evaluation of srti1−ynbyo3+δ materials for gas sensors

Ž .Sensors and Actuators B 56 1999 198–205

Evaluation of SrTi Nb O materials for gas sensors1yy y 3qd

J.C.C. Abrantes a, J.A. Labrincha b, J.R. Frade b,)

a Instituto Politecnico de Viana do Castelo, ESTG, Ap. 574, 4901 Viana do Castelo, Portugal´b ( )Ceramics and Glass Engineering Department UIMC , UniÕersity of AÕeiro, 3810 AÕeiro, Portugal

Received 25 June 1998; received in revised form 18 November 1998; accepted 23 November 1998

Abstract

Porous samples were used to evaluate the dependence of electrical conductivity on the oxygen partial pressure. Results obtained withthese porous samples revealed that Nb for Ti substitution suppresses the p-type contribution even in oxidizing conditions, and enhancesthe n-type conductivity. The conductivity of porous SrTi Nb O samples can be described by the power law dependence on py1r4,1yy y 3qd O 2

except possibly in very reducing conditions. The response time of porous samples is short, as required for resistive oxygen sensors.q 1999 Elsevier Science S.A. All rights reserved.

Keywords: Strontium titanate; Oxygen sensors; Porous materials; Defect chemistry

1. Introduction

Strontium titanate-based ceramic materials have a greattechnological potential, and remain stable in wide tempera-ture ranges. Earlier studies were related to the study of

Ž .dielectric properties at low temperatures , but high tem-w xperature applications such as resistive sensors 1–3 are

also promising. This type of application may rely on aŽtypical power law dependence of conductivity ss

y1r n.s p , expected for n-type conductivity in wide ranges1 O2

of values of oxygen partial pressure p . Donor dopantsO 2

w xsuch as La for Sr substitution 4,5 , or Nb for Ti substitu-w xtion 4,6 are expected to widen this range, even for

oxidizing conditions. Other applications such as electrodesw xfor solid oxide fuel cells 7,8 , or electrochemical mem-

w xbranes for oxygen separation 8 were also proposed re-cently, and understanding the defect chemistry at tempera-tures of about 10008C is essential for these high tempera-ture applications. Some room temperature properties mayalso be related to the high temperature defect chemistrybecause the concentrations of ionic defects may be frozenon cooling.

A combination of the dependence of conductivity on thew xoxygen partial, and Seebeck coefficient measurements 5

has been used to demonstrate that strontium titanate ismainly n-type in reducing conditions, but may become

) Corresponding author. Tel.: q351-34-370254; Fax: q351-34-25300;E-mail: [email protected]

p-type in oxidizing conditions, except possibly in donor-Ždoped materials. Addition of an acceptor species such as

.Fe or Mg for Ti substitution may enhance the concentra-w xtion of oxygen vacancies 8–10 , and is also expected to

yield mixed oxygen ion and hole transport, at least inoxidizing conditions. Alternatively, donor additives maybe used to obtain materials with n-type conductivity onlyŽ .rather than n- and p-type contributions .

Unfortunately, the response time of donor doped mate-w xrials may be long 4,11 , yielding an odd dependence of

w xconductivity on the oxygen partial pressure 4 ; this maybe due to slow response to changes in the atmosphere.However, the response time of thin films may be much

w x w xshorter 1–3 , and a preliminary work 6 confirmed thatthe response time of porous Nb-containing samples is alsomuch smaller than for dense samples.

Strontium titanate can dissolve high fractions of donorw xadditives such as lanthanum oxide 12,13 . Lanthanum

probably occupies mainly A-site positions, as suggested bythe similarity of cationic radii of La3q and Sr 2q, and

w xdemonstrated by computer simulation 14 . The oxygenw xstoichiometry 12,13 suggests that the difference in charge

between La3q and Sr 2q is compensated by negativelycharged strontium vacancies in oxidizing conditions, orelectrons in reducing conditions. On using the Kroger–Vink

Ž .notation, and assuming a source or sink for the excess ofstrontium oxide one may thus write

‘SrO’qV Ym1r2O q2eX qSr x 1Ž .Sr 2 Sr

0925-4005r99r$ - see front matter q 1999 Elsevier Science S.A. All rights reserved.Ž .PII: S0925-4005 99 00029-5

( )J.C.C. Abrantes et al.rSensors and Actuators B 56 1999 198–205 199

where ‘SrO’ denotes a generic sourcersink for the excessof SrO; this may be a Ruddelsden–Popper type of phase,

w xas suggested by several authors 12,13 . A similar mecha-nism may be proposed for Nb-containing materials.

It is also possible to obtain Sr Ti Nb O1yyr2 1yy y 3yd

w xcompositions with high Nb contents 7 , and the similarityof ionic radii of Ti4q and Nb5q suggests that Nb should

w xoccupy B-site positions 16 . On assuming that substitu-tions of Nb for Ti, and La for Sr yield similar effects onemay thus assume that strontium vacancies provide chargecompensation in Sr Ti Nb O , at least in oxidiz-1yyr2 1yy y 3yd

ing conditions, as described by:

1yyr2 SrOq 1yy TiO qyr2Nb OŽ . Ž . 2 2 5

™ 1yyr2 Sr x qyr2V Y q 1yy Ti x qyNb ØŽ . Ž .Sr Sr Ti Ti

q3O x ; 2Ž .O

The concentration of strontium vacancies is probablyŽ Žlower in compositions with unit molar A:B ratio, Sr: Ti

. .qNb s1 . In this case, electrons should play a majoreffect in charge compensation, especially in reducing con-ditions, and the relevant reaction may be described by:

SrOq 1yy TiO qyr2Nb O ™Sr x q 1yy Ti xŽ . Ž .2 2 5 Sr Ti

qyNb Ø q3O x qyeX qyr4O . 3Ž .Ti O 2

Otherwise, one may assume incorporation of the excess ofŽSrO in a Ruddelsden–Popper phase, as proposed for

w x.La-containing materials 12,13 , and the following changesin defect chemistry

SrOq 1yy TiO qyr2Nb O ™ 1yyr2 Sr xŽ . Ž .2 2 5 Sr

qyr2V Y q 1yy Ti x qyNb Ø q3O x qyr2‘SrO’Ž .Sr Ti Ti O

4Ž .

For relatively high Nb additions one may even expectsegregation of SrO-rich precipitates, as shown in the pre-sent work.

w xIrvine et al. 7 also found that reduction causes signifi-cant increase in lattice parameter of Sr Ti Nb O1yyr2 1yy y 3yd

Ž .materials at a relatively low temperature 9308C , suggest-ing changes in charge compensation even for materials

w xwith sufficient Sr deficiency. Irvine et al. 7 thus at-tributed the increase in lattice parameter to Ti3q androrNb4q ions with ionic radii exceeding both Ti4q and Nb5q.Significant fraction of Ti3q or Nb4q ions may also explainthe possibility of obtaining single phase SrTi Nb O1yy y 3

Ž w x.films with up to 100% Nb, by laser deposition 17 .Ž . Ž . Ž .Combination of Eqs. 3 and 4 also yield Eq. 1 , thus

describing the expected changes on moving from oxidizingto reducing conditions, as proposed for Sr La TiO1yx x 3qd

w xmaterials 12,13 . One should thus expect oxygen uptakeŽ .from the atmosphere on reoxidizing , and the slow re-

sponse of dense samples may thus be due to slow diffusionw xin the material. For example, Moos and Hardtl 15 claimed¨

that the concentration of vacancies remain frozen-in below

about 13508C. Nevertheless, our results show that theresponse times of very porous samples are much shorter.The results obtained with these porous samples were thusused to revise the defect chemistry of strontium titanatewith relatively high donor additions. This type of samplesmay also be suitable for gas sensors with relatively fastresponse to changes in the atmosphere.

2. Experimental procedure

ŽPowders with compositions SrTi Nb O ys1yy y 3qd

.0.01, 0.02, 0.05, and 0.1 , were prepared by solid statereaction from SrCO , TiO , and Nb O in the required3 2 2 5

proportions. The reactants were milled and mixed togetherwith ethanol and zirconia spheres in a nylon container.These powder mixtures were calcined at 11008C, for 15 h,and milled again in identical conditions to destroy agglom-erates formed during the calcination. These powders werecharacterized by X-ray diffraction, and revealed a cubicperovskite structure.

Small pellets prepared from calcined powders, wereuniaxially pressed and sintered at temperatures rangingfrom 12008C to 16008C, depending on the composition and

Ž .required density see Table 1 . Sintered samples were alsocrushed to obtain X-ray diffractograms needed to confirmthe structure and to evaluate the lattice parameter, and thetheoretical density. The real density was measured byimmersion in Hg, and the values obtained for the densestsamples compare to the corresponding theoretical density

Ž .values Table 2 .Ž .Scanning electron microscopy SEM and EDS micro-

probe analysis were used to study the effects of tempera-ture and composition on the microstructures of the sam-ples. The samples were polished and thermally etched for30 min, at a temperature which was 10% lower than the

Ž .corresponding sintering temperature in 8C . The SEMmicrostructures confirm the differences in density betweensamples shown in Table 1.

Table 1Ž .Firing conditions temperature and time of samples used in this work

Ž . Ž .Notation Composition T 8C t h

ST SrTiO 1600 203

Nb1A SrTi Nb O 1600 200.99 0.01 3.005

Nb1B SrTi Nb O 1400 40.99 0.01 3.005

Nb1C SrTi Nb O 1200 20.99 0.01 3.005

Nb2A SrTi Nb O 1600 200.98 0.02 3.01

Nb2B SrTi Nb O 1400 40.98 0.02 3.01

Nb5A SrTi Nb O 1600 200.95 0.05 3.025

Nb5B SrTi Nb O 1400 40.95 0.05 3.025

Nb5C SrTi Nb O 1480 40.95 0.05 3.025

Nb10A SrTi Nb O 1600 200.9 0.1 3.05

Nb10B SrTi Nb O 1400 40.9 0.1 3.05

( )J.C.C. Abrantes et al.rSensors and Actuators B 56 1999 198–205200

Table 2ŽSummary of relevant microstructural features relative density, average

.grain size, and seggregations of samples used in this work

Notation Relative density Grain size SeggregationsŽ . Ž .% mm

ST 99 f25Nb1A 99 f28Nb1B – f2 TiO -rich2

Nb2A 65 f0.2Nb5A 89 f0.3Nb10A 94 f1 SrO-richNb10B 60 f0.2

The electrical conductivity was measured with a four-wire DC method, at temperatures in the range 800–10008C,and as a function of the oxygen partial pressure. Thesemeasurements were performed in a furnace equipped withan electrochemical oxygen pump and a potentiometricoxygen sensor. The oxygen partial pressure was loweredby first flushing with N , and then switching on the2

Želectrochemical pump to attain reducing conditions typi-y20 .cally at least 10 atm at 8008C . The electrical measure-

ments were then recorded during a slow reoxidation of theatmosphere. Complete recovery of oxidizing conditionstook about 10 h.

A potentiometric oxygen sensor was used to monitorthe oxygen partial pressure. These sensors do not requirecalibration, and are reliable for conditions when the atmo-sphere acts as a buffer against the effect of minor leaksw x18 . Unfortunately, the buffer effect fails in a significant

Žintermediate range of oxygen partial pressure, typicallyy10 y3 .about 10 to 10 atm at 10008C , and the correspond-

ing results must thus be omitted. This effect is similar tothe sudden change of a pH signal on performing titrationsin aqueous media, which are very sensitive to minorchanges in concentration. Nevertheless, the sensor readingsare still reliable in the ranges of oxidizing, and reducingconditions.

The response time of SrTi Nb O samples was1yy y 3qd

evaluated by suddenly changing from reducing to oxidiz-ing conditions, and then measuring the time dependence ofconductivity.

3. Microstructural and structural characterization

The X-ray diffractograms of SrTi Nb O materi-1yy y 3qd

Ž .als Fig. 1 show a typical cubic perovskite structure,without any evidence of additional crystalline phasesŽ .within the detection limits of the equipment . However,some SEM microstructures reveal TiO -rich precipitates2

Ž .for 1% Nb Fig. 2B , and SrO-rich precipitates for 10% NbŽ .Fig. 2F . These Sr-rich precipitates could not be observed

Ž .on fresh fracture surfaces Fig. 2H .Ž .The microstructure obtained for 1% Nb Fig. 2B shows

Ž .segregation of a TiO -rich phase as revealed by EDS , but2

this could not be observed for samples with 2% Nb orhigher contents. The sintering behaviour of samples with1% Nb is consistent with results reported for

Ž w x.Sr Ti Nb O Cho and Johnson 19 . According1yx 1yy y 3" d

to these authors the sinterability and grain growth ofundoped and Nb-doped SrTiO can be promoted at tem-3

peratures exceeding an eutectic in the system SrTiO –3

TiO . Formation of a TiO -rich liquid phase may thus2 2

explain the sinterability and grain growth of sample Nb1AŽ . Žfired at 16008C for 20 h . Firing at 14008C below the

.expected eutectic temperature still yields dense samplesŽ .Fig. 2B , but the grain size remains much smaller than for

Ž .samples fired at 16008C above the eutectic .The titania-rich liquid phase or segregated TiO -rich2

particles suggest that the activity of titania is relativelyhigh in samples with 1% Nb. However, these findingscould not be observed for 2% Nb or higher contents, andthe activity of strontium oxide is thus likely to increase

Ž .with the fraction of Nb, as predicted by Eq. 4 ; this mayprevent the formation of a TiO -rich liquid phase, affecting2

the sintering behaviour of samples with 2% Nb. Note thatthese samples remain very porous even after sintering at16008C for 20 h, and the average grain size remains fairly

Ž .smaller than 1 mm Table 2 .Samples Nb10A show segregation of SrO-rich precipi-

tates after firing at 16008C for 20 h, and thermal etching at14408C. These SrO-rich precipitates suggest that the solu-bility of SrO may have been exceeded, and this may

Žaccount for the improved densification of the samples Fig..2F . However, one should not ignore other possible effects,

including the segregation of other species, at grain bound-w xaries. For example, Chiang and Takagi 16 reported Nb-

rich grain boundaries in SrTi Nb O .0.9 0.1 3qd

Fig. 1. X-ray diffractograms obtained for samples SrTi Nb O1y y y 3qd

Ž .samples with ys0.0; 0.01; 0.02; 0.05; and 0.1 after firing at 16008Cand for 20 h.


Ž . Ž . Ž . Ž . Ž . Ž . Ž . Ž . Ž .Fig. 2. SEM microstructures of: A Nb1A; B Nb1B; C Nb2A; D Nb5A; E Nb5B; F Nb10A; G Nb10B; H Nb10A fracture surface . SamplesNb1A, Nb2A, Nb5A and Nb10A were fired at 16008C, and thermally etched at 14408C. Samples Nb1B, Nb5B and Nb10B were fired at 14008C, andthermally etched at 12608C.


The average grain size of sample Nb10A remains muchsmaller than for samples Nb1A; this shows that the sinter-ing mechanisms of those compositions are different. Themechanism responsible for the sintering behaviour of sam-ples with 10% Nb may also play some effects on sampleswith 5% Nb. Note that the sintering behaviour of samples

Ž .with 5% Nb Fig. 2D and E is intermediate between thebehaviour observed for 2% Nb and 10% Nb.

4. Transport properties

Attempts to study the effects of the oxygen partialpressure on the transport properties of relatively denseSrTi Nb O samples yielded somewhat odd results1yy y 3qd

as shown in Fig. 3. The results obtained with densesamples deviate from any of the usual trends, and resemble

w xsome results reported for La-doped strontium titanate 4 .Fig. 4 demonstrates that the deviations observed for a

dense sample correspond to very long response times aftera sudden change from reducing to oxidizing conditions. Onthe contrary, the corresponding response time obtainedwith porous samples is short. The response times of stoi-chiometric SrTiO samples are short for both dense and3

porous samples.A simple diffusion controlled model may be derived for

dense samples on assuming that the average conductivityacross a flat sample is proportional to the average concen-

w xtration of a given species diffusing in the sample 6 . Inthis case, changes may be fastened by ensuring gas phasetransport through the open porosity, and exchange betweenthe gas phase and the surface of individual grains. The

Žtime scale for diffusion inside individual grains with.typical sizes in the range 1–10 mm thus becomes four to

six orders of magnitude smaller than the time scale re-

Fig. 3. Electrical measurements of samples with 5% Nb, at 10008C. Thevalues of relative densities are shown in the figure.

Fig. 4. Transient response of samples with 5% Nb after suddenly chang-Ž .ing from reducing to oxidizing conditions. A results for a sample with

Ž .relative density of about 89%; B results for a sample with relativedensity 64%.

quired for diffusion across dense samples with a typicalthickness of about 1 mm. Note that the kinetics of diffu-sion can be expressed as a function of a single dimension-less parameter tDrL2, where t is time, D is the diffusioncoefficient, and L is the size scale.

Nb for Ti substitution is expected to lower the concen-tration of oxygen vacancies thus explaining the differences

Ž .in transient time and size scale between dense undopedŽ .and Nb-containing samples. One may then use Eq. 1 to

describe the response of Nb-containing dense samples, onassuming oxygen uptake from the atmosphere, and slow

Ž Y .diffusion of strontium vacancies V towards the interiorSr

of the material. This hypothesis is supported by a trend tosegregate the excess of SrO at the surface, rather than inthe interior of the samples.

Sinksrsources of SrO may still be randomly distributedinside stoichiometric or acceptor doped SrTiO with suffi-3


cient concentrations of oxygen vacancies and strontiumw xvacancies, as proposed by other authors 20

V Y qVoq ‘SrO’mO x qSr x . 5Ž .¨Sr O Sr

In this case, reoxidation might be described by the usualexpression

1r2 O qVoq2eXmO x 6Ž .¨2 O

and the high mobility of oxygen vacancies in acceptorw xdoped SrTiO 9,10 may ensure short response times both3

for porous and dense samples in the actual temperatureŽ .range above 8008C .

The differences in response time between donor- andacceptor doped strontium titanate may be understood bytaking into account the expected drop in concentration of

Ž Ž .oxygen vacancies, caused by addition of niobium Eqs. 4Ž ..and 5 . The low concentration of ionic point defects thus

hinders a quick reequilibration on reoxidizing, and theoxygen chemical potential inside the material increasesslowly after a sudden change in the atmosphere. On insert-

Ž .ing this into Eq. 6 one expects a transient change inŽ .concentration of the main charge carriers electrons , and

the corresponding transient change in conductivity.Ž . Ž .A combination of Eqs. 5 and 6 may be nearly true

also for porous donor doped samples when O readily2

diffuses through the open porosity of the samples. In thisŽ . Ž . Ž .case, combination of Eqs. 5 and 6 still yields Eq. 1 .

The dependence of conductivity versus p may also beO 2

used to assess the correctness of the proposed mechanisms.In fact, the results for Nb-containing materials tend to apower law which is typical of materials with n-type be-haviour, even in oxidizing conditions. The correspondingresults for SrTiO also show a predominant n-type be-3

haviour in reduction conditions, but the behaviour in oxi-dizing conditions is clearly p-type. Nb additions are thusresponsible for suppressing the p-type contribution, as

w xreported for La-doped strontium titanate 5,21 . The n-typecontribution of Nb-doped materials increases by about oneorder of magnitude or even more because the differencesin Fig. 5 may be underestimated due to differencesin density of the samples. Note that these results wereobtained with a dense SrTiO sample, and porous3

SrTi Nb O samples.1yy y 3qd

The conductivity of samples with 5% Nb or 10% Nb isnearly pure n-type, and is described by a power law

s fs py1r4 7Ž .n n1 O 2

where s is the n-type conductivity at unit oxygen partialn1Ž .pressure. This dependence may be derived from Eq. 1 for

nearly constant concentration of strontium vacancies, andconstant activity of SrO.

Ž .The data for 1% Nb still deviate from Eq. 7 in veryreducing conditions, and also in oxidizing conditions. Sim-

Ž .ilar results were found at different temperatures Fig. 6 . InŽ .reducing conditions the deviations from Eq. 7 can be

Fig. 5. Conductivity measurements as a function of the oxygen partialŽ .pressure, at 9008C, for a dense stoichiometric sample ST , and porous

Ž . Ž . Ž .samples Nb1C 1% Nb , Nb5C 5% Nb , and Nb10B 10% Nb .

interpreted on assuming that electrons play a significantrole on the charge compensation, as reported for

w xSr La TiO materials 21 . One should thus combine1yx x 3qd

the neutrality conditionY Øw xnq2 V f Nb 8Ž .Sr Ti

Ž .with the mass action law of Eq. 1 obtaining1r2y1r4 y1r22nsY .a . p 1q0.25a pŽ . Ž .½O O2 2

y1r4y0.5a . p 9Ž . Ž .5O 2

where Y denotes the concentration of niobium and asŽ .1r20.5K a rY , K being the equilibrium constant of1 SrO 1

Ž .Eq. 1 , and a the activity of SrO. On assuming nearlySrO

constant a one may thus predict a limiting solutionSrO

nfY for sufficiently reducing conditions, and the usualpower law dependence for higher values of oxygen partialpressure. The n-type conductivity should vary with theoxygen partial pressure as predicted for the concentrationof electrons

1r2y1r4 2 y1r2sss p 1q0.25 s rs pŽ . Ž . Ž .½n1 O n1 L O2 2

y1r4y0.5 s rs . p 10Ž . Ž . Ž .5n1 L O 2

where s sm eY and s sm ea Y, e being the elec-L n n1 nŽ .tronic charge, and m the mobility. Eq. 10 thus predicts an

conductivity plateau for sufficient low values of oxygenŽ .partial pressure when s f s , as reported forn L

Ž w x.Sr La TiO at 13008C see Moos and Hardtl 21 .¨1yx x 3qd

Ž .Fig. 6 shows a slight deviation from Eq. 10 in oxidiz-ing conditions, probably due to the onset of a residualp-type contribution. This deviation can be described by a

1r4 Žtypical power law s fs p , where s denotes thep p1 O p12


Fig. 6. Conductivity measurements of a porous SrTi Nb O sam-0.99 0.01 3qd

Ž .ple Nb1C as a function of the oxygen partial pressure, at 8008C, 9008C,and 10008C.

.p-type contribution at unit oxygen partial pressure , andthe total conductivity may thus be rewritten

1r2y1r4 2 y1r2sss p 1q0.25 s rs pŽ . Ž . Ž .½n1 O n1 L O2 2

y1r4 1r4y0.5 s rs . p qs p 11Ž . Ž . Ž .5n1 L O p1 O2 2

Ž .Eq. 11 provides a reasonably good fitting for theconductivity of Nb-containing samples, as shown in Fig. 6;

Ž .this also confirms Eq. 1 , which was assumed in derivingŽ .Eq. 11 . The fitting parameter s measures the usualn1

n-type conductivity at unit oxygen partial pressure, and sLŽcorresponds to a limiting conductivity plateau for suffi-

.ciently reducing conditions .The results obtained for 5% Nb and 10% Nb fail to

Ž .show a clear conductivity plateau sfs in the workingL

conditions shown in Fig. 5. This may be explained bytaking into account that the limiting plateau is expected toincrease with the donor content, and might be displacedtowards more reducing conditions, as reported for

w xSr La TiO materials 21 .1yx x 3qd

Ž Ž ..The model Eq. 10 predicts that s should increasen1

with the donor content, as reported for La-doped strontiumw xtitanate 21 at about 13008C temperatures. However, this

is not demonstrated by our results at somewhat lowerŽ .temperatures Fig. 5 . Our results may be somewhat af-

Ž .fected by differences in sintering temperature Table 1 fordifferent samples. Note that the concentrations of ionicdefects are more likely to become nearly frozen when the

Ž .working temperatures TF10008C are much lower thanŽthe firing temperatures Ts1200, 1480 and 14008C for

.samples Nb1C, Nb5B and Nb10B, respectively . SampleNb1C was also slightly denser than samples Nb5C andNb10B.

5. Conclusions

Dense SrTi Nb O materials react slowly to1yy y 3qd

changes in the atmosphere even at temperatures as high as10008C. In this case, precise measurements should requirevery long equilibration times to avoid odd and meaninglessresults, except possibly at very higher temperatures. How-ever, this work shows that the response time of poroussamples is much shorter than for the corresponding densesamples. Porous samples were thus used to revise thetransport properties of these materials, and our results areconsistent with a defect chemistry model derived frommeasurements obtained for similar donor doped materialsŽ . w xSr La TiO 21 at temperatures of about 13008C.1yx x 3qd

Acknowledgements

This work was supported by The Portuguese Founda-Žtion for Science and Technology Program PRAXIS XXI,

.and Project 3r3.1rMMAr1760r95 .

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doped strontium titanate ceramics between 10008 and 14008C, J.Ž .Am. Ceram. Soc. 80 1997 2549–2562.

Ž .J.C.C. Abrantes graduated in Ceramic Engineering 1987 , and obtained aŽ .MSc degree in Materials Engineering 1991 from the University of

Aveiro, Portugal. He is a member of the academic staff of the PolytechnicInstitute of Viana do Castelo, Portugal. His main research interest are thetransport properties of ionic and mixed conductors, and materials forresistive sensors.

Ž .J.A. Labrincha graduated in Ceramic Engineering 1985 , and obtained aŽ .PhD degree in Materials Science and Engineering 1993 from the

University of Aveiro, Portugal. He is an Associate Professor at theUniversity of Aveiro. His research interests include novel materials forsolid oxide fuel cells, oxygen sensors, and other electrochemical applica-tions of oxygen ion or protonic conductors.

Ž .J.R. Frade graduated in Chemical Engineering 1978 from the UniversityŽ .of Coimbra, Portugal, and obtained a PhD degree 1983 from the

Ž .University of Sheffield, UK, and ‘Agregacao’ 1995 from the University˜of Aveiro, Portugal. He is a Full Professor at the University of Aveiro.His research interest include solid state reactions and related topics, thedevelopment and applications of high temperature oxygen ion and pro-tonic conductors, mixed conductors for electrochemical oxygen mem-branes, and semiconducting materials for gas sensors.

evaluation of srti1−ynbyo3+δ materials for gas sensors

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