Solid State Ionics 226 (2012) 53–58

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Solid State Ionics

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Ionic and electronic conductivity of 3 mol% Fe2O3-substituted cubic Y-stabilized ZrO2

K.V. Kravchyk a,b, O. Bohnke b,⁎, V. Gunes b, A.G. Belous a, E.V. Pashkova a,J. Le Lannic c, F. Gouttefangeas c

a V. I. Vernadskii Institute of General and Inorganic Chemistry, 32/34 Palladina Avenue, 03680, Kyiv 142, Ukraineb Institut des Molécules et Matériaux du Mans (IMMM CNRS 6283), Oxides and Fluorides Department, Université du Maine Av O.Messiaen, 72085 Le Mans Cedex9, Francec Centre de Microscopie Electronique à Balayage et microAnalyses (CMEBA), Université Rennes 1, Beaulieu Av du Général Leclerc 35042 Rennes Cedex, France

⁎ Corresponding author. Tel.: +33 243833354; fax: +E-mail address: odile.bohnke@univ-lemans.fr (O. Bo

0167-2738/$ – see front matter © 2012 Elsevier B.V. Allhttp://dx.doi.org/10.1016/j.ssi.2012.07.014

a b s t r a c t

a r t i c l e i n f o

Article history:Received 6 December 2011Received in revised form 19 June 2012Accepted 20 July 2012Available online 10 September 2012

Keywords:Stabilized cubic ZrO2

Ionic conductivityElectronic conductivityIron oxide

The effect of the addition of a small amount of iron oxide (3 mol%) in Y-stabilized ZrO2 has been investigated. Theaging of the obtained compoundhas been studied and the sintering temperature has beendetermined. Thebulk andtotal conductivities of (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 have been studied by means of impedance spectroscopy inthe temperature range from 425 K to 775 K. The electronic conductivity has been studied by the Hebb-Wagnertechnique using a blocking Pt microelectrode. The investigation has been carried out in a wide range of oxygenactivity, 10−25baO2

b103, and from 770 K to 1020 K. These data have been compared to the compound withoutiron oxide, YSZ (ZrO2)0.90–(Y2O3)0.10. This study demonstrates that the addition of a small amount of Fe2O3

decreases the sintering temperature and increases the stability of the compound without increasing the electronicconductivity.

© 2012 Elsevier B.V. All rights reserved.

1. Introduction

Materials based on oxygen-conducting cubic and tetragonal phasesof ZrO2 are of significant scientific and practical interests due to theirapplication in different electrochemical devices, namely solid oxide fuelcells, oxygen sensors and oxygen separationmembranes [1–3]. Differentoxides have been used for stabilizing tetragonal and cubic ZrO2 phases atroom temperature (alkaline and rare-earth oxides), among which, themost accepted nowadays is yttrium oxide [4,5]. However, materialsbased on zirconium oxide stabilized by yttrium oxide (Y-stabilizedZrO2 or YSZ) have some disadvantages such as a high sintering temper-ature (T≥1600 °C) and a low chemical stability inwet atmosphere [6,7].

The effect of the addition of iron oxide (Fe2O3) to YSZ has beeninvestigated by several authors. It has been shown that the partialsubstitution of Y3+ ions by the smaller Fe3+ ions (rYc.n.6

3+ =0.892 Ǻ;rFec.n.63+ =0.645 Ǻ [8]) decreases the sintering temperature [9–11]. Ithas also been reported, by the authors of the present paper, that thechemical stability of YSZ can be improved in ambient air by the sub-stitution of Fe3+ ions for Y3+ ions [12]. These two properties arehighly beneficial to the stabilized zirconium oxide.

A part from these positive effects, it iswell known that partial substi-tution of Y3+ by isovalent or aliovalent ions influences the O2− iontransport properties of the stabilized ZrO2 ceramics. Several authorsinvestigated the effect of the addition of iron oxide on the electricalproperty of Y- or Ce-stabilized ZrO2 [10,13–19]. These studies are

33 243833506.hnke).

rights reserved.

mainly focused on ionic conductivity and much more rarely on theeffect of Fe2O3 addition on the electronic conductivity. Verkerk et al.[10,15] have shown that iron oxide decreases the bulk and grain bound-ary conductivities in YSZ, despite its positive effect on sintering temper-ature. On the other hand, Foschini et al. [18] mentioned that theaddition of iron oxide in Ce-stabilized ZrO2 increases the grain-boundary conductivity. These results seem rather contradictory butthese authors did not perform anymeasurements of the electronic con-ductivity. Boukamp et al. [13] have shown that iron oxide, added in agreat amount (27 mol% FeO1.5), increases the electronic conductivityof YSZ owing to the presence of mixed valence of Fe3+ and Fe2+ ions.These authors also showed that the introduction of a second phase inthe sintered compounds, as either an amorphous or a nanocrystallineoxide, substantially influences the electrical properties of YSZ.

The aim of this paper is then to investigate the effects of the additionof a small amount of Fe2O3 to the binary system (ZrO2)0.90–(Y2O3)0.10 bysubstituting Fe3+ for Y3+. Small amount of Fe2O3 may prevent theoccurrence of electronic conductivity while retaining the chemicalstability in air and the low sintering temperature. Mössbauer andX-ray diffraction studies performed, some years ago, on the solid solu-tion (ZrO2)0.90–(Y2O3)1−x–(Fe2O3)x by the authors of this papershowed that a pure cubic phase of ZrO2 is obtained for x≤0.03. At thiscomposition, iron exists only as Fe3+ ions in octahedral coordination[12]. According to these results, the chemical stability in air and theelectrical properties of two compounds, with the same amount ofZrO2, are carefully compared in this study: one compound with10mol% of Y2O3 and no Fe2O3, namely (ZrO2)0.90–(Y2O3)0.10, and asecond one with 7 mol% of Y2O3 and 3 mol% of Fe2O3, namely(ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03. The influence of a small amount

Fig. 1. Concentration dependence of cubic phase for (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 and(ZrO2)0.92–(Y2O3)0.08 powder samples, kept under ambient atmosphere, as a function oftime.

54 K.V. Kravchyk et al. / Solid State Ionics 226 (2012) 53–58

of ferric oxide addition on ionic and electronic conductivity has beenstudied.

2. Experimental

The powders of (ZrO2)0.90–(Y2O3)0.10 and (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 have been prepared by precipitation of cation hydrox-ides following the procedure previously described [12]. The precip-itates have been dried at 353 K and then annealed at 1673 K. Thesintering of the pellets has been performed for 2 h in air at 1873 K for(ZrO2)0.90–(Y2O3)0.10 and at 1723 K for (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03.Despite a lower sintering temperature, the compactness of the iron-substituted YSZ was higher (98%) than the one of YSZ (87%).

X-Ray Diffraction (XRD) measurements were carried out with aDRON 4-07 powder diffractometer (CuKα radiation, 40 kV, 18 mA).XRD patterns of the powders, after annealing at 1673 K, were collect-ed in the angular range 2Θ from 10° to 150° in a step-scan mode witha step size of 0.02° and a counting time of 10 s per data point. SiO2

(2Θ calibration) and Al2O3 (intensity standard) have been used asexternal standards [20].

Scanning electron microscopy (SEM) and energy dispersive spec-trometry (EDS) have been used to investigate the microstructure ofthe sintered pellets and to determine the chemical composition ofthe grain and grain boundaries of each compound.

Impedance spectroscopy technique has been used to determinethe bulk conductivity and the total conductivity of the sintered(ZrO2)0.90–(Y2O3)0.10 and (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 samples.In this paper, the term “total conductivity” means (bulk+grainboundary) conductivity. We would like to remind that compleximpedance spectroscopy allows determining the bulk and total con-ductivity of the samples due to all the mobile charge carriers. Thistechnique is not able to determine the nature of the charged carriersmoving in the bulk and in the grain boundaries (either ions or elec-trons). However this technique, thanks to the sweep in frequency,is able to distinguish between bulk, grain boundary and electrodeinterface relaxations. Impedance spectroscopy was carried out inthe frequency range from 2 MHz to 1 Hz using a 1260 FrequencyResponse Analyzer and a 1296 Dielectric Interface from Solartron.The measurements were performed in the temperature range from425 to 775 K under dry air in a two-probe cell (DataLine). It hasbeen checked that the electrochemical system remains linear up to250 mV (r.m.s) even at high temperatures; therefore a 100 mV(r.m.s) applied ac voltage was chosen. Sputtered Pt was used as elec-trodes. The samples had the following dimensions: diameter≈0.7 сmand 0.4 cm and thickness≈0.1 сm and 0.2 cm. Data analysis wasperformed with the Complex Non linear Least Squares fit (CNLS) soft-ware package “Zview” of Solartron (version 3.3b) based on the LEVMsoftware of J.R. MacDonald. In all the fitting, the χ2 value is ensured tobe less than 10−4.

The electronic conductivity of both samples (ZrO2)0.90–(Y2O3)0.10and (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 has been carried out by usingthe Hebb–Wagner polarisation technique [21,22]. A Pt microelectrode(contact radius of 120±20 μm) was used as the blocking electrodefor oxygen ions as suggested by S. Lübke and H. D. Wiemhöfer [23].It has been pressed against the sample. To make Pt electrode totallyblocking and to avoid any electrochemical reaction between O2−

from the sample and O2 from external gas phase, a glass sealing(Heraeus, type IP 041) was applied onto the entire surface of the pel-let around the tip of the microelectrode and also on the sample sur-face to totally encapsulate it and to prevent the oxygen exchangefrom the gas phase. The blocking character of the Pt microelectrodehas been checked thoroughly as described in the text. On the oppositeside of the pellet, a two-phase mixture of Cu2O/CuO was used as areversible O2− electrode. The pellet was prepared from reagent-quality materials in the molar ratio 1/2. The steady-state current–voltage measurements were performed in the temperature range

from 770 K to 1020 K under dry N2 atmosphere. The applied voltage(U) ranged from −880 mV to +380 mV in a stepwise of 10 mV. Thecurrent (I) was recorded under steady state conditions. In this potentialrange the YSZ compounds were electrochemically stable. All the exp-eriments have been carried out from low oxygen activity (negativepotential) to high oxygen activity (positive potential). This potentialallowed us to investigate the property of the studied oxides in a largerange of O2 activity (aO2

) from 10−25 to 103.

3. Results and discussion

Fe2O3 alone is not able to stabilize the cubic phase of ZrO2

(c-ZrO2). However the partial substitution of Fe3+ for Y3+ increasesthe Y2O3 efficiency as a stabilizer of c-ZrO2. In the binary ZrO2–Y2O3

system, c-ZrO2 is fully stabilized for a Y2O3 content of 8–9 mol%[24]. In the system ZrO2–Y2O3–Fe2O3, c-ZrO2 is fully stabilized for atotal content of Y2O3 and Fe2O3 of 10–12 mol% [12]. The increase ofthe content of aliovalent substituents, such as Y3+ and Fe3+, inc-ZrO2 may increase the ionic conductivity of the ceramic by increas-ing the number of oxygen vacancies. This would be favorable for theelectrical property of the ceramic.

A part from the increase of the number of oxygen vacancies, thesubstitution of Fe3+ for Y3+ influences the chemical stability of YSZin air. It is known that the structure of Y-stabilized ZrO2 degradeswith time [25]. The transformation from cubic (c-ZrO2) and tetrago-nal (t-ZrO2) phases of ZrO2 to monoclinic phase (m-ZrO2) is observed.According to Alekseenko et al. [26] the degradation of stabilizedY2O3–ZrO2 structure, from martensite (t-ZrO2) to monoclinic(m-ZrO2), can be explained by the yttrium segregation into thegrain boundaries with further interaction with water present in air.XRD investigations carried out on (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03and (ZrO2)0.92–(Y2O3)0.08 powders, annealed at T=1673 K, revealedthat phase composition of (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 powdersample does not change after aging for three years in air, as shownin Fig. 1. This was not the case for the (ZrO2)0.92–(Y2O3)0.08 powdersample, kept under the same conditions, which shows a continuoustransformation from the cubic phase (c-ZrO2) to the monoclinicphase (m-ZrO2). The stability of the high-temperature phases ofZrO2 in iron-containing systems may be explained by decreasing itshydrophilicity. This decrease may be induced by a decrease of yttriasegregation into the grain boundaries. This decrease of segregation byiron content is confirmed by SEM and EDS analyses performedon (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 and (ZrO2)0.90–(Y2O3)0.10 samples(Fig. 2 and Table 1). The analysis of the chemical compositionin the middle of the grains (referred as (1) in Fig. 2a for(ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 and (3) in Fig. 2b for (ZrO2)0.90–(Y2O3)0.10) and in their grain boundaries (referred as (2) inFig. 2a for (ZrO2)0.90–Y2O3)0.07–(Fe2O3)0.03 and (4) in Fig. 2b for

Fig. 2. SEM micrographs of sintered pellets of a) (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 andb) (ZrO2)0.90–(Y2O3)0.10 samples (the numbers on the micrographs indicate pointswhere SEM–EDS analyses were performed).

-Z "

(oh

m)

-Z "

(oh

m)

1 Hz

675 K

χ2 = 8.7 10-5

Z ' (ohm)

Z ' (ohm)

0

105

105

105

2x105

2x105

2.0x105

3x105 4x1050

2 MHz

1 Hz265 kHz

650 Hz

625 K

χ2 = 1.4 10-4

00

8x104

4x104

4x104

2 MHz

1 Hz

600 kHz

1.5 kHz

650 Kχ2 = 9.7 10-5

L R(b)

CPE(b)

R(gb)

CPE(gb)

CPE elect

Ws

55K.V. Kravchyk et al. / Solid State Ionics 226 (2012) 53–58

(ZrO2)0.90–(Y2O3)0.10) clearly showed the decrease of yttrium concentra-tion in the grain boundaries of the iron containing sample. These resultsare summarized in Table 1.

The substitution of Fe3+ for Y3+ has also an influence on thesintering temperature. A maximum compactness, i.e. ≥98% of thetheoretical value, has been obtained for the iron substituted samplesintered at 1723 K although a compactness of only 87% has beenachieved at 1873 K for (ZrO2)0.90–(Y2O3)0.10. Therefore iron oxide,even in small amounts (≈3 mol%), significantly decreases thesintering temperature of ZrO2 ceramics. As a consequence of thedecrease of the sintering temperature, the grain size of the ceramiccontaining iron is smaller than that of the Y doped-ZrO2 ceramic, asobserved in the micrographs of Fig. 2.

The substitution of iron oxide for yttrium oxide has therefore sev-eral positive impacts on the stabilized Y doped-ZrO2 ceramic: itincreases the substitution rate and then the number of oxygen vacan-cies, it prevents the aging of the powder in ambient air, it decreasesits sintering temperature and then limits the grain growth.

Table 1Grain size and yttrium concentrations in the middle of grains and in the grain bound-aries of (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 and (ZrO2)0.90–(Y2O3)0.10 samples.

Sample Area Points inFig. 2

Yttriumconcentration,mass %

(ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03Grain size≈200 to 500 nm

Grain 1 7.2Grain boundary 2 6.6

(ZrO2)0.90–(Y2O3)0.10Grain size≈1 to 2 μm

Grain 3 9.8Grain boundary 4 11.1

The electrophysical properties, i.e. bulk and total conductivitiesas well as electronic conductivity, have been determined on(ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 (98% compactness) and (ZrO2)0.90–(Y2O3)0.10 (87% compactness) samples. The bulk and total conductivitieshave been determined from impedance spectroscopy carried out underdry air, in the temperature range from 425 to 775 K with an applied acvoltage of 100 mV (r.m.s). Figs. 3 and 4 present typical impedance dia-grams, plotted in the Nyquist plane, and recorded at 625, 650 and 675 Kfor (ZrO2)0.90–(Y2O3)0.10 and (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 samplesrespectively. At these temperatures, the diagrams consist of twosemi-circles at high and intermediate frequency and a straight line(with an angle slightly greater than 45°) at low frequency. Thehigh-frequency semi-circle disappears from the experimental frequencywindow (2 MHz–1 Hz) as temperature increases. On the other handthe low frequency straight line disappears totally when temperature de-creases below 575 K as well as the intermediate semi-circle at evenlower temperature. It is the reason why the impedance diagrams,recorded at these three temperatures, have been chosen to be presentedin this paper. Impedance data analysis was performedwith the CNLS soft-ware “Zview” of Solartron. The electrical equivalent circuit, made of in-ductance (L), resistances (R) and constant phase elements (CPE), has

0

-Z "

(oh

m)

0

2 MHz

1 MHz

3 kHz

Z ' (ohm)1054x104 6x104 8x104

2x104

2x104

Fig. 3. Impedance diagrams for (ZrO2)0.90–(Y2O3)0.10 sample (f=1.37 cm−1) recordedunder dry air at different temperatures, 625 K, 650 K, 675 K (black dots). The electricalmodel used for the fitting of the experimental data is shown in the inset as well as theresult of the fitting as a line.

625 K

94 kHz

600 Hz

1 Hz

650 K

200 kHz

1.5 kHz

1 Hz

675 K

420 kHz

3 kHz

1 Hz

-Z "

(oh

m)

-Z "

(oh

m)

-Z "

(oh

m)

0

105

2x105

0

0

8x104

4x104

4x104

2x104

0

Z ' (ohm)

Z ' (ohm)

Z ' (ohm)

0

105

105

105 2x105

2x105

3x105 4x105

0

4x104 6x104 8x1042x104

Fig. 4. Impedance diagrams of (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 sample (f=0.27 cm−1)recorded under dry air at different temperatures, 625 K, 650 K, 675 K.

1000/T K-1

1.2 1.4 1.6 1.8 2.0 2.2 2.4 2.6

Sig

ma*

T (

S K

cm

-1)

10-8

10-7

10-6

10-5

10-4

10-3

10-2

10-1

100

101 770

K

(ZrO2)0.90 - (Y2O3)0.10

(ZrO2)0.90 - (Y203)0.07 - (Fe2O3)0.03

Fig. 5. Temperature dependences of bulk conductivity (black symbols) and total conductiv-ity (open symbols) for (ZrO2)0.90–(Y2O3)0.10 (triangles) and (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03(circles) samples. The stars indicate the bulk conductivity (black) and total conductivity(white) of Y-substituted ZrO2 reported by Guo et al. [27].

56 K.V. Kravchyk et al. / Solid State Ionics 226 (2012) 53–58

been used for all the samplefitting. It is shown in the inset of Fig. 3. A con-stant phase element (CPE) instead of a pure capacitance C is used in themodel to take into account the fact that the semicircles are not centeredon the real axis. They are indicative of a distribution of the relaxationtime constants. Two (R,CPE) elements in series are used to fit the imped-ance diagrams, one being related to the sample bulk [R(b)//CPE(b)] andthe other one to the sample grain boundary [R(gb)//CPE(gb)]. An induc-tance (L) is used to take into account the contribution of the electricalwires at high frequency. This impedance influences the diagrams only atlow temperature. Finally, the electrode effect at low frequency has beenmodelized by a CPE (CPEelect) in parallel with a finite length Warburgelement with a short circuit terminal (Ws) which corresponds to a diffu-sion process at the electrode/electrolyte interface in parallel with thecharge of the double layer. This diffusion process is clearly observed atand above 650 K. The choice of the fitting model is guided by the formof the impedance diagrams (a curvature is observed at high temperature)and by theχ2 value, which reflects the goodness of the fitting. Its value issmaller than 10−4 when a Warburg element is used instead of a CPEalone. Below 650 K a CPE alone is used for the fitting. Fig. 3 shows, as anexample, the impedance data (dots) and the fitting curves (lines), theχ2 values are indicated in the diagrams. This model perfectly applies inthe whole temperature range investigated, from 425 to 775 K and forall the samples.

From the values of R(b) and R(gb), the bulk conductivity and thegrain boundary conductance of the ceramics have been determined.

The total conductivity has been obtained from the fitting model.According to the values of the capacitance, the high frequencysemi-circle can be ascribed to the motion of the charged species intothe grains of the ceramic (Cg≈10−10 F) and the intermediate frequencyone to their motion into the grain boundaries (Cgb≈10−8 F). Finally,the straight line at low frequency can be related to the polarisation ofthe electrode interface (Celect≈10−5 F). Cg and Cgb remain almost con-stant in the temperature range investigated although Celect increaseswith temperature. This result confirms that the two first capacitancesare related to polarisations in the material and the last one to thepolarisation of the electrolyte/electrode interface with a blocking elec-trode. As shown in Figs. 3 and 4, the frequency of the maximum of theimaginary impedance increases as temperature increases meaningthat the charged species motion is thermally activated.

Fig. 5 shows the plots of the bulk conductivity (black symbols) andthe total (grain+grain boundary) conductivity (open symbols) as afunction of the inverse of temperature for the (ZrO2)0.90–(Y2O3)0.10and (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 samples. The total conductivityhas been obtained by normalizing the total resistance to the samplethickness and area. Bulk and total conductivities for both samples fol-low the Arrhenius law. It can be observed that the total conductivityis only slightly smaller than the bulk one meaning that the resistanceof the grain boundary is not very high and that sintering is effective.The activation energy of the conduction mechanism in the grains(Ea) can be determined from the Arrhenius law:

σ ¼ σ0

Texp� Ea

RT

� �: ð1Þ

The activation energy is very close for both samples, i.e. 1.19 (with Fe)and 1.17 eV (without Fe), confirming the O2− ion migration. Theappearance of oxygen conductivity can be explained by the formationof oxygen vacancies during the dissolution of Y2O3 and Fe2O3 in ZrO2.

57K.V. Kravchyk et al. / Solid State Ionics 226 (2012) 53–58

The dissolution of these oxides in ZrO2 (in small amounts) can bewrittenas follows, according to the Kröger and Vink notation:

Y2O3→2Y′Zr þ 3O�

O þ V ••O ð2Þ

Fe2O3→2Fe′Zr þ 3O�O þ V ••

O ð3Þ

where Y′Zr(or Fe′Zr) denotes a Y3+ ion (or a Fe3+ ion) on zirconium crys-tallographic sitewith a negative excess charge,OO

× denotes an oxygen ionon its regular lattice position andV ••

Oa vacancy in an O2− site with a dou-bled positive excess charge. For comparison, the value of conductivity of(ZrO2)0.92–(Y2O3)0.08, as reported by Guo et al. [27], has been indicatedby stars in Fig. 5 (black star=bulk conductivity, white star=total con-ductivity). It can be shown that the substitution of Fe3+ for Zr4+

decreases the bulk conductivity of one order of magnitude and increasesvery slightly the activation energy. This result can be linked to the unitcell volume of the oxide that decreases as substitution occurs. The unitcell volume decreases from 136.2 Ǻ3 for x=0 to 134.6 Ǻ3 for x=0.03after sintered at 1673 K [12]. This result confirms the substitutionprocess. It is in good agreement with the ionic radius of these ions,i.e. rZrc.n.84+ =0.87 Ǻ; rYc.n.6

3+ =0.892 Ǻ; rFec.n.63+ =0.645 Ǻ [8].The stability of (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 in reducing and

oxidizing conditions has been investigated by measuring the electronicconductivity as a function of oxygen partial pressure. The experimentalresults are compared with those of (ZrO2)0.90–(Y2O3)0.10 obtained withthe same method and in the same conditions of temperature and atmo-sphere. Activity instead of partial pressure is used in this work. In thisstudy, the standard state for the oxygen activities is taken as the oxygengas with a partial pressure of PO2

=1.013 bar. The electronic conductivityhas been carried out using the Hebb–Wagner polarisation technique[21,22] with blocking Pt microelectrode [23] in the temperature rangefrom 770 K to 1020 K. The goodness of the current–voltage curves isbased on the assumption that the oxygen ions are totally blocked at thePt microcontacts under steady state conditions. The sample was totallyencapsulated by a glass sealing, with a particular attention around thetip of the Ptmicroelectrode, in order to avoid any electrochemical reactionbetween the oxygen of the external gas phase and the oxygen ions of the

Fig. 6. Oxygen partial pressure dependence of the electronic conductivity of (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 at different temperatures and of (ZrO2)0.90–(Y2O3)0.10 at 870 K(inset).

sample. Any leakage of oxygen through the glass sealing may lead to apartial non-blocking character of the Pt electrode and therefore to an in-crease of the recorded steady state current. This point has been thorough-ly checked before collecting data by measuring the steady state currentunder dryN2 andunder dry air (PO2

≈0.20 bar) at 1020 K. The I–Ucurveswere perfectly similar particularly in the high oxygen activity (aO2 ) do-main, indicating that in these experiments the microelectrode is an ionblocking electrode. From the experimental I–U curves, electronic conduc-tivity, σe, can be calculated according to the following equation:

σe aO2

� �¼ 1

2πrdIdU

ð4Þ

where r refers to the radius of the Pt electrode that we assumed to bespherical.

The oxygen activity at the Pt blocking electrode has been calculatedby using the Nernst equation:

U ¼ RT4F

lna′O2

a″O2

!ð5Þ

where a′O2is the oxygen activity at the Pt blocking electrode; a″O2

is theoxygen activity at the reversible electrode/ceramic interface (thisvalue was taken from [23]). On the basis of the experimental I–U data,dependences of electronic conductivity on oxygen activity wereobtained at Т=770, 870, 970, and 1020 K in the oxygen activity thatrange from 103 to 10−25 or 10−18 depending on the working tempera-ture, as shown in Fig. 6. At low oxygen activities, the increase of elec-tronic conductivity is related to the appearance of n-type conductivityat the Pt microcontact according to the following relationship:

At low oxygen activity : OxO→V ••

O þ 2e′ þ 12O2 gð Þ: ð6Þ

Electrons are produced as a consequence of the oxidation of oxygenions at the Pt electrode. Therefore n-type conductivity (σn) appears inthe sample as well as oxygen vacancies V ••

O . The rate of the reactionincreases as oxygen activity decreases in agreement with reaction (6)which is driven to the right-hand side.

On the other hand, at high oxygen activities, the increase of elec-tronic conductivity is related to the appearance of p-type conductivityat the Pt microcontact according to the following relationship:

At high oxygen activity : V ••O þ 1

2O2 gð Þ→Ox

O þ 2h•: ð7Þ

Holes are produced as a consequence of the reduction of oxygenatom at the Pt electrode. Therefore p-type conductivity (σp) appearsin the sample and oxygen vacancy content decreases. The rate of thereaction increases as oxygen activity increases in agreement with reac-tion (7) which is driven to the right-hand side.

As long as we can assume that the concentration of electrons andholes is small compared to the dopant concentration, the number ofoxygen vacancies can be considered to remain constant. Therefore,according to relationships (6) and (7), the dependence of the concen-tration of electrons and holes, and then σn and σp, can be determinedas a function of the oxygen activity under steady state regime:

σp∝ h•� �∝a1=4O2

; σn∝ e′h i

∝a�1=4O2

; σe ¼ σn þ σp : ð8Þ

According to this classic defect model, the electronic conductivity σe

vs oxygen activity plots should display a slope of +0.25 in the highoxygen activity domain and a slope of−0.25 in the low oxygen activitydomain. Further at a given temperature, the electronic conductivityshould show a minimum separating the p-and n-type conductivities.At this particular oxygen activity, the conductivity of electrons andholes will be equal (σn=σp).

58 K.V. Kravchyk et al. / Solid State Ionics 226 (2012) 53–58

Fig. 6 shows that the electronic conductivity of (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 sample, is low, i.e. ≈10−7 S cm−1 at T=770 Kand≈10−6 S cm−1 at T=870 K in a wide range of oxygen activity(10−25baO2

b1). At T=970 K and 1020 K, a minimum of electronic con-ductivity around 10−6–10−5 S cm−1 is observed in the oxygen activityregion: 10−15baO2

b1. The comparison of the electronic conductivitiesof (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 and (ZrO2)0.90–(Y2O3)0.10 samples(inset of Fig. 6) does not evidence any increase of the electronic conduc-tivity at T=870 K, when 3 mol% of iron oxide is added to the compound.Boukampet al. [13] reported ahigh electronic conductivity of ZrO2–Y2O3–

Fe2O3 system with a high amount of Fe, i.e. 27 mol% of FeO1.5, and theynoted the presence of a second phase, α-Fe2O3. In the present investiga-tion the amount of iron oxide is low (3 mol% of Fe2O3). This small amountdoes not affect the electronic conductivity but has significant positiveeffects on the stabilization and aging of YSZ and on its sinteringtemperature.

A part from these first observations, the plots of Fig. 6 show sever-al peculiarities which are not in agreement with the classical defectmodel described above. At low oxygen activity, after a linear increaseof conductivity with a slope of 0.20 (close to the theoretical valueindicated by the solid line), a plateau is observed on the curves indi-cating that the concentration of electrons towards lower oxygen ac-tivities is limited. The plateau is well observed at 870 and 970 K. Apossible explanation is that the oxygen activity is no longer controlledby the voltage but is influenced by a trapping effect involvingredox-active ions such as Fe3+. The reduction of Fe3+ ions to Fe2+

at the Pt microcontact consumes the electrons formed by reaction (6)and limits their concentration. This reduction avoids the increase ofelectronic conductivity and decreases the slope of the σe vs aO2

plotas observed in Fig. 6. As oxygen activity decreases and since Fe3+ isnot able to diffuse into the material, all Fe3+ is reduced at the Ptmicroelectrode and therefore the trapping is no more efficient. Theslope of the curve slightly increases to reach 0.25, as expected.

At high activity of oxygen, in the p-type region, holes are formedaccording to Eq. (7). The electronic conductivity increases again(with a slope of only 0.10). Afterwards the conductivity does notreach a plateau but decreases meaning that a reaction takes placethat consumes the holes. Since Fe2+ is present at the Pt microcontact,we can assume that these ions are oxidized to Fe3+ and even to Fe4+.Once more this oxidation may limit the electronic conductivity in thep-type region. After all the ions are oxidized, reaction (7) only takesplace leading to an increase of the slope. These results suggest thatthe process of reduction (oxidation) of iron ions may take placealong with the incorporation (extraction) of oxygen into the oxideas oxygen activity decreases (increases). These results are in goodagreement with the analysis of Sasaki et al. [28] on the influence ofthe trapping effect of minor redox-active ions in YSZ including iron.These authors confirmed, by using EPR technique, the decrease ofFe3+ concentration (and increase of Fe2+ respectively) as oxygenpartial pressure decreases below 10−10 bar, in Y-stabilized ZrO2.They also determined the energy diagram of the Fe3+/Fe2+ redoxions showing the electrochemical potential of the electrons that liesaround the center of the band gap of YSZ corresponding to an oxygenpartial pressure around 10−12 bar.

4. Conclusion

The addition of a small amount of Fe2O3 (3 mol%) to Y-stabilizedZrO2 leads decreasing the sintering temperature and strongly increas-ing the stability of the cubic phase towards wet atmosphere. It has

been observed by SEM and EDS analyses that the presence of ironlimits the segregation of Y3+ into the grain boundary. This may notonly increase the chemical stability of the Zr–Y–Fe compound but alsoavoid the grain growth during sintering. We have shown by impedancespectroscopy that the bulk conductivity of (ZrO2)0.90–(Y2O3)0.07–(Fe2O3)0.03 decreases of one order of magnitude compared to the oneof (ZrO2)0.90–(Y2O3)0.10. The grain boundary resistance is howeversmall leading to a total conductivity very close to the bulk one. Despitethe presence of Fe3+ that can be reduced to Fe2+, the electronic con-ductivity remains low in a large oxygen activity domain. This has beendetermined by Hebb–Wagner technique using a blocking microelec-trode. The trapping of electrons by Fe3+ at low oxygen activity maylimit the n-type electronic conductivity by consuming electrons comingfrom the extraction of oxygen from the oxide. Therefore a plateau isobserved in the σe vs aO2

plot. This process is certainly reversible andFe2+ may be oxidized at high oxygen activity leading again to a limita-tion of the p-type electronic conductivity. The reduction–oxidation pro-cess of iron ions may then extend the oxygen activity domain intowhich the oxide has a minor electronic conductivity compared to theionic one.

Acknowledgments

K. V. Kravchyk would like to express gratefulness to the Universityof Maine for the financial support and for the opportunity to performpresent investigations in the Laboratory of Oxides and Fluorides ofthe University. This work has been done in the framework of PHCDNIPRO program.

References

[1] K. Kawamura, K. Watanabe, T. Hiramatsu, A. Kaimai, Y. Nigara, T. Kawada, J. Mizusaki,Solid State Ionics 144 (2001) 11.

[2] J.H. Lee, J. Kim, S.W. Kim, H.W. Lee, H.S. Song, Solid State Ionics 166 (2004) 45.[3] A.G. Belous, K.V. Kravchyk, E.V. Pashkova, O. Bohnke, C. Galven, Chem. Mater. 19

(2007) 5179.[4] H. Kaneko, F. Jin, H. Taimatsu, J. Am. Ceram. Soc. 76 (1993) 793.[5] S. Corman, V.S. Stubican, J. Am. Ceram. Soc. 68 (1985) 174.[6] S. Ran, L. Winnubst, W. Wiratha, D.H. Blank, J. Am. Ceram. Soc. 89 (2006) 151.[7] M. Yoshimura, T. Noma, K. Kawabata, S. Somija, J. Mater. Sci. Lett. 6 (1987) 465.[8] R.D. Shannon, Acta Cryst. A32 (1976) 751–757.[9] S. Lawson, C. Gill, G.P. Dransfield, J. Mater. Sci. 30 (1995) 3057.

[10] M.J. Verkerk, A.J.A. Winnubst, A.J. Burggraaf, J. Mater. Sci. 17 (1982) 3113.[11] V.G. Peichev, S.Y. Pliner, A.A. Dobizha, Steklo Keram. 3 (1991) 26.[12] A.G. Belous, E.V. Pashkova, K.V. Kravchyk, V.P. Ivanitskii, O.I. V'yunov, J. Phys.

Chem. C 112 (2008) 3914.[13] B.A. Boukamp, T.P. Raming, A.J.A. Winnubst, H. Verweij, Solid State Ionics 158

(2003) 381.[14] D.P.F. De Souza, A.L. Chinelatto, M.F. De Souza, J. Mater. Sci. 30 (1995) 4355.[15] M.J. Verkerk, B.J. Middelhuis, A.J. Burggraaf, Solid State Ionics 6 (1982) 159.[16] N. Matsui, M. Takigawa, Solid State Ionics 40 (41) (1990) 926.[17] U. Vohrer, H.D. Wiemhofer, W. Gopel, B.A. Hassel, A.J. Burggraaf, Solid State Ionics

59 (1993) 141.[18] C.R. Foschini, D.P. Souza, P.I. Filho, J.A. Varela, J. Eur. Ceram. Soc. 21 (2001) 1143.[19] T.S. Zhang, J. Ma, Y.J. Leng, S.H. Chan, P. Hing, J.A. Kilner, Solid State Ionics 168

(2004) 187.[20] In: Certificate of Analysis, Standard Reference Material 1976, Instrument Sensitiv-

ity Standard for X-ray Powder Diffraction, National Institute of Standards & Tech-nology, Gaithersburg, 1991, p. 1.

[21] M.H. Hebb, J. Chem. Phys. 20 (1) (1952) 185–190.[22] J.B. Wagner, C. Wagner, J. Chem. Phys. 26 (6) (1957) 1597–1600.[23] S. Lübke, H.D. Wiemhöfer, Solid State Ionics 117 (1999) 229.[24] H.G. Scott, J. Mater. Sci. 10 (1975) 1527.[25] A.G. Belous, A.N. Makarenko, E.V. Pashkova, B.S. Khomenko, Inorg. Mater. 35 (1999)

1341.[26] V.I. Alekseenko, G.K. Volkova, J. Tech. Phys. 70 (2000) 57.[27] X. Guo, R. Waser, Prog. Mater. Sci. 51 (2006) 151–210.[28] K. Sasaki, J. Maier, Solid State Ionics 134 (2000) 303.