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Available online at www.sciencedirect.com

Journal of the European Ceramic Society 33 (2013) 2241–2250

Functional nanoceramics for intermediate temperature solid oxide fuel cellsand oxygen separation membranes

V. Sadykov a,∗, V. Usoltsev a, N. Yeremeev a, N. Mezentseva a, V. Pelipenko a, T. Krieger a,V. Belyaev a, E. Sadovskaya a, V. Muzykantov a, Yu. Fedorova a, A. Lukashevich a, A. Ishchenko a,

A. Salanov a, Yu. Okhlupin b, N. Uvarov b, O. Smorygo c, A. Arzhannikov d,M. Korobeynikov d, Ma.K.A. Thumm e,f

a Boreskov Institute of Catalysis, Novosibirsk State University, Novosibirsk, Russiab Institute of Solid State Chemistry, Novosibirsk, Russia

c Powder Metallurgy Institute, Minsk, Belarusd Budker Institute of Nuclear Physics, Novosibirsk State University, Novosibirsk, Russia

e Karlsruhe Inst. Technol., Karlsruhe, Germanyf Novosibirsk State University, Novosibirsk, Russia

Available online 4 February 2013

bstract

his work reviews results of research aimed at design and characterization of mixed ionic–electronic conducting perovskite–fluorite nanocompositexide ceramics. Nanocrystalline oxides were prepared via Pechini route, nanocomposites – via ultrasonic dispersion of their mixture in organicolvents with addition of surfactants. Genesis of the real structure of nanocomposites at sintering by conventional as well as advanced (microwave or-beam treatment) techniques was studied in details by structural methods. Applied preparation procedures ensured nano-sizes of perovskite/fluoriteomains even in dense ceramics and a high spatial uniformity of their distribution. Redistribution of elements between perovskite and fluoriteomains without formation of new phases was revealed. Characterization of nanocomposite transport properties by oxygen isotope heteroexchange

nd conductivity or weight relaxation demonstrated that perovskite–fluorite interfaces are paths for fast oxygen diffusion. Best perovskite–fluoriteombinations tested as cathode layers or dense oxygen separation layers in asymmetric supported membranes demonstrated performance promisingor the practical application.

2013 Elsevier Ltd. All rights reserved.

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eywords: Nanocomposites; Synthesis; Structure; Conductivity; Diffusion

. Introduction

Synthesis of inexpensive mixed ionic–electronic conductingMIEC) materials, including composites comprised of one elec-ronic conductor (perovskite-like oxides) and one good ionic

onductor (doped ceria, zirconia, etc.) is a very important taskn design of advanced cathodes of solid oxide fuel cells (SOFC)nd oxygen separation membranes.1–11 Traditional method of

∗ Corresponding author at: Boreskov Institute of Catalysis, Siberian Branch ofhe Russian Academy of Sciences, Prospect Akademika Lavrentieva, 5, Novosi-irsk 630090, Russia. Tel.: +7 383 3308763; fax: +7 383 3308056.

E-mail addresses: sadykov@catalysis.ru, sadykovy@academ.orgV. Sadykov), Uvarov@solid.nsc.ru (N. Uvarov), smorygo@rambler.ruO. Smorygo), arzhannikov@phys.nsu.ru (A. Arzhannikov),anfred.thumm@kit.edu (Ma.K.A. Thumm).

sfodairubeo

955-2219/$ – see front matter © 2013 Elsevier Ltd. All rights reserved.ttp://dx.doi.org/10.1016/j.jeurceramsoc.2013.01.007

omposites preparation by ball-milling a mixture of oxides1–4

ould not guarantee the uniform spatial distribution of parti-les of constituting phases required for a good percolation,nd, hence, high mixed ionic–electronic conductivity. Impreg-ation/infiltration of a porous electrolyte layer by perovskiteuspension or precursor solution7 is certainly more promisingor manufacturing nanostructured cathode layers, but controlf the spatial distribution of phases in composite could beifficult as well. Different versions of the sol–gel methodre potentially more efficient in providing required uniformntermixing of phases in nanocomposites.12,13 Polymerized cit-ic acid–ethylene glycol polyester precursor (Pechini) methodsed for synthesis of nanocrystalline complex oxides14 could

e applied for synthesis of nanocomposites as well. A pow-rful ultrasonic treatment of the mixture of nanocrystallinexides in solvents with addition of surfactants15 is promising

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242 V. Sadykov et al. / Journal of the Europ

or providing a uniform intermixing of both phases in nanocom-osites.

Another important problem in synthesis of MIEC compositess that high (up to 1200–1400 ◦C) temperatures required for sin-ering coarse-grained perovskite–fluorite composites into denseeramics often result in formation of new phases (La2Zr2O7yrochlore) with a low conductivity and oxygen mobility.16 Sin-ering of nanocrystalline oxides or their composites into denseeramics is known to proceed at much lower temperatures thusreventing undesired interaction between phases. New tech-iques of sintering based on application of radiation-thermalreatment by e-beams or microwave heating17,18 could furtherecrease the temperature of nanocomposites sintering into denseayers thus preventing formation of new phases with a low con-uctivity and oxygen mobility.

This work reviews results of research aimed at design andharacterization of MIEC nanocomposite oxide ceramics com-rised of perovskite-like (La1−xSrxMe1

1−yMe2yO3, Me1,2 =

n, Co, Fe, Ni; La1−xBixMnO3+δ, Pr2NiO4) and fluorite-ike (Sc0.2Ce0.1Zr0.79O2−y, Gd-doped ceria, Bi1.5Y0.3Sm0.2O3,iErO3−δ) phases prepared by using these new approaches.19–29

ffects of systems composition and sintering parameters onheir functional characteristics (transport properties and reac-ivity of surface sites) are considered. Examples of successfulesting of MIEC nanocomposites as SOFC cathodes and oxy-en separation layers in asymmetric supported membranes areiven.

. Materials and methods

Nanocrystalline oxides with perovskite-like and fluorite-liketructures were synthesized by Pechini route using metal (Me)itrates, citric acid (CA), ethylene glycol (EG) and ethylene-iamine (ED) as reagents. The molar ratio of CA:EG:ED:Meas 3.75:11.25:3.75:1. CA and metal nitrates were dissolved in

thylene glycol at 80 ◦C and in distilled water at room temper-ture, respectively. The prepared solutions were mixed togethert room temperature under stirring followed by addition ofD, stirring further for 60 min, heating to 70 ◦C and keep-

ng at this temperature for 24 h for the gel formation. The gelas calcined in the 500–700 ◦C temperature range under air to

emove organics and produce nanocrystalline oxides. Powderedc0.2Ce0.1Zr0.79O2−y electrolyte further referred to as ScCeSZas a commercial product of DKKK (Japan).24

The composites were prepared by ultrasonic treatment of theixture of powders in isopropanol with addition of polyvinyl

utyral using a T25 ULTRA-TURRAX (IKA, Germany)omogenizer. Suspensions were then either dried and pressednto pellets or supported by slip casting/air brushing as thin lay-rs on different substrates (anode half-cells.22,28 Ni–Al alloyoam substrates for membranes23,25). Supported functional lay-rs and pellets were sintered in air at temperatures up to 1300 ◦Csing either conventional sintering in furnaces or advanced

adiation-thermal sintering by e-beams (RTS)/microwave sinter-ng (MS). MS was carried out using a system based on gyrotronith frequency 24 GHz specially designed for heating of mate-

ials. Samples were heated by the focused radiation (power

(a(o

eramic Society 33 (2013) 2241–2250

.5–1.5 kW) with heating rate 50◦/min followed by dwellingt the final temperature for 30 min and then cooling to roomemperature. RTS was carried out on an ILU-6 accelerator2.4 MeV electron pulses, 8–20 Hz pulse frequency, heating rate0–40 ◦/min, time of treatment 10–240 min).29

The structural/microstructural features of composites andheir transport properties have been studied by using XRD,RTEM and SEM with EDX, electrical conductivity mea-

urements, oxygen isotope heteroexchange, conductivity/weightelaxation techniques using earlier described procedures.19–29

XRD patterns were obtained with an ARLX’TRA diffrac-ometer (Thermo, Switzerland) using Cu Kα monochromaticadiation (λ = 1.5418 A) in 2θ range 5–90◦.

Transmission electron microscopy (TEM) micrographs werebtained with a JEM-2010 instrument (lattice resolution 1.4 A,cceleration voltage 200 kV). Analysis of the local elementalomposition was carried out by using an energy-dispersive EDXpectrometer equipped with Si(Li) detector (energy resolution30 eV).

Details of the microstructure of sintered layers/pellets weretudied with the scanning electron microscope JEOL JSM-460LV equipped with an EDX-INCA/Energy-350 (Oxfordnstr.) spectrometer.

The oxygen mobility and surface reactivity of powdered sam-les were characterized by the oxygen isotope exchange usingoth static and flow reactors with MS control of the gas phasesotope composition.22,23

For dense pellets, conductivity was measured with a MO-10icro-Ohmmeter at 25 Hz frequency. The oxygen chemical dif-

usion coefficients and amount of easily removed oxygen werestimated by analysis of their weight relaxation (using a STA 409C “LUXX” NETZSCH machine) or conductivity relaxationfter step-wise change of O2 partial pressure.22

Performance of functional layers sintered by using differentechniques was tested using button-size solid oxide fuel cellsperating with wet H2–air feeds and reactors equipped withsymmetric supported membranes.22,23

. Results

.1. Phase composition and microstructure

For all perovskite–fluorite nanocomposites prepared vialtrasonic treatment in isopropanol including La0.8Sr0.2MnO3LSM)–ScCeSZ, LaCo0.5Fe0.5O3 (LFC)–Ce0.9Gd0.1O2 (GDC),aFe0.7Ni0.3O3 (LFN)–GDC, La0.8Sr0.2Fe1−xNixO3 (LSFNx)–DC, La0.3Bi0.7MnOx–BiErO3−δ (LBM–BE), La0.3Bi0.7nOx–Bi1.5Y0.3Sm0.2O3 (LBM–BYS), La0.8Sr0.2Fe0.8o0.2O3 (LSFC)–GDC, sintering at temperatures below000 ◦C was not accompanied by appearance of new phases.he same was true for LSM–ScCeSZ, LFN–GDC, LFC–GDCystems sintered to complete density at higher temperaturessing both traditional and advanced sintering techniques22,24,29

see Figs. 1 and 2 for LFN–GDC system). For LSFN–GDCnd Pr2NiO4–GDC nanocomposites sintered at higher1200–1300 ◦C) temperatures, admixture of NiO phase wasbserved.20–23,27 Hence, substitution of La by Sr or Pr

V. Sadykov et al. / Journal of the European Ceramic Society 33 (2013) 2241–2250 2243

20 30 40 50 60 70 80

Inte

nsity, a.u

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Fig. 1. X-ray diffraction pattern of samples of LFN–GDC system calcined at13

appLnz

prdpFffeStUeefloi

F7(p

Table 1Effect of sintering temperature on the lattice parameters of perovskite and fluoritephases in LSM–ScCeSZ composite.

Sintering T (◦C) Lattice parameter (Å)

Fluorite Perovskite

a a c

700 5.0927 5.499 13.365900 5.0965 5.513 13.3781

e(So(Gtblri

parrtocpiode

300 ◦C. (1) LFN, (2) 70% LFN + 30% GDC, (3) 50% LFN + 50% GDC, (4)0% LFN + 70% GDC, (5) GDC.

pparently decreases stability of a complex perovskite-likehase in composites with GDC. For LSFN0.2–ScCeSZ com-osite sintered at 1200 ◦C, such phases as NiO, La2Zr2O7,aSrFeO4 were observed as well,20 thus demonstrating pro-ounced chemical interaction between perovskite and dopedirconia as well.

Interaction between perovskite and fluorite phases in com-osites leads to redistribution of cations between their domainsevealed by TEM with EDX (Fig. 3) even when new phaseso not appear. This is reflected in the shift of diffractioneaks positions with respect to those of pure phases (see, i.e.,igs. 1 and 2). The highest degree of transfer of rare-earth cationsrom perovskite-like oxide into fluorite phase was observedor Pr2NiO4–GDC system (Fig. 3c), which can be apparentlyxplained by well-known high solubility of Pr cations in ceria.30

ome amount of transition metal cations is incorporated intohe surface layer of fluorite phase as well19–29 (cf. Fig. 3c).sually, for fluorite phase in nanocomposites, the lattice param-

ter increases with the sintering temperature (Table 1), which isxplained by progressive incorporation of big La cations into the

uorite lattice.22 At the same sintering temperature, reflectionsf fluorite phase are shifted to lower angles (lattice parameterncreases) with increasing perovskite content in composites thus

20 30 40 50 60

0

200

400

600

800

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P

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F

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F

F

P

P

P

d

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a

b

c

P

ig. 2. Effect of sintering on XRD patterns of LFN–GDC composite. (a) GDC00 ◦C, (b) LFN 700 ◦C, (c) LFN–GDC after conventional sintering at 1300 ◦C,d) LFN–GDC after radiation-thermal sintering at 1100 ◦C. P – perovskite-typehase, F – fluorite-type phase.

otst(fo

(pr

3

pp1icp

100 5.0990 5.526 13.381

videncing a larger extent of La transfer into the fluorite latticeFig. 1). The lattice parameters of LSM phase in composite withcCeSZ increase with sintering temperature due to incorporationf bigger Zr and Sc cations into the B sublattice of perovskite22,24

Table 1). For L(S)FC or L(S)FN perovskites in composites withDC, variation of the lattice parameters with sintering tempera-

ure is more complex due to concurrent processes of exchange ofig La and Sr cations with smaller Ce and Gd cations (decreasesattice parameter) and oxygen loss from the lattice leading toeduction of Co3+/Ni3+ cations into Co2+/Ni2+ state with biggeronic radii (increases lattice parameter).22

In all composites, domain sizes of perovskite and fluoritehases increase with sintering temperature remaining in nanor-nge even in completely dense materials (Fig. 4), so they can beeally termed as nanocomposites. For pure perovskite and fluo-ite phases sintered at the same temperature, X-ray sizes are 2–3imes larger than those in nanocomposites.22,27 The domain sizef GDC tends to decrease with the content of perovskite phases inomposite (Fig. 5) which can be assigned to the dilution effect byerovskite matrix.27 For perovskite phase, the domain size variesn a complex manner reaching a minimum at the equal contentf both phases (Fig. 5). The increase of the perovskite phaseomain size at a high content of GDC apparently could not bexplained by any simple dilution effect. It implies some increasef perovskite sinterability in composites with GDC excess dueo pronounced redistribution of cations between phases (videupra). Indeed, disordering of perovskite structure by genera-ion of cation vacancies due to transfer of big A site cationsLn, Sr) into GDC domains is expected to accelerate the sur-ace diffusion of transition metal cations thus favoring growthf perovskite domains at sintering.

The surface image of polished 1:1 nanocomposite pelletsFig. 6) demonstrates a good uniformity of spatial distribution oferovskite (dark regions) and fluorite (bright regions) domainsequired for percolation and high conductivity/oxygen mobility.

.2. Transport properties

Total (mainly, electronic) specific conductivity of nanocom-osites increases with the sintering temperature due toorosity annealing and percolation improving.19–22,27–29 For:1 perovskite–fluorite nanocomposites sintered at ∼1200 ◦C,

t reaches the values up to 10–40 S/cm, being higher for Sr-ontaining systems and nanocomposites with a higher content oferovskite phase (Figs. 7 and 8). Though this level is by an order

2244 V. Sadykov et al. / Journal of the European Ceramic Society 33 (2013) 2241–2250

F ) nanoo Ox; (b

oscppps

tbtwi

Fa(p

ivpiamippF

ig. 3. High resolution TEM images of LFC–GDC (a, b) and Pr2NiO4-GDC (cf domains by EDX: (a) 1 – La0.61Fe0.21Co0.17Ox, 2 – La0.08Ce0.74Gd0.08Co0.05

f magnitude lower than specific conductivity of pure perov-kites (Fig. 8), it is still sufficient for their application as SOFCathode materials.1–3 Radiation-thermal sintering at lower tem-eratures as compared with those for conventional sinteringrovides the same level of specific conductivity of nanocom-osite pellets (Fig. 8), apparently due to close density of pelletsintering by different techniques.

The oxygen mobility in nanocomposites characterized byhe dynamic extent of oxygen heteroexchange XS (the num-er of monolayers exchanged up to a given temperature in the

24,26

emperature-programmed mode in a static reactor ) increasesith the sintering temperature (Fig. 9) apparently correlat-

ng with porosity annealing and increasing perovskite–fluorite

ig. 4. Effect of sintering temperature on X-ray particle sizes of perovskite (P)nd fluorite phase (F) in nanocomposites La0.8Sr0.2Fe0.6Ni0.4O3–GDC1), La0.8Sr0.2Fe0.7Ni0.3O3–GDC (2), LSM–ScCeSZ (3) with 1:1erovskite–fluorite ratio.

Xdhc

FC

composites sintered at 700 ◦C (a), 1300 ◦C (b) and 1250 ◦C (c). Compositions) 1 – La0.1Ce0.8Gd0.08Ox, 2 – La0.6Fe0.2Co0.2Ox; (c) Ce0.48Pr0.47Gd0.02Ni0.01.

nterface. In general, XS values are large and rather close forarious nanocomposites, while the difference in XS betweenerovskites (i.e., LSM and LSFN) is very big: at 650 ◦C XSs ∼1–2 monolayers for LSM with a low bulk oxygen mobilitynd ∼200 monolayers for LSFN0.4, while for GDC it is ∼50onolayers.22 Hence, non-additive increase of oxygen mobil-

ty in nanocomposites as compared with perovskite and fluoritehases is apparent. For LSFN–GDC and LSFC–GDC nanocom-osites sintered at 1250–1300 ◦C, the values of XS given inig. 9 correspond to the fraction of exchanged bulk oxygenV ∼ 0.2. Microwave sintering of LFN–GDC composites intoense ceramics at lower (1100 ◦C) temperatures provides even

◦

igher oxygen mobility (XV ∼ 0.5 at 650 C, Fig. 10), whichan be explained by smaller domain sizes, and, hence, more

ig. 5. Effect of La0.8 Sr0.2Fe0.7Ni0.3O3 content (wt.%) in nanocomposite withe0.9Gd0.1O2−δ sintered at 1200 ◦C on the domain sizes of both phases.

V. Sadykov et al. / Journal of the European Ceramic Society 33 (2013) 2241–2250 2245

Fig. 6. The image of the surface of LFN–GDC (1:1) nanocomposite sintered at1200 ◦C in back-scattered electrons.

FcL

dt

r

F11

Fig. 9. Effect of sintering temperature on the nanocomposites oxygen mobil-ity characterized by the dynamic extent of exchange.21,22 (1) LSM–ScCeSZ(1:1), (2) LSCF–GDC (7:3), (3) LSFN0.3–GDC (1:1). PO2 = 4 Torr, heating ramp5◦/min.

Fig. 10. Temperature dependence of the dynamic extent of exchange X forLh

ie

ig. 7. Effect of perovskite content in nanocomposites with GDC on their spe-ific conductivity at 1000 K. (1) LSFN0.3–GDC sintered at 1200 ◦C and (2)FC–GDC sintered at 1250 ◦C.

eveloped perovskite–fluorite interface in nanocomposite sin-ered by microwave heating.

Specific rate of oxygen heteroexchange (characterizes theeactivity of surface sites in activation of O2 molecules)

ig. 8. Temperature dependence of specific conductivity of GDC (1), LFC–GDC:1 nanocomposite (2, 3) and LFC (4) sintered in the conventional furnace at300 ◦C (1, 3, 4) or by radiation-thermal sintering at 1100 ◦C (2).

teho

FoL(

V

FN–GDC composites sintered by microwave heating at 1100 ◦C. PO2 = 4 Torr,eating ramp 5◦/min. LFN:GDC = 7:3 (1), 3:7 (2) and 1:1 (3).

ncreases with sintering temperature as well (Fig. 11). It isxplained by progressive transfer of transition metal cations onhe surface of electrolyte domains creating new surface sites with

nhanced activity.22–24,27 Bi-containing nanocomposites have aigher specific rate of exchange which can be due to participationf Bi cations in activation of O2 molecules as well.

ig. 11. Effect of nanocomposite sintering temperature on the specific rate ofxygen heteroexchange at 1000 K. (1) La0.3Bi0.7MnOx + BiErO3−δ (1:1), (2)a0.3Bi0.7MnOx + Bi1.5Y0.3Sm0.2O3 (1:1), (3) La0.8Sr0.2Fe0.8Co0.2–GDC (7:3),

4) LSM–ScCeSZ (1:1), (3) LSFN0.3–GDC (1:1).

2246 V. Sadykov et al. / Journal of the European Ceramic Society 33 (2013) 2241–2250

Fig. 12. Typical variation of mole fractions of 18O (α) and C16O18O molecules(X ) at the reactor exit after switching the stream of He for the stream of 1%C1

toiapPdndbppamsapeGab

Fd

dabtcmmrfcphSi(

Psis

TC

O

L

PP

P

LL

L

118O2 in He at 600 ◦C over Pr2NiO4–GDC (1:1) nanocomposite sintered at200 ◦C. Pretreatment in O2 at 650 ◦C, points – experiment, lines – fitting.

Oxygen isotope exchange experiments in the flow reac-or (SSITKA mode) allowed to estimate the coefficients ofxygen self-diffusion in nanocomposites (DO) by fitting exper-mental transient curves (Fig. 12) following earlier describedpproaches.26 In agreement with published data,31 amongerovskite-like oxides the highest DO values were found forr2NiO4 in which oxygen mobility is determined by fastiffusion of oxygen interstitials (Table 2). For LSFN–GDCanocomposites, experimental data could be satisfactorilyescribed only in frames of model with a fast exchangeetween the surface and mobile oxygen species located at theerovskite–fluorite interface followed by a slower exchange witherovskite and fluorite domains.26 This model clearly revealstomic-scale factors providing non-additive increase of oxygenobility in nanocomposites as compared with separate perov-

kite and fluorite phases. For Pr2NiO4–GDC nanocomposite,nalysis of 18O2 isotope transients has not allowed to estimatearameters of oxygen diffusion along all these pathways, appar-

ntly due to rather close and high DO values for Pr2NiO4 andDC. The average DO values determined for this nanocomposite

re lower than those for Pr2NiO4 (Table 2). This can be explainedy a strong chemical interaction between Pr2NiO4 and GDC

DPfG

able 2haracteristics of oxygen mobility and reactivity for powdered oxides by SSITKA of

xide/T sintering (◦C)

a0.8Sr0.2Fe0.7Ni0.3O3+δ–GDC (1:1)/1200 (18O2)

r2NiO4/1100 (18O2)

r2NiO4–GDC/1200 (18O2)

r2NiO4–GDC/1200 (C18O2)

a0.3Bi0.7MnOx/800 (18O2)

a0.3Bi0.7MnOx + Bi1.5Y0.3Sm0.2O3/800 (18O2)

a0.3Bi0.7MnOx + Bi1.5Y0.3Sm0.2O3/800 (C18O2)

ig. 13. Temperature dependence of Dchem and kchem for dense pellet of Pr2NiO4

etermined by conductivity relaxation method.

omains in nanocomposite (vide supra) disrupting the cooper-tive process of oxygen migration in Pr2NiO4 which involvesoth interstitial and regular oxygen positions.32 Moreover, inhe case of a high bulk oxygen mobility, the surface reactionould be the rate-limiting stage of the oxygen exchange, thusaking difficult reliable estimation of DO.22 Indeed, experi-ents with C18O2 exchange characterized by a much higher

ate of the surface reaction allowed to estimate the oxygen dif-usion parameters in both GDC and Pr2NiO4 domains (Table 2)haracterized by close DO values. For LBM + BYS nanocom-osite, the average DO values estimated by 18O2 SSITKA areigher than in perovskite but lower than in BYS (Table 2). C18O2SITKA for this system allowed to distinguish a fast diffusion

n BYS domains and a much slower diffusion in LBM domainsTable 2).

Estimation of Dchem and kchem for Pr2NiO4 andr2NiO4–GDC nanocomposite (Figs. 13 and 14) revealedome decrease of both surface reactivity and oxygen mobilityn nanocomposite in agreement with SSITKA data (videupra). In the high-temperature region, some decrease ofchem is observed apparently due to oxygen loss from the

r2NiO4+δ lattice.33 However, in all temperature range, Dchemor Pr2NiO4–GDC (Fig. 14) is higher than for 70% LSFC–30%DC nanocomposite (Fig. 15) well-known for its high

18O2 and C18O2.

Texchange (K) D (cm2 s−1)

973 Perovskite >2 × 10−14

Fluorite >3 × 10−13

Interface >5 × 10−8

1000 >6 × 10−11

1000 >10−11

873 Perovskite ∼1 × 10−11

Fluorite ∼3 × 10−11

1000 2 × 10−13

1000 3 × 10−12

873 LBM < 10−13

BYS > 10−10

V. Sadykov et al. / Journal of the European Ceramic Society 33 (2013) 2241–2250 2247

F Pr2NiO4–GDC (1:1) nanocomposite determined by conductivity relaxation method.

optLtatpoddlnetpi

3o

a

F(LL

Fig. 16. Oxygen chemical diffusion coefficient Dchem and chemical exchangec1

s

ig. 14. Temperature dependence of kchem (a) and Dchem (b) for dense pellet of

xygen mobility.1 LBM–BYS nanocomposite without Sr in theerovskite lattice demonstrates the oxygen mobility close tohat in 70% LSFC–30% GDC nanocomposite, while Dchem forSFN–GDC is lower by an order of magnitude (Fig. 15). In

he high-temperature range, the bulk lattice oxygen mobilitynd surface reactivity of LFC–GDC nanocomposites is closeo those of 70% LSFC–30% GDC nanocomposite. Hence,erovskite–fluorite nanocomposites could provide a high latticexygen mobility and reactivity even when perovskites are notoped by Sr, which is important for preventing performanceegradation caused by blocking of surface sites by Sr carbonateaayer. The increase of both Dchem and kchem in LFN–GDCanocomposite with GDC content (Fig. 16), similar to thearlier observed trend for LSFN–GDC system,21 agrees withhe model explaining enhanced oxygen mobility in nanocom-osite by fast oxygen migration along the perovskite–fluoritenterface.23,26

.3. Performance of nanocomposite layers in SOFC andxygen separation membranes

Nanocomposite layers supported on SOFC YSZ/NiO + YSZnode substrates as well as on Ni–Al foam substrates of oxygen

ig. 15. Temperature dependence of Dchem (cm2 s−1, filled symbols) and kchem

cm s−1, empty symbols) for dense pellets of 70% LSFC–30% GDC (1), 30%FC–70% GDC (2), 50% LFC–50% GDC (3), 70% LFC–30% GDC (4),BM–BYS (5), LSFN0.3–GDC (1:1) (6).

aoacc(G

TMa

C

L(

L(

L((

onstant kchem vs. GDC content estimated by the weight relaxation method at050 ◦C for LFN–CGO composite sintered at 1200 ◦C.

eparation membranes were successfully sintered by microwavend/or radiation-thermal treatment. These techniques allowed tobtain required density of nanocomposites even at rather moder-te sintering temperatures (Fig. 17), thus preventing undesirablehemical interaction with substrates.25,28,29 For solid oxide fuelells with thin layers of YSZ electrolyte on anode substrates

Table 3), the maximum power density was increased whenDC interlayer was replaced by a thin dense layer of MIEC

able 3aximum power density of thin film solid oxide fuel cells operating on wet H2

s fuel and air as oxidant.

ell/[Reference] T (◦C) Maximum powerdensity (mW/cm2)

SFN0.2 (50 �m)/GDC (∼15 �m)/YSZ20 �m)/Ni-YSZ26

700 380750 510800 580

SFN0.2 (15 �m)/LSM–ScCeSZ10 �m)/YSZ (5 �m)/Ni-YSZ26,28

600 210700 480800 680

SM (50 �m)/LSM–ScCeSZ (10 �m)/YSZ5 �m)/Ni-YSZ/Ni–Al foam substrate1 mm)25,28

600 150700 450800 650

2248 V. Sadykov et al. / Journal of the European Ceramic Society 33 (2013) 2241–2250

Fig. 17. SEM image demonstrating typical structure of LFN–GDC (a) and La0.3Bi0.7MnOx–BiErO3−δ layers (b) on Ni–Al foam-based substrates after radiation-thermal sintering at 900 ◦C.

F ties (L l side

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spitwtwith electrolyte (YSZ), surface segregation of alkaline-earthoxides and surface sites poisoning by carbonates.19,37–40 Alter-native approach used in our research is based upon design of

ig. 18. Temperature dependence of CH4 conversion and products selectiviSFN–GDC/MnFe2O4–GDC layers and top porous Pt/PrSmCeZrO layer at fue

anocomposite. The cell performance was stable for middle-erm (up to 100 h) testing times.22,28,29

For design of supported asymmetric membranes, binaryi–Al foam substrates with graded porosity described

n details elsewhere28 were used. Membranes for trans-ormation of CH4 into syngas by oxygen separatedrom air were prepared by covering these substrates byacroporous–mesoporous–microporous layers of LSFN–GDC

anocomposite followed by a dense layer of MnFe2O4–GDCnd a porous layer of Pt/PrSmCeZrO catalyst.23 At temperaturesn the range of 900–1000 ◦C, high CH4 conversions and syngaselectivity were obtained (Fig. 18). The oxygen flux throughembranes achieves values up to 6–10 mL O2/cm2 min, which

s sufficient for the practical application.9–11

For supported asymmetric membrane comprised ofBM–BYS layers on binary Ni–Al foam substrate, the oxy-en flux from air into Ar stream was up to 5 mlO2/cm2minFig. 19),29 which is also promising for the practical application.

. Discussion

For successful application of MIEC oxide materials as

OFC cathodes and oxygen separation membranes, they shouldrovide a high oxygen mobility and a high reactivity of the sur-ace sites, reasonably high level of conductivity, the chemicalnd thermophysical compatibility with electrolyte and chemical

Fpl

a) and oxygen flux j (mL O2/cm2 min) (b) through membrane with MIEC. Feed 21%CH4 in Ar, flow rate 2 L/h, air flow rate 2 L/h.

tability in real operation conditions.1–9 While for single-phaseerovskite-like structures substitution of La by Sr, Ba, etc.ncreases the oxygen mobility and surface reactivity by genera-ion of oxygen vacancies (disordering of anion sublattice) alongith a high electronic conductivity,34–36 it leads also to degrada-

ion caused by strong chemical interaction of doped perovskites

ig. 19. Temperature dependence of oxygen flux under air/He gradient for sup-orted asymmetric membrane comprised of La0.3Bi0.7MnOx–Bi1.5Y0.3Sm0.2O3

ayers on binary Ni–Al foam substrate.

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V. Sadykov et al. / Journal of the Europ

anocomposite MIEC materials structured at nanoscale due toxistence of perovskite–fluorite interfaces. Specificity of thetructure of interfaces is determined both by redistribution ofations between coexisting phases and heteroepitaxy betweenerovskite–fluorite surface faces. While depletion of the surfaceayers of perovskite domains by La cations due to their pref-rential transfer into the surface layer of fluorites apparentlyenerates disordered oxygen layers with a plenty of vacancies,heir enrichment by transition metal cations with disorderedoordination spheres helps to decrease the average metal-oxygentrength as well as to enhance the electronic conductivity due tolustering of transition metal cations. In frames of known modelsf oxygen diffusion in nanostructured perovskites,22,35,36,41,42

his rearrangement decreases the barrier for the oxygen ionigration, thus explaining fast oxygen diffusion along these

nterfaces demonstrated in our research. Similarly, substitutionf a part of La by Bi with a weaker metal-oxygen bond andff-center positions in the lattice helps to further decrease thectivation barrier of migration thus increasing the oxygen mobil-ty along the perovskite–fluorite interface.43 On the other hand,eteroepitaxy of the surface planes of stacked perovskite anduorite nanodomains22 apparently helps to prevent appearancef large-angle grain boundaries which could serve as barriers forhe oxygen migration. This model apparently explains why it isossible to ensure a good level of oxygen mobility in nanocom-osites of GDC with undoped LFN (LFC) perovskites as well asn LSM–ScCeSZ nanocomposite despite very low bulk oxygenobility in LSM.22 As the result, the average parameters char-

cterizing the oxygen mobility in nanocomposites (XS, Dchem)re quite close for different systems studied here. Note alsohat Dchem is estimated by analysis of the weight/conductivityelaxation caused by removal of a small (∼1%) fraction of theverall amount of oxygen contained in nanocomposites,22 whichupports hypothesis on the location of this oxygen at perov-kite/fluorite interfaces.

This model seems not to work in the case of nanocompos-tes with Ruddlesden–Popper oxides (Pr2NiO4), in which highxygen mobility is provided by oxygen interstitials. Indeed, inhe case of Pr2NiO4–GDC nanocomposites only some decreasef oxygen mobility was observed as compared with Pr2NiO4hase. However, these nanocomposites ensure a higher sta-ility of the cathode microstructure than Pr2NiO4 known toecompose under air at intermediate temperatures into theixture of Pr4Ni3Oy + PrOx with poorly controlled textural

roperties.33

From the surface reactivity point of view, nanocompos-tes were reliably demonstrated to possess enhanced catalyticctivity in the oxygen heteroexchange as a basic stage of O2olecules activation on cathode materials. This is again caused

y redistribution of cations between the perovskite and fluoriteomains. The most active sites of perovskite surface involvedn dissociation of O2 molecules are coordinatively unsaturatedransition metal cations.22 Along with disordering of the surface

f perovskite domains and its enrichment by transition metalations considered above, another important factor is migra-ion of some transition metal cations on the surface of fluoriteomains creating new surface sites.22

eramic Society 33 (2013) 2241–2250 2249

From the technological point of view, another importantesult of our research is that applied procedures allowed simplend inexpensive routes for synthesis of mixed ionic–electroniconducting nanocomposites with uniform spatial distribu-ion of domains of constituting phases. Sintering of greenanocomposites into dense ceramics proceeds at rather mod-rate (1100–1200 ◦C) temperatures, especially when modernintering techniques (microwave sintering, radiation-thermalintering) are applied. As the result, domain sizes remain in theanorange even in dense ceramics, which provides developederovskite–fluorite interfaces.

Testing of thin supported MIEC nanocomposite layers asOFC cathodes and oxygen permselective layers in asymmet-ic supported membranes demonstrated their high and stableerformance promising for the practical application.

. Conclusions

A procedure of mixed ionic–electronic conductingerovskite–fluorite nanocomposite synthesis based uponltrasonic treatment of the mixture of nanocrystalline powdersn organic solvent provides fine and uniform intermixing ofomponents required for a good percolation. New advancedintering techniques based upon the radiation-thermal oricrowave heating ensure a high density of nanocomposites

lready after sintering at 1000–1100 ◦C. As a result, highxygen mobility and surface reactivity were achieved due to aositive role played by developed perovskite–fluorite interfacesith specific structure and composition. A high and stableerformance of these nanocomposites as functionally gradedathode layer in thin film SOFC as well as oxygen-separationayers in asymmetric supported membranes promising for theractical application was demonstrated.

cknowledgements

Support by OCMOL FP7 Project, Integration Projects SBAS – Belarus NAS, Russian Federation Government Grant N1.G34.31.0033 and Federal Program “Scientific and Educa-ional Cadres of Russia” is gratefully acknowledged.

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