CERAMICSINTERNATIONAL

Available online at www.sciencedirect.com

http://dx.doi.org/0272-8842 & 20

nCorrespondinBelgrade, Serbia

E-mail addre

Ceramics International 40 (2014) 9285–9292www.elsevier.com/locate/ceramint

Electrical properties of multidoped ceria

M. Stojmenovića,n, S. Boškovića, M. Žunićb,c, J.A. Varelac, M. Prekajskia, B. Matovića, S. Mentusd,e

aInstitute of Nuclear Sciences “Vinča”, Mihajla Petrovića – Alasa 12-14, 11001 Belgrade, University of Belgrade, SerbiabInstitute for Multidisciplinary Research, Kneza Višeslava 1, 11030 Belgrade, University of Belgrade, Serbia

cInstituto de Quimica, UNESP-LIEC, CMDMC, Rua Prof. Francisco Degni, 55, CEP 14800-900 Araraquara, SP, BrazildFaculty of Physical Chemistry, Studenski trg 12-16, 11158 Belgrade, University of Belgrade, Serbia

eSerbian Academy of Sciences and Arts, Knez Mihajlova 35, 11000 Belgrade, Serbia

Received 30 December 2013; received in revised form 28 January 2014; accepted 30 January 2014Available online 8 February 2014

Abstract

Multidoped nanosized ceria powders were prepared by either modified glycine nitrate procedure (MGNP) or self-propagating reaction at roomtemperature (SPRT). As the dopants to CeO2, trivalent rare earth oxides such as Nd2O3, Sm2O3, Gd2O3, Dy2O3 and Y2O3 were used, with thetotal molar fraction of 20%. The pressed powder pellets were subjected to the densification by sintering at 1500 1C, in an air atmosphere.A single-phase crystalline form was evidenced by X-ray diffractometry for both sintered materials. By means of complex impedancemeasurements, the conductivity of the sintered samples was determined as a function of temperature. At 700 1C, the conductivity amounted to2.19� 10�2 and 1.40� 10�2 Ω�1 cm�1 for the SPRT and for the MGNP sample, respectively. The corresponding values of activation energiesof conductivity amounted to 0.72 (MGNP) and 0.59 (SPRT) eV in the temperature range 550–700 1C.& 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Keywords: A. Sintering; C. Ionic conductivity; D. CeO2; E. Fuel cells

1. Introduction

Owing to its fluorite structure and valuable capabilities, thenanopowders of cerium (IV) oxide (CeO2) doped by trivalent rareearth ions such as Nd3þ , Sm3þ , Gd3þ , Dy3þ , Y3þ and Yb3þ

became a very interesting ceramic electrolytes for solid fuel cells(SOFC—solid oxide fuel cells) [1–4]. These powders displayed anenhanced ionic conductivity thanks to a high concentration ofoxygen vacancies [5]. The number of trivalent rare earth ions,introduced into crystal lattice of CeO2 by partial replacement ofCe4þ ions, is practically equal half the number of oxygen vacancies[6,7]. In previous studies, the oxides of rare earth elements such asLa [5,8–10], Sm [5,9–15], Gd [16–19], Dy [20,21] and Yb [22],and Y [23–25] were used as the dopants most frequently, thanks toa good solubility in ceria. In addition to the number of latticevacancies, some other factors have been found to influence the

10.1016/j.ceramint.2014.01.15114 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

g author at: INN Vinca, Lab. 170, P. O. Box 522, 11001. Tel.: þ381 11 340 8860; fax: þ381 11 340 8224.ss: mpusevac@vinca.rs (M. Stojmenović).

conductivity of ceria-based solid electrolytes. The lattice distortion,caused by the difference in ionic radii of Ce4þ and dopant ions wasfound to be one of the factors [26,27]. The reduction of meanparticle diameter and the unification of particle dimensionsimproved the conductivity of nanocrystalline materials [28]. Furtherto this, the reduction of mean particle diameter may acceleratesintering, and thus help to save energy in the production ofelectrolytes for fuel cells. The requirements for SOFC characteristicssuch as high efficiency, low operating temperature, and low costproduction process for obtaining pollutant-free ceramics, broughtdoped CeO2 into a focus of interest of scientific community.In our previous study [29] two different methods were used to

synthesize the doped CeO2 nanopowders (where the number ofdopants varied from 1 to 6). The relative fraction of each of dopantvaried, however, their total molar percent was kept constant at a levelof 20% (x¼0.2), since, according to the previous literature reports,the concentration of dopants of 10–20 mol% enabled the highestionic conductivities [27,30]. The physicochemical characteristics ofthe samples synthesized by two ways were analyzed in acomparative manner. The difference in either synthesis procedure

M. Stojmenović et al. / Ceramics International 40 (2014) 9285–92929286

or the number and concentration of dopants caused the difference incrystallite size, microstrain, specific surface area and porosity. Theincrease in dopant concentration usually reduced the mobility ofoxygen vacancies, and, consequently, the conductivity, too. Also,our recently published study [31] of sintered doped ceria samples ofthe constant composition, revealed that the samples with fivedopants displayed the highest ionic conductivity at 450 1C, whilethe addition of only one more dopant reduced the magnitude orderof ionic conductivity to that of pure CeO2. According to the authors'best knowledge, the conductivity of doped ceria with such a largenumber (5 or 6) of dopants has not been studied so far.

In accordance with the above mentioned facts, the subjectof this study was the ionic conductivity of the sinteredmultidoped ceria samples with five dopants in the temperaturerange 400–700 1C, with a particular attention to the depen-dence of conductivity on the density and the microstructure.

2. Material and methods

The experimental work was focused to the ceria based nano-powders with five dopants, having composition Ce0.8Nd0.01Sm0.04Gd0.04Dy0.04Y0.07O2�δ. The synthesis was performedby either modified glycine nitrate procedure (MGNP) [32] orself-propagating reaction in room temperature (SPRT) envir-onment [33].

The MGNP method involved the use of an aqueous solutionof δ-amino acetic acid-glycine (NH2CH2COOH) (Fluka), asolution of mixture of Ce(NO3)4 and Ce(CH3COO)4 in a moleratio 1:1, and a solution of nitrates [Me(NO3)3 � 6H2O] of Nd,Sm, Gd, Dy, Y (Aldrich, USA). The volume ratios provided toobtain the total mole fraction of dopants (� ) of 0.2 in the finalproduct. The reaction was carried out in a thoroughly cleanedsteel reactor. The concentration ratio of glycine vs. nitratessatisfied the equation:

4NH2CH2COOHþ½ð1�xÞCeðNO3Þ3 U6H2Og þxMeðNO3Þ3 U6H2O�þ2O2

-2Ce1� xMexO2� δþ22H2O↑þ5N2↑þ8CO2↑ ð1ÞThe prepared solution was heated first at 90 1C to remove

water, and further to 540 1C, to induce glycine ignitionaccompanied by nitrate and acetate decomposition. Theobtained dry powdery product was calcined at 600 1C over aperiod of 4 h to remove the traces of residual carbon. Theessence of this method of synthesis is that cerium originatedpartly from nitrate and partly from acetate, what provided arather mild decomposition reaction on heating, compared tothe agile original glycine-nitrate synthesis method [34]. Theyield of the product was very close (96–99%) to that calculatedon the basis of amounts of used salts.

In the SPRT method, the initial reactants were nitrate salts ofCe, Nd, Sm, Gd, Dy, Y (Aldrich, USA), and NaOH (Vetpromchemicals). The amounts of reactants needed for synthesis wascalculated using the equation:

2½ð1�xÞCeðNO3Þ3 U6H2OþxMeðNO3Þ3 U6H2O�þ6NaOHþð1=2�δÞO2

-2Ce1� xMexO2� δþ6NaNO3þ15H2O ð2ÞThe reactants were mixed in an alumina mortar for 15 min,

enabling rapid progress of the reaction at room temperature in air.

Then the mixture was kept in air for about 3 h, enabling thereaction to finish according to Eq. (2). The essence of this methodis an exothermic reaction in solid state, yielding the ceramicpowders at almost room temperature. The obtained suspension wasthen transferred in distilled water and subjected to centrifugation at3000 rpm, for 10 min by using a Centurion 1020D centrifuge.Rinsing procedure was repeated four times with water and twicewith ethanol. Then the synthesized nanopowders were subjected todrying in an oven at 100 1C. The resulting powder had theexpected stoichiometry. This method is pronouncedly cost-effective compared to the other methods.Green body compaction was done by cold pressing at

105 MPa. The sintering process was carried out in the tubefurnace at the temperature of 1500 1C in air atmosphere at aheating rate of 4 1C min�1. The density values were calculatedon the basis of geometric dimensions of the samples.The identification of crystalline phases was performed by

X-ray powder diffraction (XRPD) at room temperature usingSIEMENS D500 XRPD diffractometer and Cu-Kα1,2 radiation.The data for structure refinement were collected in the range of20–1001 2θ, using steps of 0.021 2θ and scan time of 12 s/step.The structure of powders was refined using computer programFullProf [35–37] which adopts the Rietveld calculation method.The TCH pseudo-Voigt profile function was used. To take intoaccount the instrumental broadening, the XRD pattern of astandard specimen CeO2 was fitted by the convolution of theexperimental TCH pseudo-Voigt function [38].For scanning electron microscopy (SEM) analysis, the electron

microscope model FE-SEM Jeol JSM 6330F (Japan) was used.The samples were pre-coated with a several nanometers thicklayer of gold before observation. For coating procedure, a deviceFine Coat JFC-1100 ION SPUTTER company JEOL was used.The images were recorded in SEI mode at a magnification� 5000 with the accelerating voltage of 10 kV. EDS analysiswas carried out at the invasive electron energy of 30 keV bymeans of QX 2000S device, a product of the company OxfordMicroanalysis Group. The maximum resolution was 0.4 nm.The electric conductivity was measured by the complex

impedance method in the temperature range 400–700 1C. Themeasuring cell was placed in a vertical furnace open to airatmosphere. To ensure good electrical contact, both sides of thesintered samples were coated with the silver paste. The temperaturewas increased steeply with the increments of 50 1C. The measuringdevice was a frequency response analyzer (FRA Solartron 1260Impedance/Gain Phase Analyzer) coupled with a dielectric inter-face (Solartron 1296). Operating frequencies were in the range0.1 Hz–5 MHz and the peak-to-peak AC voltage amplitude was50 mV. The impedance plots obtained experimentally were fittedby means of the software ZViews for Windows (Version 3.2b).

3. Results and discussion

3.1. Structure details of ceria samples

The phase composition of sintered samples was determinedby Rietveld refinement. The data reported by Kuemmerleand Heger [39] was used as starting structural model for the

Table 1Crystallite size (D), lattice parameters (a), microstrain (e), stacking factors (Rp), (Rwp), (RB) and (Ch

2) of sintered ceria ceramics obtained by the MGNP and SPRTmethods.

COMPOSITION, sintered samples a (Å) Ce–O (Å) D (nm) e (� 103Å) Rp Rwp RB Ch2

Ce0.8Nd0.01Sm0.04Gd0.04Dy0.04Y0.07 O2�δ (MNGP) 5.420336(3) 2.3470743(8) 283 4 19.80 11.40 2.17 3.77Ce0.8Nd0.01Sm0.04Gd0.04Dy0.04Y0.07 O2�δ (SPRT) 5.424358(6) 2.3488159(1) 811 10 20.60 14.90 2.48 5.66

Fig. 1. Plots of the final Rietveld refinements for the sintered samples:(a) MGNP synthesized Ce0.8Nd0.01Sm0.004Gd0.04Dy0.04Y0.07O2�δ powderand (b) SPRT synthesized Ce0.8Nd0.01Sm0.004Gd0.04Dy0.04Y0.07O2�δ powder,showing the comparison between observed (○) and calculated (full line)relative reflection intensities as well as their difference. The vertical barsindicate the positions of reflection lines for ceria structure.

M. Stojmenović et al. / Ceramics International 40 (2014) 9285–9292 9287

refinement. The results (lattice parameter (a), crystallite size (D),microstrain (e) and cation-anion bond length), including corre-sponding agreement factors, are presented in Table 1. Reasonablevalues of R factors indicate a successful refinement.

In spite of both high number (five) and high concentration(0.01oxo0.07) of dopants, sintered ceramics displayed asingle phase composition, as evidenced by X-ray measure-ments (Fig. 1). The lattice constant, ao, as well as the ⟨Ce–O⟩bond distance are almost constant in both the samples.

The only difference was noticed in crystallite size, namelylarger crystallites were observed for ceramics obtained by theSPRT method. This is in accordance with the properties of thestarting powders [29]. Namely, in the case of powders obtainedby the SPRT method, smaller mean crystallite size caused adistributed disorder generated by interrupting the periodicity ofcrystallites by interfaces, which in enlarged extent created bothshort and long range disorder. Therefore, the uniformity ofparticle packing displayed a pronounced influence on densifi-cation during sintering; the powders obtained by the SPRTwere more agglomerated and thus less suitable for sintering.

3.2. Densification and microstructure

The success of nanopowders consolidation is intimatelyrelated to the control of the competition between densificationand coarsening. The green structure, i.e., pore size distributionplays an important role in achieving high density. Densifica-tion is retarded or inhibited for wide pore distribution. In sucha case, big pores became larger while small pores shrink anddisappear.

The density of the samples after sintering is presented inTable 2. The samples synthesized by the MGNP method achievedhigher initial density (89.77%) than the samples obtained by theSPRT method (87.72%). The main reason of such differences liesin the fact that the particles of the samples obtained by the MGNPmethod remained more uniform after calcination. This yielded to abetter packing, lower inner activity and thus to a higher densityupon sintering procedure. On contrary, the samples obtained by theSPRT method were not subjected to calcinations, thus non-uniformparticles with much vacancies were produced. Consequently, theydelivered less compact particle packing, higher inner activity, andconsequently, lower sintering density.

This agrees well with the microstructural observation (Fig. 2).Although the sintered ceramics obtained from both SPRT andMGNP powders consisted of compacted polygonal crystals, thecurved grain boundaries indicate that the sintering in some pointswas not finished. The microstructures differ obviously due to thedifference in the synthesis method. Narrow size grains and

negligible porosity were characteristic for sintered materialsobtained from MGNP powders. On the other side, sintersobtained from SPRT powders displayed bimodal size distributionand some amount of pores. The fraction of small grains aroundlarger ones indicated that the mechanism of Ostwald ripening wasdominantly responsible for the grain growth observed.Having in mind that high sintering temperature was used;

one may conclude that the obtained ceria ceramics arecharacteristic of high thermal stability, which may be veryimportant from the point of view of its potential use as solidelectrolytes at elevated temperatures.

Table 2Geometric densities of the ceramics sintered during 1 h at 1500 1C in air atmosphere.

COMPOSITION,sintered samples

Theo. density[g/cm3]

m [g] D after sintering1500 1C 1 h [cm]

S after sintering1500 1C 1 h [cm2]

h after sintering1500 1C 1 h [cm]

Density after sintering1500 1C 1 h [g/cm3]

TD(%)

Ce0.8Nd0.01Sm0.04Gd0.04Dy0.04Y0.07O2�δ (MGNP)

7.016 0.2471 0.844 0.5595 0.07 6.3 89.8

Ce0.8Nd0.01Sm0.04Gd0.04Dy0.04Y0.07O2�δ (SPRT)

7.016 0.3467 0.946 0.7029 0.08 6.2 88.4

Fig. 2. SEM micrographs of doped ceria samples of the compositionCe0.8Nd0.01Sm0.004Gd0.04Dy0.04Y0.07O2�δ obtained by MGNP (a) and bySPRT and (b) methods, sintered at 1500 1C for 1 h in air atmosphere.

Fig. 3. Complex impedance plots of the sintered samples: (a) MGNP measuredat 450 1C and (b) SPRT measured at 400 1C in air atmosphere.

M. Stojmenović et al. / Ceramics International 40 (2014) 9285–92929288

It is reasonably to expect that the difference in microstruc-ture, namely smaller average grain size (AGS) and largernumber of grain boundaries of sintered MGNP sample incomparison to the sintered SPRT sample, discussed in thissection, and may cause the differences in their electricalproperties, which was considered in the next section.

3.3. Electrical conductivity

Original Nyquist plots (impedance diagrams) of the sinteredsamples having the composition Ce0.8Nd0.01Sm0.04Gd0.04Dy0.04Y0.07O2�δ obtained by both the MGNP and SPRTmethods, recorded within the temperature range of 400–700 1C, with the increments of 50 1C, are presented inFigs. 3 and 4.

At temperature 400 1C (Fig. 3(b)) for the SPRT sample andat temperature 450 1C for the MGNP sample (Fig. 3(a)),dominant part of the impedance diagram presented the high

frequency semicircle, originating from the parallel combinationof electrolyte resistance and (somewhat frequency dependent-distributed) geometric capacitance. In this range of tempera-tures the Nyquist plot did not indicate the separation betweenthe grain boundary and the intergranular impedance.The low frequency part of impedance diagram, being

usually linear in the systems predominantly capacitive incharacter, was not observed at lower temperatures. However,on rising temperature (see Fig. 4), the high frequencysemicircle (well visible at 400–450 1C) disappeared underprevailing effect of inductance of metal conductors (estimatedto be roughly 50 μH), while a low-frequency distorted semi-circle arose, which may be explained as a response of theresistance of oxygen oxidoreduction in combination with the

Fig. 4. Complex impedance plots of the sintered samples: (a) MGNP and(b) SPRT measured in the temperature range 400–700 1C in air atmosphere.The working temperatures are indicated at each diagram. The arrows indicatethe points on the real axis corresponding to the readings RbþRig.

M. Stojmenović et al. / Ceramics International 40 (2014) 9285–9292 9289

somewhat dispersed double layer capacitance. This explana-tion directed the way how to determine the electrolyteresistance from the Nyquist plots. In accordance with thisexplanation, the Nyquist plots were fitted with the equivalentcircuit with two constant phase elements (CPEs), given as theinsets in Fig. 3.The resistive component of the high temperature impedance

response may originate from the charge transfer resistance ofthe oxidoreduction process involving atmospheric oxygen andoxygen ions in solid electrolyte:

12O2þ2e�-O2� ð3Þ

In Table 3 the values of conductivities for samples MGNPand SPRT were presented at different temperatures. The valuesof conductivity are always higher for the SPRT sample. InTable 3 the ratio s(SPRT)/s(MGNP) is presented too. Onemay see that the difference in conductivity between the SPRTand the MGNP samples decreased with rising temperature.One may seek for the explanation of this behavior within thescope of the microstructure of the samples. Since the SPRTsample had lower density, smaller total surface area of grainsmay be the reason of its higher conductivity. At lowertemperatures, the influence of grain boundaries was morepronounced, causing also the higher ratio s(SPRT)/s(MGNP).As the temperature increased, the contribution of grainboundaries in the total resistance decreased, thus the ratios(SPRT)/s(MGNP) decreases too.Activation energies (Ea) of conductivity were calculated

from the Ahrrenius plots according to the equation:

lnðsTÞ ¼ ln A� Ea

k

1T

ð4Þ

where s is the conductivity, T is the absolute temperature, A isthe pre-exponential factor and k is the Boltzmann constant.In Fig. 5(a) and (b) the conductivity vs. temperature depen-

dence of MNGP and SPRT samples was converted into theArrhenius plots. Generally, the MGNP sample displayed higheractivation energy than the SPRT one. Some differences appear indifferent temperature regions. In the range 550–700 1C, activationenergy of conductivity 0.59 eV and 0.72 eV were found for theSPRT sample and MGNP sample, respectively. In the tempera-ture range 400–550 1C, activation energies of 0.82 eV and0.94 eV were determined for the samples SPRT and MGNP,respectively. Such temperature dependence of activation energycould be explained by the fact that the ionic conduction isthermally activated process. We should always keep in mind thatthe samples observed presented in chemical sense the samematerials prepared in different ways, thus only the differences inphysical properties may cause the differences in activation energy.Observing the microstructures of these samples (Fig. 2), one mayconclude that MGNP sample displayed smaller AGS, whichmeant larger number of grain boundaries. Although the grainboundary resistance was not perceived by EIS measurements athigh temperatures, it seems reasonable to assume that a largergrain boundary surface influences the activation energy of thesample MGNP.

M. Stojmenović et al. / Ceramics International 40 (2014) 9285–92929290

The typical charge transfer resistance values (Rct) of oxygenoxidoreduction reaction were obtained by fitting procedure.The values Rct were used for calculation of area specificresistance (ASR) using the following equation:

ASR¼ 12ðRchemAelectÞ ð5Þ

where Aelect is the electrode/electrolyte area of the cell and thefactor 1/2 accounts for the fact that a symmetric cell wasmeasured [40]. The values of ASR are directly dependent on

Table 3Conductivity of the samples synthesized by the MGNP and SPRT methods, measu

COMPOSITION,sintered samples

400 1C s[Ω�1 cm�1]

450 1C s[Ω�1 cm�1]

500 1C s[Ω�1 cm�1]

Ce0.8Nd0.01Sm0.04Gd0.04Dy0.04Y0.07O2�δ (MGNP)

2.14� 10�4 6.09� 10�4 1.50� 10�3

Ce0.8Nd0.01Sm0.04Gd0.04Dy0.04Y0.07O2�δ (SPRT)

6.46� 10�4 1.66� 10�3 3.75� 10�3

The ratio s(SPRT)⧸s(MGNP)

3.01 2.72 2.50

Fig. 5. Conductivities of the (a) MGNP and (b) SPRT samples presented in theform of Arrhenius plots.

the rate limiting step of the electrochemical processes involvedin the reaction at the electrode/electrolyte interface.In Fig. 6 the Arrhenius plot of ASR for MGNP and SPRT is

presented. The activation energies for the electrochemicalprocesses at the electrode/electrolyte interface are the samefor both samples (Ea¼2.86 eV) in the temperature range 550–700 1C. For the range 400–550 1C one may state that theactivation energies are mutually equal too within the limits ofexperimental error. These values are obviously, as expected,correlated not to the microstructure but only to the chemicalcomposition of materials.

red at different temperatures in air atmosphere.

550 1C s[Ω�1 cm�1]

600 1C s[Ω�1 cm�1]

650 1C s[Ω�1 cm�1]

700 1C s[Ω�1 cm�1]

3.47� 10�3 6.44� 10�3 1.03� 10�2 1.40� 10�2

7.07� 10�3 1.18� 10�2 1.66� 10�2 2.19� 10�2

2.03 1.83 1.61 1.56

Fig. 6. Arrhenius plots of ASR values for the samples: (a) MGNP and(b) SPRT measured in the temperature range 400–700 1C in air atmosphere.

M. Stojmenović et al. / Ceramics International 40 (2014) 9285–9292 9291

4. Conclusion

The Ce0.8Nd0.01Sm0.04Gd0.04Dy0.04Y0.07O2�δ nanopowderswere obtained by either modified glycine nitrate procedure orself-propagating reaction at room temperature. After pressinginto pellets and sintering by calcination at 1500 1C, the MGNPsample displayed more uniform, more densely packed parti-cles, with higher surface area of grains compared to the SPRTone. A unique fluorite structure characteristic of CeO2 wasevidenced regardless on the synthesis route.

The conductivity was measured in the temperature range400–700 1C by the complex impedance method. Higher valueswere measured for SPRT sample at all temperatures. For instance,at 700 1C, 2.19� 10�2 S cm�1 was found for this sample. Thedifference in conductivity between MGNP and SPRT sampleswas attributed to the difference in the surface area of grains. Thegrain boundaries present the barriers for ionic motion, which islarger for MGNP sample due to the larger total grain surface area.The contribution of grain boundaries to the total (grain boundaryþbulk) resistance decreased with the temperature increase,causing the decrease of the difference in conductivity betweenSPRT and MGNP samples.

The values of activation energy of conductivity, Ea, decreasedwith the temperature increase. The higher activation energy ofMGNP sample compared to the one of SPRT sample was foundat all temperatures. Although the impedance measurements do notallow distinguishing the role of grain boundaries at highertemperatures, we may suggest that higher surface area of grainscauses the Ea value of MGNP sample to be higher than that ofSPRT sample.

Acknowledgments

This work has been supported by the Ministry of Educationand Science of Serbia (Project no. III 45012) and Fundacão deAmparo à Pesquisa do Estado de Saõ Paulo-FAPESP (Projectno. 2010/20574-3).

References

[1] B.C.H. Steele, Materials for IT-SOFC stacks, Solid State Ion. 134 (2000)3–20.

[2] B.C.H. Steele, Material science and engineering; the enabling technologyfor commercialization of fuel cell systems, J. Mater. Sci. 35 (2001)1053–1068.

[3] O. Yamamoto, Solid oxide fuel cells: fundamental aspects and prospects,Electrochim. Acta 45 (2000) 2423–2435.

[4] S.M. Haile, Fuel cell materials and components, Acta Mater. 51 (2003)5981–6000.

[5] T. Mori, J. Drennan, J.H. Lee, J.G. Li, T. Ikegami, Oxide ionicconductivity and microstructures of Sm- or La-doped CeO2-basedsystems, Solid State Ion. 154–155 (2002) 461–466.

[6] J.R. McBride, K.C. Hass, B.D. Poindexter, W.H. Weber, Raman andX-ray studies of Ce1�xRExO2�y, where RE¼La, Pr, Nd, Eu, Gd and Tb,J. Appl. Phys. 76 (1994) 2435–2441.

[7] S. Tsunekawa, T. Fukuda, A. Kasuya, X-ray photoelectron spectroscopyof monodisperse CeO2�x nanoparticles, Surf. Sci. Lett. 457 (2000)L437–L440.

[8] T. Mori, J. Drennan, Y. Wang, J.H. Lee, J.G. Li, T. Ikegami, Electrolyticproperties and nanostructural features in the La2O3–CeO2 system,J. Electrochem. Soc. 150 (2003) A665–A673.

[9] T. Mori, J. Drennan, Y. Wang, J.G. Li, T. Ikegami, Influence of nano-structure on electrolytic properties in CeO2 based system, J. Therm. Anal.Calorim. 70 (2002) 309–319.

[10] T. Mori, Y. Wang, J. Drennan, G. Auchterlonie, J.G. Li, T. Ikegami,Influence of particle morphology on nanostructural feature and conduct-ing property in Sm-doped CeO2 sintered body, Solid State Ion. 175(2004) 641–649.

[11] Y. Wang, T. Mori, J.G. Li, Synthesis, characterization and sinterability of10 mol% Sm2O3-doped CeO2 nanopowders via carbonate precipitation,J. Eur. Ceram. Soc. 26 (2006) 417–422.

[12] Y. Wang, T. Mori, J.G. Li, Y. Yajima, Low-temperature fabrication andelectrical property of 10 mol% Sm2O3-doped CeO2 ceramics, J. Sci.Technol. Adv. Mater. 4 (2003) 229–238.

[13] N.K. Singh, P. Singh, D. Kumar, O. Parkash, Electrical conductivity ofundoped, singly doped, and co-doped ceria, Ionics 18 (2012) 127–134.

[14] T. Mori, T. Ikegami, H. Yamamura, Application of a crystallographicindex for improvement of the electrolytic properties of the CeO2–Sm2O3

system, J. Electrochem. Soc. 146 (1999) 4380–4385.[15] J. van Herle, D. Seneviratne, A.J. McEvoy, Lanthanide co-doping of solid

electrolytes: AC conductivity behavior, J. Eur. Ceram. Soc. 19 (1999)837–841.

[16] R. Buchanan, T. Mori, Y. Wang, D.R. Ou, J. Drennan, Fabrication ofdense GdxCe1�xO2�x/2 sintered bodies with nano-size grain and itsconducting properties, Mater. Trans. Japan 30 (2006) 955–958.

[17] T. Mori, J. Drennan, Y. Wang, G. Auchterlonie, J.G. Li, Influence ofnanostructural features on electrolytic properties of Gd doped CeO2 solidelectrolytes, J. Ceram. Soc. Japan 112 (2004) S642–S648.

[18] I. Riess, D. Braunshtein, D.S. Tannhauser, Density and ionic conductivity ofsintered (CeO2)0.82(GdO1.5)0.18, J. Am. Ceram. Soc. 64 (1981) 479–485.

[19] Y. Wang, T. Mori, J. Drennan, J.G. Li, Y. Yajima, Low temperaturesynthesis of 10 mol% Gd2O3-doped CeO2 ceramics and its characteriza-tion, J. Ceram. Soc. Japan 112 (2004) S41–S45.

[20] T. Mori, T. Kobayashi, Y. Wang, Synthesis and characterization ofnanohetero-structured Dy doped CeO2 solid electrolytes using combina-tion process of spark plasma sintering and conventional sintering, J. Am.Ceram. Soc. 88 (2005) 1981–1984.

[21] Y. Wang, T. Mori, J.G. Li, J. Drennan, Synthesis, characterization, andelectrical conduction of 10 mol% Dy2O3-doped CeO2 ceramics, J. Eur.Ceram. Soc. 25 (2005) 949–956.

[22] F. Ye, T. Mori, D.R. Ou, M. Takahashi, J. Zou, J. Drennan, Ionicconductivities and microstructures of ytterbium-doped ceria, J. Am.Ceram. Soc. 154 (2007) B180–B185.

[23] T. Mori, J. Drennan, Y. Wang, G. Auchterlonie, J.G. Li, A. Yago,Influence of nano-structural feature on electrolytic properties in Y2O3

doped CeO2 system, J. Sci. Technol. Adv. Mater. 4 (2003) 213–220.[24] D.R. Ou, T. Mori, F. Ye, M. Takahashi, J. Zou, J. Drennan, Micro-

structures and electrolytic properties of yttrium-doped ceria electrolytes:dopant concentration and grain size dependences, Acta Mater. 54 (2006)3737–3746.

[25] Y. Wang, T. Mori, J.G. Li, T. Ikegami, Y. Yajima, Low-temperaturepreparation of dense 10 mol%-Y2O3-doped CeO2 ceramics using powderssynthesized via carbonate coprecipitation, J. Mater. Res. Soc. 18 (2003)1239–1246.

[26] B.C.H. Steele, Appraisal of Ce1�yGdyO2�y/2 electrolytes for IT-SOFCoperation at 500 1C, Solid State Ion. 129 (2000) 95–110.

[27] H. Inaba, H. Tagawa, Ceria-based solid electrolytes, Solid State Ion. 83(1996) 1–16.

[28] J.P. Gonjal, R. Schmidt, J.E. Canuto, P.R. Alvarez, E. Moran, Increasedionic conductivity in microwave hydrothermally synthesized rare-earthdoped ceria Ce1�xRExO2�(x/2), J. Power Sources 209 (2012) 163–171.

[29] M. Stojmenović, S. Bošković, S. Zec, B. Babić, B. Matović, D. Bučevac,Z. Dohčević-Mitrović, F. Aldinger, Characterization of nanometricmultidoped ceria powders, J. Alloys Compd. 507 (2010) 279–285.

[30] M. Dudek, J. Molenda, Preparation and properties of CeO2-basedelectrolytes, Ceramics 84 (2004) 177–182.

M. Stojmenović et al. / Ceramics International 40 (2014) 9285–92929292

[31] M. Stojmenović, S. Bosković, D. Bučevac, M. Prekajski, B. Babić,B. Matović, S. Mentus, Electrical characterization of multidoped ceriaceramics, Ceram. Int. 39 (2013) 1249–1255.

[32] S.B. Bošković, B.Z. Matović, M.D. Vlajić, V.D. Kristić, Modifiedglycine nitrate procedure (MGNP) for the synthesis of SOFC nanopow-ders, Ceram. Int. 33 (1) (2007) 89–93.

[33] S. Bošković, D. Đurović, Z. Dohčević-Mitrović, Z. Popović,M. Zinkevich, A. Aldinger, Self-propagating room temperature synthesisof nanopowders for solid oxide fuel cells (SOFC), J. Power Sources 145(2005) 237–242.

[34] S. Mentus, D. Jelić, V. Grudić, Lantahnum nitrate decomposition byboth temperature programmed heating and by citrate gel combustion.Comparative study, J. Therm. Anal. Calorim. 90 (2007) 393–397.

[35] J. Rodriguez-Carvajal, FullProf computer program, 1998 ⟨http://charybde.saclay.cea.fr/pub/divers/fullprof.98/windows/winfp98.zip⟩.

[36] J. Rodriguez-Carvajal, Recent advances in magnetic structure determina-tion by neutron powder diffraction, Physica B 192 (1993) 55–69.

[37] J. Rodríguez-Carvajal, Recent developments of the program FULLPROF,in commission on powder diffraction (IUCr)., Newsletter 26 (2001)12–19.

[38] D. Balzar, N. Audebrand, M.R. Daymond, A. Fitch, A. Hewat,J.I. Langford, Size-strain line-broadening analysis of the ceria round-robin sample, J Appl. Crystallogr. 37 (2004) 911–924.

[39] E.A. Kuemmerle, G. Heger, The structures of C–(Ce2 O3þd), Ce7 O12

and Ce11 O20, J. Solid State Chem. 147 (1999) 485–500.[40] U.P. Muecke, K. Akiba, A. Infortuna, T. Salkus, N.V. Stus,

L.J. Gauckler, Electrochemical performance of nanocrystalline nickel/gadolinia-doped ceria thin film anodes for solid oxide fuel cells, SolidState Ion. 178 (2008) 1762–1768.