was obtained when the same structuralmodel was used for Rietveld refinement. The oxygen site occupancyagain refined too close to 1 and was set to 0.996 corresponding toSr2CrNbO5.973 as indicated by TG measurement in CO. Table 2 sum-marizes the final refined structure data. It is noted that the latticeparameter a decrease slightly from 7.8796(0) to 7.8695(7) A. The cellvolume contracts 0.25% during the reduction, which may be causedby the loss of lattice oxygen. On the other hand, the reduction of Crfrom +3 to +2 and Nb from 5+ to 4+ may result in the lattice expan-sion, however, the contribution of M–O bond length expansion on

Fig. 3. RT X-ray diffraction patterns of Sr2CrNbO6 after reducing at 800 ◦C in CO(99.95%) for 12 h, showing observed (continuous line), calculated (circles). The shortvertical lines indicate the angular position of the allowed Bragg reflections. At thebottom in each figure the difference plot, Iobs − Icalc, is shown.

d Physics 111 (2008) 335–340 337

Table 2Structure and profile parameters for Sr2CrNbO6 reduced in CO (99.95%) at 800 ◦C for12 h

Atom Site x y z Occupancy Uiso (A2)

O 24e 0.2506(2) 0 0 0.996 0.0462(1)Sr 8c 0.25 0.25 0.25 1 0.0351(3)Cr(1) 4a 0 0 0 0.70(3) 0.0026(2)Cr(2) 4b 0.5 0.5 0.5 0.30(1) 0.0031(1)Nb(1) 4a 0 0 0 0.30(1) 0.0039(2)Nb(2) 4b 0.5 0.5 0.5 0.70(3) 0.0036(1)

Symmetry, space group Cubic, Fm3mCell parameters a = 7.8695(7) A, V = 487.35(1) A3

2� (◦) 10–120 (0.02 steps)Reflections 68Refined parameters 19R/Rw (%) 0.039/0.076, �2 = 2.917

reduction is often exceeded by opposing contributions of factorssuch as lattice oxygen loss leading to the contraction.

The morphology and element distribution of the samples wereanalyzed by a field-emission scanning electron microscopy withan energy-dispersive spectroscopy analysis. The SEM images of thesintered pellets before and after the reduction in CO are shown inFig. 4. It is found that the as-sintered sample is compact with parti-cle size 0.5–1 �m (Fig. 4a and b). After the reduction, the surface ofmaterial exhibits the increment of porosity (Fig. 4c and d). Fig. 4bshows the agglomeration tendency of the particles compared tothat in Fig. 4d. From the EDS analysis, there are no significantchanges observed after the reduction among the constituent ele-ments.

The valence state of chromium ions is measured by XPS. Fig. 5ashows a representative of the Al K� excited core level spectrumof Sr2CrNbO6 in the Cr2p region. The peaks for Cr2p are typicalbroad doublets. As seen in Fig. 5a, the binding energy is 575.5 eVfor Cr2p3/2 and 585.6 eV for Cr2p1/2 compared to the early reports[21]. The spin-orbit splitting is therefore 10.1 eV. The charge trans-fer (2p → 3d) signals of Cr2p3/2 are seen at 581.8 eV. The chromiumions in as-prepared sample are present as Cr(III). After the reductionat 800 ◦C in CO, the peaks shift of 1.1 eV is observed to the low bind-ing energy side of the Cr2p3/2,1/2 peaks, which is attributed to thereduction of chromium ions and corroborates the TG result. For theas-prepared material, the peaks for Nb3d5/2 and Nb3d3/2 are 206.4and 209.2 eV, respectively (Fig. 5b). The niobium ions are proposedto be Nb(V) according to the literature (206.6 eV for Nb3d5/2) [22].

There are no obvious peak shift observed for the reduced product,this gives an indication of no lower valence niobium ions after thereduction.

The conductivity of Sr2CrNbO6 prepared in air was measuredby a.c. impedance spectra and d.c. four-probe methods. The a.c.impedance spectrum of Sr2CrNbO6 at 255 ◦C in air is shown inFig. 6a. It indicates that the electronic conduction is dominantbecause only one arc was observed [14]. The semicircle is broad-ened and is associated with a capacitance of around 1 nF calculatedwith the equation of ωRC = 1 (where ω = 2�f and ω is the angularfrequency). The arc relates to both bulk and grain boundary resis-tances, with the contribution of grain boundary being dominantat temperatures where the arc is observed. It is convenient to plotthe impedance data as the real part of impedance vs. imaginary partdivided by the frequency (Z′′/f) as shown in Fig. 6b. The linear regionin such plots indicates a relaxation process, with the slope and they-intercept of the line giving the relaxation frequency and a char-acteristic resistance, respectively [23]. In Fig. 6b the Z′ vs. Z′′/f plotof Sr2CrNbO6 at 255 ◦C shows two linear regions, indicating thepresence of two relaxation process. In polycrystalline conductingmaterials the relaxation process could correspond to the bulk (fb),

338 T. Xia et al. / Materials Chemistry and Physics 111 (2008) 335–340

rther

Fig. 4. FE-SEM micrographs of Sr2CrNbO6 (a and b) formed at 1200 ◦C; (c and d) fupellets before and after the reduction in CO are shown in (e) and (f), respectively.

the grain boundary (fgb) or the electrode process (fel). Calculationof the capacitance gives a value (1 × 10−9 F), which is typical of thegrain boundary capacitance [24]. These impedance data were car-

ried out nonlinear curve fitting using an equivalent circuit model.For the temperature studied in air, there was no possibility to sep-arate bulk and grain boundary components, so only the total acconductivity was investigated. The activation energy is 0.59(2) eVfrom 800 to 255 ◦C. The total conductivities of Sr2CrNbO6 in airare 2.03 × 10−2 S cm−1 and 5.94 × 10−5 S cm−1 at 800 and 255 ◦C,respectively. The lower conductivity may be attributed to the exis-tence of pores and secondary insulating phases confirmed by theSEM and XRD analysis, respectively. In the literatures [25,26], ithas been reported that the microstructural defects, such as poresbehave very similarly to the grain boundary. Given the porousnature of the sintered specimen in the present study (relative den-sity of ∼86%), Bruggeman Asymmetric medium theory can be usedto describe the ceramic materials having both high conductivity(Sr2CrNbO6) and low conductivity (pores) phases interconnectedin three dimensions (3D). According to the following equation:

�m = �h(1 − f )3/2 (1)

where �m is the measured conductivity in the present experiment,�h is the conductivity of high conductivity phase (Sr2CrNbO6) and

heating at 800 ◦C in CO (99.95%) for 12 h. EDS results obtained for the surfaces of

f is the volume fraction of low conductivity phase (pores), themeasured conductivity will decrease with increasing the volumefraction of pores. Therefore, the more pores result in the decrease

of conductivity of conducting material.

The previous results have shown that the secondary insulatingphases which here will be called microstructure defects behavevery similar to the grain boundary [26,27]. A parallel model viewsthe insulating phases as blockers which statistically immobilizea fraction of charge carriers [26]. The insulating phases can dis-courage the conducting connectivity between conducting particles,which is called the blocking process. These insulating phases ordi-narily tend to segregate to grain boundaries in order to lower thestrain and electrostatic energy of the system. This might decreasethe grain boundary conductivity of solid materials. Thus, this modelwell supports for the interpretation of results obtained for thepresent material. On the other hand, the effective medium theory(EMT) also demonstrates that the secondary insulating inclusionscan block the mobility of charge carriers and lower the conductivityof conducting material [28,29].

The electrical conductivity of Sr2CrNbO6 was also investigatedby the d.c. four-probe technology. The conductivity of Sr2CrNbO6 ishigher in N2 than in air. The total conductivities in N2 and air weremeasured as 3.19 × 10−2 S cm−1 and 2.43 × 10−2 S cm−1 at 800 ◦C,respectively. At low temperature, e.g. 300 ◦C, the total conductivity

T. Xia et al. / Materials Chemistry and Physics 111 (2008) 335–340 339

Fig. 6. Impedance spectrum (255 ◦C) of Sr2CrNbO6 sintered at 1200 ◦C for 8 h; (a)complex impedance plane plots. The solid line in (a) is the nonlinear fit to theimpedance data using an equivalent circuit as shown in the circuit diagram. (b)Z′ vs. Z′′/f representation. The solid lines in (b) are the fits to the linear regions of thespectrum. R represents the resistance and CPE means the constant phase elements.

increases slightly with decreasing p(O2), indicating that the mate-rial exhibits a n-type conduction. The electrons for the conductionmay come from loss of lattice oxygen and creation of oxygen vacan-cies. This can be expressed by the following equations [31,32]:

O×O ⇔ 1

2O2 + 2e′ + V ··

O, (2)

Hence, neglecting the mobility activation energy for electrons,the classical power dependence should be � ∼ [e′]PO−1/4

2 . The cre-ated electrons may be trapped by chromium, which increases theconcentration of charge carriers, resulting in the increment of elec-

Fig. 5. Al K� excited core level spectra for as-prepared Sr2CrNbO6 and after reducingat 800 ◦C for 12 h in CO (99.5%): (a) in a Cr2p binding energy region and (b) in a Nb3dbinding energy region.

is 1.11 × 10−3 S cm−1 in N2, which is almost one order of magni-tude higher than that in air (1.15 × 10−4 S cm−1). The conductivitiesand activation energy (Ea) of Sr2CrNbO6 in various atmospheresare shown in Fig. 7. The activation energy and conductivity in airmeasured by a.c. and d.c. methods are consistent with each other.In CO atmosphere, the total conductivity deceases and the study

is described below. The slope change at around 535 ◦C on the con-ductivity curve in CO can be explained by the change in the extendof reduction and is consistent with the TG analysis at the sametemperature.

According to the early reports, the conductivity of Sr2CrNbO6may involve the 3d orbitals, then the energy level configurationsfor chromium ions are: Cr3+ : t3

2ge0g, Cr4+ : t2

2ge0g, and Cr2+ : t3

2ge1g.

As illustrated by Tuller [30], in perovskite compound, the 3eg orbitalof B-site cations, e.g. Crn+, may overlap with the nearby 2p� orbitalfrom the split O2p to form � bonds (eg–p�–eg bond). The 3t2gorbital of Crn+ also may overlap with the 2p� of O2− to form theweaker �-bonds (t2g–p�–t2g bond). In addition, the other possi-bility of d-orbital interactions may be the direct t2g–t2g overlap,which is unlikely due to the long Cr–Cr distance. It is supposedthat electrons or electron holes may hop from one Crn+ to a nearbyCrn+ through the � or � bonds, which results in the electronconduction. The intrinsic conduction process may be written as:Cr3+ + Cr3+ → Cr4+ + Cr2+ [21].

Fig. 8 shows the electrical conductivity of Sr2CrNbO6 as a func-tion of oxygen partial pressure p(O2) at 800 ◦C. At p(O2) valuesvarying from 1 to 10−3 atm, it is found that the conductivity

Fig. 7. Arrhenius plots of conductivities for Sr2CrNbO6 sintered at 1200 ◦C for 8 h invarious atmospheres by a.c. impedance and d.c. four-probe methods. The table inthe inset shows the activation energies (eV) in the different atmospheres. Ea(1) athigh temperature (above 535 ◦C) and Ea(2) at low temperature (below 460 ◦C).

340 T. Xia et al. / Materials Chemistry an

[

219–228.

Fig. 8. Isothermal conductivity (800 ◦C) of Sr2CrNbO6 sintered at 1200 ◦C for 8 h as afunction of oxygen partial pressure (log(PO2) = 0 to −16). The dotted line illustratesthe expected response outside these regions.

tronic conductivity. As shown in Fig. 8, the conductivity decreaseswith decreasing p(O2) when p(O2) values below 10−7 atm. Theconsideration which may influence the electronic and ionic con-ductivity is the concentration of oxygen vacancies. With furtherdecreasing the p(O2) and increasing oxygen deficiency, the 3D hop-ping path for electrons or holes through Cr–O–Cr close contactsmight be affected because some are blocked by oxygen vacan-cies leading to the lower conductivity which is consistent with theobserved change of conductivity shown in Figs. 7 and 8. On theother hand, with the increase of Vo defects, the oxide ion conduc-tivity may increase if the formation of defect clusters and orderingof oxygen vacancies are negligible.

In the conditions of a small variation of the total oxygen content,electron hole conductivity should decrease with the oxygen partialpressure as � ∼ [h·] ∼ PO1/4

2 . For the porous specimens, the equi-libration kinetics is faster. The conductivity of this material alsoexhibits an oxygen partial pressure dependence which in O2 + N2and CO + CO2 mixtures in Fig. 8 can be approximately described asbeing proportional to P(O2)±1/25. The similar oxygen partial pres-sure dependencies have been obtained by Venugopal et al. [33], andHolt and Kofstad [34]. This is an unusual oxygen pressure depen-

dence of oxides and it is hard to explain in terms of conventionaldefect equilibra. Furthermore, one may suggest that the defect sit-uations that are reached under these conditions reflect the slowlychanging state, e.g. slow equilibration kinetics. It should be notedthat only electronic hopping for chromium ions is considered inthis model, but a small residence of electrons on niobium ionscould affect the conductivity. There are no apparent dependences ofconductivity upon oxygen partial pressure observed under oxidiz-ing conditions, which would be consistent with ionic conduction.Although the impedance data at low temperature suggest that elec-tronic conduction is dominant, possibly via the grain boundaryelement, oxide ionic conductivity will become more important athigh temperature in oxidizing atmosphere.

4. Conclusions

The double perovskite Sr2CrNbO6 was prepared by a solid-state reaction. XRD pattern shows that the perovskite phaseof Sr2CrNbO6 exists with small amounts of secondary phases.Sr2CrNbO6 has a primitive cubic structure with space group Fm3m(2 2 5). It is chemical stable in H2 and CO until at least 700

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d Physics 111 (2008) 335–340

and 800 ◦C, respectively. In a strong reducing atmosphere (CO),it remains the cubic structure with space group Fm3m (2 2 5),a = 7.8677(2) A and V = 487.02(6) A3. a.c. impedance and d.c. con-ductivity measurements indicate that electronic conduction isdominant in this compound. The conductivities of this material are3.19 × 10−2 S cm−1 and 2.43 × 10−2 S cm−1 at 800 ◦C in N2 and airatmosphere, respectively. The lower conductivity may be explainedby the existence of pores and secondary insulating phases. The con-ductivity exhibits a small oxygen pressure dependence [P(O2)1/25],and this may be interpreted to reflect slow equilibration kineticswith decreasing oxygen activity. On the other word, the electronor electron holes may hop through the Cr–O � bonds (eg–p�–eg)and/or � bonds (t2g–p�–t2g). The increasing oxygen vacancies instrong reducing atmosphere may make the hopping of electronsand electron holes more difficult, consequently leading to thedecrease of electronic conductivity. Sr2CrNbO6 might be a candi-date for oxygen permeation if the ion conductivity is reasonablyhigh, which need the intensive investigation.

Acknowledgements

We thank National Natural Science Foundation of China forfunding the NSFC Project (20701013), also thank the financialsupport received from Heilongjiang Education Department andHeilongjiang University. We are also indebted to Mrs. Meiye Li andShuyun Wang for FE-SEM and XRD measurements.

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