3D core–shell architecture from infiltration and beneficial reactive sintering as highly efficient...

Journal ofMaterials Chemistry A

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aState Key Laboratory of Materials-Orien

Chemistry & Chemical Engineering, Nanji

Mofan Road, Nanjing 210009, P.R. China.

8317 2242; Tel: +86 25 8317 2256bDepartment of Mechanical Engineering, T

Technology, Clear Water Bay, Kowloon, HoncDepartment of Chemical Engineering, Curti

Cite this: DOI: 10.1039/c3ta13253f

Received 16th August 2013Accepted 19th October 2013

DOI: 10.1039/c3ta13253f

www.rsc.org/MaterialsA

This journal is © The Royal Society of

3D core–shell architecture from infiltration andbeneficial reactive sintering as highly efficient andthermally stable oxygen reduction electrode

Dengjie Chen,a Guangming Yang,a Francesco Ciucci,b Moses O. Tadec

and Zongping Shao*ac

Solid oxide fuel cells (SOFCs) as alternatives for energy conversion have the capacity to overcome low

energy conversion efficiency, highly detrimental emissions from traditional fuel utilization and the limited

reserves of fossil fuels crisis. Herein, a 3D core–shell architecture has been fabricated from solution

infiltration in combination with high-temperature reactive sintering and evaluated as the oxygen

reduction electrode for SOFCs. The resultant electrode is composed of a stable porous Sm0.2Ce0.8O1.9

scaffold as the core for bulk oxygen ion diffusion, and a connective Sm,Ce-doped SrCoO3�d perovskite

film as the shell for efficient oxygen reduction reaction and partial current collection. The significant

enhancement in conductivity, chemical and thermal compatibility with such core–shell structured

electrodes can deliver promising and stable power outputs. An anode-supported solid oxide fuel cell

with such a core–shell structured cathode exhibits a peak power density of 1746 mW cm�2 at 750 �C,which is comparable to the most promising cathodes ever developed. In addition, both a symmetrical

cell and a fuel cell demonstrate favourable short-term stability during 200 h operation at 700 �C. Thecombined strategy involving infiltration and high-temperature reactive sintering (accompanied by ion

diffusion) appears to be a promising approach to fabricate cathodes with high electrochemical

performance and stability.

1. Introduction

Improvement of energy conversion efficiency and emissionreduction during the utilization of fossil fuels have becomeimportant issues for realizing a sustainable future.1 Among thevarious advanced energy conversion technologies, electro-chemical energy conversions have received particular attentionrecently for their superior efficiency and low emissions.2–5 As ahigh-temperature electrochemical device, solid oxide fuel cells(SOFCs) are known for their additional advantages of fuel ex-ibility and high value of exhaust heat.6–10 Current developmentsin SOFCs focus on reducing the operation temperature to theintermediate range (500–800 �C), which can lead to prolongedlifetime and reduced materials and operation costs, thusmaking this fascinating technology more competitive withcurrently matured power generation systems.11,12 However, thesluggish oxygen reduction reaction (ORR) kinetics has become

ted Chemical Engineering, College of

ng University of Technology, No. 5 Xin

E-mail: [email protected]; Fax: +86 25

he Hong Kong University of Science and

g Kong, China

n University, Australia

Chemistry 2014

the main obstacle for reducing the operation temperature ofSOFCs.13,14

Up to now, several ways have been exploited to increase thereduced temperature performance of the oxygen reductionelectrode of SOFCs, including the adoption of mixed oxygen-ionand electronic conducting oxides as electrode materials toextend active reaction sites, electrode surface modication withnanocatalysts to increase oxygen surface exchange kinetics, andelectrode microstructure optimization through building porous3D architectures to increase the effective electrode surface areaand to minimize concentration polarization.15–22 Althoughmany highly active electrodes with outstanding performance forORR at intermediate temperature have appeared from time totime in literature, materials that are suitable for practical use inSOFCs are still lacking up to now. This is due to the fact that apractical electrode of SOFCs should meet several strictrequirements while most of the available electrodes cannotmeet all of them simultaneously. In addition to the high electro-catalytic activity, it should have comparable thermal expansionto the electrolyte ensuring high thermo-mechanical compati-bility, high electrical conductivity ensuring efficient currentcollection, negligible chemical reactions with other cellcomponents ensuring sufficient chemical compatibility, andhigh structural and micro-structural stability ensuring longoperation stability.

J. Mater. Chem. A

http://dx.doi.org/10.1039/c3ta13253f

http://pubs.rsc.org/en/journals/journal/TA

Journal of Materials Chemistry A Paper

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Recently, SrCoO3�d (SC) based perovskites have attractedmuchattention owing to either their high oxygen vacancy concentrationor electronic conductivity e.g., Ba0.5Sr0.5Co0.8Fe0.2O3�d (BSCF),SmxSr1�xCoO3�d and Sr1�xCexCoO3�d.15,16,23–28 However, SC itselfis unstable and is a poor electrocatalyst for ORR due to theimmobilized oxygen ions in some phase structures.25,29,30 Dopingin A and/or B sites is usually adopted to stabilize the phasestructure and increase the ionic or/and electronic conductivity ofSC.25,27–29,31 Among various elements used for doping, doping Smin A sites and/or Ce in B sites successfully limits the phase tran-sition and the doped SC materials exhibit relatively high electro-chemical performance as electrodes.25,27,28 Nevertheless, theirpractical application has so far been hindered by several criticaltechnical barriers, such as the mismatch in thermal expansionwith the electrolyte, high reactivity with the electrolyte leading tothe formation of an insulating interfacial layer. Several strategieshave been tried to moderate the above-mentioned problems.Among them, the deposition of active nano-structured particlesonto a porous scaffold by solution inltration has been demon-strated to be a promising approach.21,22,32 Although it is somewhatsuccessful in enhancing initial kinetics of ORR, the long-termoperational stability is questionable due to the high instability ofnanoparticles at elevated temperatures.33,34 In addition, isolatedparticles could block the electron transfer in the electrode, espe-cially with the adoption of porous electrolyte scaffolds, thusintroducing current collection issues.

Thin lms are thermally much more stable than nano-particles. Recently, well-dened thin lms have been consid-ered as model electrodes to explore the fundamental ORRmechanism.35–37 However, they are not suitable as practicalelectrodes due to their low active surface area. More recently,some successes were achieved by coating a thin layer on aporous cathode scaffold to form 3D lm electrodes.38–41 Forexample, under cathodic polarization, the electrochemicalperformance and operational stability of theLa0.6Sr0.4Co0.2Fe0.8O3�d cathode was enhanced aer coatingwith a dense La0.85Sr0.15MnO3 thin-lm layer.38 Recently, Zhouand Shao et al. have also demonstrated that a dense La2NiO4+d

thin lm modied with cone-shape particles (thus improvingthe active surface area) effectively improved the activity andoperation stability of a BSCF electrode under CO2-containing airfor ORR.41 However, both La0.85Sr0.15MnO3 and La2NiO4 havevery low ionic conductivity, thus the thickness of the thin lmshould be carefully controlled to minimize ion-transportationpolarization resistance at intermediate temperatures. In addi-tion, special apparatus is required for fabricating such thinlms, making it difficult for large scale preparation.

In our previous paper, we demonstrated that the interfacialphase reaction between SC and Sm0.2Ce0.8O1.9 (SDC) could beutilized to develop a high-performance cathode.42 Here wereport an innovative electrode consisting of a highly conductiveand active shell (Sm,Ce-doped SrCoO3�d) coating on a porousstable core (SDC) for ORR and its facile fabrication. The shellwas fabricated by a simple solution inltration in combinationwith a benecial interfacial reaction at high temperature. Theresulted electrode exhibited good electrical conductivity,comparable thermal expansion coefficient (TEC) to electrolyte,

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high ORR activity, and superior durability at intermediatetemperatures. The ndings may potentially provide a newstrategy to design highly active and stable electro-catalysts fromconventional processes for various applications in solid-oxidedevices.

2. ExperimentalFabrication of anode substrates

The cell used in this study is in an anode-supported thin lmelectrolyte conguration. NiO + YSZ (60 : 40, in weight ratio)anode substrate was prepared by a tape-casting process, whileYSZ electrolyte layer was deposited onto the anode substrateby air-driven spraying deposition. The slurry for tape castingwas prepared by two milling processes and the detailedprocedures to fabricate disk-shape anode substrates can befound in our previous work.43 The disk-shaped anodesubstrates were red at 1100 �C for 2 h in air to removeorganic solvents and to create sufficient mechanical strength.The YSZ electrolyte slurry with solid content of �5 wt% forwet spraying was prepared by ball milling an appropriateamount of YSZ powders in organic solvents at the rotationspeed of 400 rpm for 0.5 h in a planetary ball mill (Fritsch,Pulverisette 6). Aer spray-depositing YSZ slurry onto theanode substrates, the obtained NiO–YSZ anode-supportedsubstrate and YSZ electrolyte bilayer cells were red at 1400 �Cfor 5 h in air with a ramp rate of 3 �C min�1 to allow densi-cation of the electrolyte layer.

Preparation of porous scaffold

Porous SDC scaffold was fabricated by spraying deposition of aSDC layer over the electrolyte and then sintering at elevatedtemperature. SDC powder was synthesized by amodied sol–gelmethod,44 then 90 wt% SDC and 10 wt% starch were mixed withan appropriate amount of glycerol, ethylene glycol and iso-propyl alcohol to form a slurry under ball milling at a rotationspeed of 400 rpm for 0.5 h. The obtained slurry was spray-deposited onto the surface of the electrolyte, followed by sin-tering at 1300 �C for 5 h in air to form porous scaffolds with aporosity of �50%.

Fabrication of core–shell structured electrodes

Core–shell structured electrodes were realized by a facile tech-nique that combined the solution inltration and high-temperature reactive sintering. To fabricate connective thin lmcoating on the porous SDC scaffold, appropriate nitrate salts ofSr and Co (1 : 1, in molar ratio) were rst dissolved in mixedsolvents of water and ethanol (3 : 1, in volume ratio) to obtainabout 0.8 mol L�1 SC precursor solution. The ethanol was usedas surfactant to improve the wettability of the solution on theSDC scaffold because the surface tension of ethanol is muchlower than that of water. Next, glycine, at a molar ratio of glycineto total metal cations of 2 : 1, was added into the solution forcomplexing the metal cations. The inltration process wascarried out on a hot plate (60 �C) to accelerate the gas expulsionand enhance the liquidity of solution in the porous scaffold. SC

This journal is © The Royal Society of Chemistry 2014


Fig. 1 Cross-sectional SEM images of (a) as-prepared SDC scaffold,and the SSC-infiltrated SDC cathodes after sintering at (b) 800, (c) 900,(d) 1000 and (e) 1100 �C for 2 h.

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solution was inltrated into the porous scaffold using amicroliter syringe. The load of nominal SC was controlled at�20 wt% according to the equation SC% ¼mSC/(mSC + mSDC)� 100%, in which mSC and mSDC were based on the amountsweighed with an electronic balance. Aer the inltration of theappropriate amount of solution into the SDC porous scaffold,the electrodes were red at different temperatures (800, 900,1000 or 1100 �C) for 2 h.

Basic characterizations

The phase structures, microscopic features and electronicconductivities of the SC solution inltrated SDC electrodessintered at different temperatures were characterized by X-raydiffraction (XRD, D8 Advance), scanning electron microscopy(SEM, QUANTA 200 and JEOL S4800), transmission electronmicroscopy (TEM, JEOL JEM-2100) and four-terminal DCtechnique using a digital source meter (Keithley 2420). Theinterfacial ion diffusion was characterized by assessing solu-tion-deposited SC thin lms on dense SDC pellets with anSEM-energy dispersive X-ray spectrometer (EDS, JEOL S4800).The thermal expansion behaviour and thermal stability of thecore–shell structured electrode were studied by using a dila-tometer (NETZSCH DIL 402 C), thermo-gravimetry (TG) anddifferential scanning calorimetry (DSC, NETZSCH, STA 449 F3),respectively.

Electrochemical measurement

Anode-supported fuel cells and symmetric cells were fabricatedfor electrochemical performance and stability evaluation. Elec-trochemical impedance spectroscopy (EIS) of the symmetriccells with core–shell structured electrodes under open circuitconditions was obtained by using a potentiostat (Solartron1287) and a frequency response analyzer (Solartron 1260A) inthe frequency range from 10�1 Hz to 105 Hz with 10 mV as thesignal amplitude. Current–Voltage (I–V) and Current–Power(I–P) polarization curves of anode-supported fuel cells withcore–shell structured electrodes were collected by using a digitalsource meter (Keithley 2420) with a four-probe conguration.The cells were measured at 500–750 �C with humidiedhydrogen (3 vol% H2O) as the fuel at a ow rate of 80 ml min�1

and ambient air served as the oxidant gas. To avoid potentialdegradation of the Ag paste current collector, a mesh-like Agwas applied as the current collector during the stabilityevaluation.45

3. Results and discussion

Shown in Fig. 1a is an SEM image of the as-prepared SDCscaffold. It has a highly porous structure, constructed from wellfused micrometer particles with smooth grain surface. Theporosity was around 50%. The impregnation method has beenwidely applied for the preparation of perovskite oxide inltratedelectrodes in literature. We also tried the preparation ofconventional Sm0.5Sr0.5CoO3 (SSC) inltrated SDC scaffold aselectrode of SOFCs using a mixed solution of Sm(NO3)3,Sr(NO3)2 and Co(NO3)3 at appropriate cation ratios. Fig. 1b–e


show the SEM images of the electrodes aer ring at 800, 900,1000 and 1100 �C, respectively. For the electrode calcined at 800�C, we can clearly observe that clusters of nanoparticles with thesize of �100 nm modied the surface of the SDC scaffold.Previously we have demonstrated from the XRD pattern that theprecursor of SSC was successfully transformed to perovskiteoxide with an orthorhombic structure even aer calcination at800 �C.46 It suggests the SSC precursor did not react with theSDC scaffold. Aer calcination at 1000 �C we can still observenanoparticles over the scaffold surface, although the particlesize was apparently increased to �150 nm. Aer calcination at1100 �C, the nanoparticles disappeared, implying the formationof a thin lm over the SDC scaffold. However, the electrode wasseriously sintered, as demonstrated by the sharp decrease inporosity of the electrode (Fig. 1e).

The SEM images of SC solution-impregnated SDC elec-trodes aer ring at different temperatures between 800 and1100 �C are shown in Fig. 2. For the SC solution-impregnatedSDC scaffold aer ring at 800 �C, similar to the SSC solutionimpregnated SDC, we can clearly observe many nanoparticlesover the surface of the SDC scaffold. However, such nano-particles were seriously fused once the ring temperature wasincreased to 900 �C and the number of nanoparticles wasgreatly reduced. No nanoparticles were observed over thesurface of the electrode aer ring at 1000 �C or higher; inaddition, the electrode surface became smooth again. Itimplies the nanoparticles were likely sintered and formed intoa thin lm over the SDC scaffold. In contrast, nanoparticlescould still be observed aer ring at 1000 �C with the

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Fig. 2 Cross-sectional SEM images of the SC-infiltrated SDC cathodesafter sintering at 800, 900, 1000 and 1100 �C for 2 h.

Fig. 3 TEM results under different magnifications of the core–shellstructured electrode.


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Sm0.5Sr0.5CoO3�d (SSC) solution, suggesting that SC solution ismore reactive than SSC to form a thin lm on the SDCscaffold.

To support the potential formation of a thin lm over theSDC scaffold surface in the electrodes aer ring at 1000 �C orhigher, some particles from the 1000 �C red electrode were cutand subjected to TEM observation. As shown in Fig. 3, a corepart (darker area) and a shell part (brighter area) with thicknessless than 100 nm can be observed. The lattice spacing of theoutermost part was �0.270 nm, which corresponds well to the(110) diffraction planes of the cubic perovskite structure or (110)for the hexagonal perovskite (BaNiO3-type) structure of SC (PDF#48-0875). To further clarify the chemical composition of thecore–shell particles, EDS was performed and the correspondingdata are shown in Table 1. According to the atomic distributionof core–shell particles from 3 different samples, the atomiccontents of Sm and Ce are dominant in the darker area, whilecontents of Sr and Co elements are higher than those of Sm andCe elements in the brighter area. Therefore the darker area canbe regarded as the SDC-rich region, while the brighter area(outer layer) can be recognized as the SC-rich region. Incombination with the SEM images shown in Fig. 2, it is likely

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that 3D core–shell structured electrode was formed aer thehigh temperature ring ($1000 �C).

To obtain information about the phase composition of thevarious inltrated electrodes aer ring at different tempera-tures, XRD characterization was performed and the results areshown in Fig. 4. For a clear comparison, the XRD patterns of acrystalline SCphasewith ahexagonal BaNiO3-type structure, SSCwith an orthorhombic perovskite structure and SDCwith a cubicuorite structure are also presented. Our previous investigationhas demonstrated that SC reacted with SDC at elevated temper-ature inSC+SDCcomposite.42For the electrodered at 800 �C, inaddition to the characteristic peaks of the SDC phase, a vacancy-ordered hexagonal perovskite phase prevailed, which can beassigned to the SC phase. Thus, the nanoparticles observed onthe SDC scaffold surface were likely SC, formed from the SCprecursor. It implies the phase reaction between SC precursorand SDC crystalline phase was negligible at 800 �C. It alsosuggests, for later fuel cell application, no obvious interfacialphase reaction between cell components should happen if theoperation temperature is 800 �C or lower. Once the ringtemperature was elevated to 900 �C, a new peak at 2q of 40.89�

appeared, which matched the diffraction peak at the (111) planeof Sm/Ce-doped SrCoOx cubic perovskite phase,25,27 while thepeak intensity of BaNiO3-type phase at 2q ¼ 44.24� reduced. Forthe electrodes red at 1000 �C or higher, all peaks related to theBaNiO3-type phase structure disappeared, leading to theformation of amixture of a cubic perovskite phase and a uorite-type SDC. It suggests the minimum temperature for the fullconversion of BaNiO3-type SC phase to cubic phase is around1000 �C. Such phase transformation strongly suggests that asmall fraction of Ce and/or Sm ions likely diffused from the SDCscaffold and dissolved into the SC lattice, since doping either SmorCe ions into SChas been identied tobe effective in stabilizingthe oxide in either cubic or tetragonal perovskite structure.25,29

Meanwhile, Sr and/or Co ionsmay also back diffuse and dissolveinto the SDC during high-temperature sintering. Interestingly, itwas proposed that incorporation of either Sr or Co ions into rareearth-doped ceria could improve ionic conductivity.47,48



Table 1 Atomic distribution (%) of core–shell particles determined by EDS. Only Sm, Ce, Sr and Co elements are included for percentagecalculation

Region Sm (%) Ce (%) Sr (%) Co (%)

Darker core Particle I: 17.0 Particle I: 73.7 Particle I: 5.6 Particle I: 3.7Particle II: 19.1 Particle II: 79.1 Particle II: 0.5 Particle II: 1.3Particle III: 17.8 Particle III: 74.2 Particle III: 4.6 Particle III: 3.4

Brighter shell Particle I: 10.4 Particle I: 7.8 Particle I: 41.8 Particle I: 40.0Particle II: 8.8 Particle II: 4.8 Particle II: 47.7 Particle II: 38.7Particle III: 9.9 Particle III: 5.1 Particle III: 44.5 Particle III: 40.5

Fig. 4 XRD patterns of SC, SSC, SDC and SC-infiltrated SDC aftercalcination at 800, 900, 1000 and 1100 �C for 2 h, as well as thecomposite sample fired at 1000 �C and tested at 700 �C for 200 h.

Fig. 5 TG/DSC behaviors of the SC-infiltrated SDC sample fired at1100 �C.


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It is well known that SC has multiple phase structures,depending on the temperature, surrounding atmosphere andthermal history,49 and some obvious endothermic orexothermic peaks were reported in the DSC curves of SC due tothe phase transition during the heating and cooling process.The inltrated electrode red at 1100 �C was also subjected toTG/DSC characterization to gain information about potentialthermally induced phase transition. The sample red at 1100 �Cwas selected in this experiment since the phase transitiontemperature may be higher than 1000 �C. As shown in Fig. 5, noendothermic or exothermic peaks were observed in the DSCcurve within the investigated temperature range, suggesting thelack of rst order phase transitions in the sample. It furthersupports that Sm and/or Ce was likely diffused from the SDCscaffold and doped into SC, thus stabilizing the cubic latticestructure of the oxide.

To obtain direct evidence of thermally induced ion diffu-sion, a dense SDC pellet (to simulate the SDC particle in scaf-fold) deposited with a SC layer was prepared and co-red at1000 �C for 2 h, then subjected to EDS line-scanning. Accordingto Fig. 6, diffusion of Sm and Ce ions from the dense SDC layerinto the SC layer with the diffusion length around 1000 nm wasclearly demonstrated. The back diffusion of SC to the denseSDC layer could also be observed. It thus strongly supports thatthe Sm and Ce in SDC scaffold likely diffused into the SC layerand facilitated the phase transition of BaNiO3-type structure tocubic structure in the SC solution impregnated electrodes aer


calcination at 1000 �C or higher. Since the actual thickness ofthe SC lm in real electrodes was much smaller, the compo-sition of entire SC lms in the core–shell structured electrodescan be considered homogeneous. The successful diffusion ofSm and Ce from SDC with the formation of a thin and homo-geneous Sm,Ce-incorporated SC layer (shell) over the SDCscaffold was further supported from the EDS surface elementalmapping of an inltrated electrode red at 1000 �C for 2 h asshown in Fig. 7. Homogeneous distribution of Sr, Co, Sm andCe elements was observed, which was indicative of a connectivelm coating on the SDC surface on the one hand, andsuccessful diffusion of Sm and Ce from the scaffold to thesurface layer on the other hand. The simultaneous formation ofthin lm and cubic lattice structure for the electrode red at1000 �C suggests the critical role of reactive sintering for theformation of the thin lm.

As an electrode of SOFCs, high electric conductivity isrequired to ensure efficient current collection, while SC wasreported to have poor electronic conductivity due to structuralordering.29 The electronic conductivities of the various inl-trated electrodes red at 800, 900, 1000 and 1100 �C for 2 h, SCand SDC + SC composite were then measured by the 4-probe DCmethod. The SDC + SC composite was made from the corre-sponding powders which were synthesized by a combinedEDTA–citrate complexing process.44 The composite was thensintered at 1100 �C for 2 h and a rectangular bar with thedimensions of �2 mm � �5 mm � �10 mm and porosity of

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Fig. 6 SEM-EDS line scan concentration profiles of the SC thin filmcoated on the SDC dense layer after sintering at 1000 �C for 2 h.

Fig. 7 Cross-sectional SEM-EDS elemental mapping of the SC-infil-trated SDC cathode after sintering at 1000 �C for 2 h. Fig. 8 Electronic conductivities of (a) SC, (b) SC-infiltrated SDC

cathodes after sintering at 800, 900, 1000 and 1100 �C for 2 h and (c)SDC + SC composite after sintering at 1100 �C for 2 h.


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�50% was nally obtained for comparison. To make a directcomparison, the porosity of all samples was controlled at�50%through adjusting the amount of pore formers. The SC sample(Fig. 8a) showed modest conductivity of 4–10 S cm�1 withsemiconducting behaviour, i.e., an increase in conductivity withthe increase of temperature, over the investigated temperaturerange of 300–900 �C. As shown in Fig. 8b, for the inltratedelectrode red at 800 �C, the maximum conductivity was stilllower than 1 S cm�1, which can be explained by the poorintrinsic conductivity of SC and its isolated nanoparticlemorphological structure. For the samples red at 900 �C or

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above, the electronic conductivities are dramatically increasedby more than ten-fold; this is in good agreement with theformation of thin lms and the incorporation of Sm and Ce intoSC with the formation of cubic perovskite oxide. The formationof cubic perovskite led to the increase in electronic conductivitysince a cubic structure can maximize the overlapping betweenthe 2p orbital of oxygen with the 3d orbital of cobalt in perov-skite, which is benecial for the small polaron hopping. Theincrease in electronic conductivities with increasing ring



Fig. 10 ASRs of the SC-infiltrated SDC cathodes calculated from EIS inNyquist polts under open circuit condition between 550 and 750 �C.


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temperatures can be explained by a greater amount of Sm andCe ions diffused from the SDC scaffold into the SC shell sincethe cation diffusion was thermally activated. A decrease inconductivity with increasing test temperature, i.e. metallic-likeconducting behavior, was observed for the inltrated electrodesred at 900 �C or higher, which was very different from SC, butsimilar to Sm/Ce-doped SC as reported in literature.25,29 It alsogives indirect evidence for the formation of Sm/Ce-incorporatedSC in the inltrated electrodes. A much lower electronicconductivity of the SC + SDC composite (Fig. 8c) indicates thatthe connective thin lm is more effective in improving elec-tronic conductivity, mainly due to geometric effects.

Thermomechanical stability is another important concernfor SOFC electrode development. Large differences in TECs ofcell components may introduce large internal stress duringoperation, which could cause rapid degradation in cell perfor-mance due to delimitation. A signicant advantage of an inl-trated electrode is that the TECmismatch between the electrodeand electrolyte will be small if the scaffold is made frommaterials that have a similar TEC to the electrolyte. For theinltrated electrode, the thermal expansion behaviour is moresimilar to the scaffold, while the TEC of a composite electrode isa weighted average of its components. Shown in Fig. 9 are thethermal expansion curves of a porous SDC scaffold and a SCinltrated SDC electrode red at 1000 �C. The calculated TEC ofporous SDC is 12.9 � 10�6 K�1 at the temperature range of 200–800 �C, which matches well the reported value of SDC in liter-ature.50 For the inltrated electrode, a TEC of 15.6 � 10�6 K�1

was calculated from the thermal expansion curve. The specialcore–shell structured conguration effectively reduced the TECof the electrode, making it more compatible with the electrolyte.

Electrochemical activity for ORR of the various inltratedelectrodes was rst evaluated by EIS in a symmetric cellconguration, and the corresponding area specic resistances(ASRs) are shown in Fig. 10. The electrode red at 800 �C dis-played the largest polarization resistance, which can beexplained by the negligible interfacial reaction between the SCelectrode and the SDC electrolyte and the isolated SC particles

Fig. 9 Thermal expansion behaviour of the porous SDC and SC-infiltrated SDC bars sintered at 1000 �C.


to result in the low current collection efficiency. As expected, theelectrode performance increased obviously when the ring wasperformed at 1000 �C, to reach ASRs of 0.021, 0.04, 0.076, 0.17and 0.39 U cm2 respectively at 750, 700, 650, 600 and 550 �C.The observed ASRs are very encouraging, and are comparable tothe reported lowest polarization resistances of single-phaselandmark BSCF electrode and better than the SC + SDCcomposite electrodes prepared by co-ring of pre-mixedpowders or other conventional solution-inltrated compositeelectrodes.16,21,22,32,33,42,46,51,52 The different concentrations of theSm and Ce ions in the SC shell and the low surface area for ORRare two possible reasons for the slightly larger ASRs of theelectrode red at 1100 �C.

The electrochemical performance of the electrode red at1000 �C was further characterized in an anode-supported thin-lm YSZ (�15 mm) electrolyte single cell with humidiedhydrogen (3% H2O) as the fuel and ambient air as the oxidantgas. Shown in Fig. 11 are the I–V and corresponding I–P curvesof the cell. The open circuit voltages (OCVs) are close to thetheoretical values, reaching �1.05 to �1.09 V from 550 to 750

Fig. 11 Current–voltage and corresponding current–power curves ofthe cell consisting of the Ni–YSZ anode, YSZ electrolyte and SC-infiltrated SDC cathode sintered at 1000 �C with operating tempera-tures from 550 to 750 �C.

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�C. Peak power densities of 317, 573, 932, 1357 and 1746 mWcm�2 were achieved at 550, 600, 650, 700 and 750 �C, respec-tively. The cell with the novel core–shell structured electrodereported here had much better electrochemical performancethan similar SOFCs with single phase mixed conducting elec-trode or normal composite electrode, which already showedvery promising performance.25,53 The core–shell structuredelectrode is thus highly attractive and suitable for SOFCs atintermediate temperatures. The electrochemical performanceof the fuel cell may be further improved by adopting electrolyteswith high ionic conductivity at low temperature and optimizingthe microstructure in the anode.54,55

In addition to electrochemical activity, operational stabilityis also a big concern in developing new cathode mate-rials.33,34,56,57 Fig. 12a displays the impedance spectra from asymmetric cell with the core–shell structured electrode (red at1000 �C) aer operation for 4, 108 and 208 h at 700 �C. The ASRsremained constant even aer testing for 208 h, suggesting thatthis core–shell structured electrode is thermally stable. Thestability under full cell conditions is also critically important forpractical applications since the cathode experiences a loweroxygen partial pressure under fuel cell operating conditions dueto cathodic polarization.58 The short-term stability of thefabricated electrode was further investigated in real fuel cell

Fig. 12 (a) Impedance spectra of a symmetric cell under OCVconditions and (b) cell voltages of a fuel cell at a constant currentdensity with the SC-infiltrated SDC cathode sintered at 1000 �C afteroperating for �200 h.

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operating conditions. The time dependence of the cell voltage isrelatively stable at a constant current density of 650 mA cm�2 at700 �C (Fig. 12b), demonstrating that the fabricated electrode isstable for ORR during �200 h operation in fuel cell operatingconditions. Aer the performance test with the symmetric cell,the phase composition of the electrode was examined by XRD(Fig. 4). As expected, no extra diffraction peaks appeared, sug-gesting this in situ newly formed perovskite in the electrode wasstable during the intermediate-temperature operation.

Based on the above results, a mechanism for the formationof the 3D core–shell structured electrode and an explanation forthe improved electrode performance were proposed and areschematically shown in Fig. 13. Aer the inltration of SCsolution into the porous SDC scaffold, SC nanoparticles wereformed over the SDC scaffold surface at relatively low temper-ature (800 �C; Fig. 13a), under such conditions no reactivesintering was observed, similar to the inltrated electrodesreported in literature. With the further increase of ringtemperature to 1000 �C or higher, the diffusion of Sm and Cecations from the SDC scaffold occurred which were thenincorporated into the SC lattice to stabilize the oxygen vacancydisordered cubic perovskite phase on the one hand and inducedreactive sintering to convert the nanoparticles to a thin lmcoating over the SDC scaffold on the other hand (Fig. 13b). Forthe normal inltration–sintering process, due to the lack ofreactive sintering, only isolated nanoparticles were formed overthe scaffold, which had relatively poor electrical conductivity(Fig. 13c). Although similar interfacial reaction was alsoobserved in SC + SDC composite electrode, theSmxSr1�xCo1�yCeyO3�d phase was formed only at the SC–SDCgrain boundary (Fig. 13d), thus a large amount of SC wasrequired to allow such a phase to form a continuous electronicpathway, resulting in the large TEC of the electrode. Theformation of the cubic perovskite shell by reactive sintering inthis study resulted in increased electronic and oxygen-ion

Fig. 13 Schematic diagrams of (a & b) preparation processes for thenanoparticle-decorated electrode and core–shell structured elec-trode including solution infiltration and in situ reactive sintering, andORRs at (c & d) traditional electrodes and (e) core–shell structuredelectrode.




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conductivity of the electrode (Fig. 13e); consequently, improvedelectrode activity was realized as compared to normalcomposite electrode and conventionally inltrated electrodes.In addition, the core–shell structured electrode congurationeffectively improved the thermal stability by avoiding thefurther sintering and thermomechanical stress (Fig. 13c)because of the similar thermal expansion to the electrolyte; as aresult, an excellent operational stability was also achieved.

4. Conclusions

In conclusion, a core–shell structured cathode containing aconnective lm shell (Sm2+, Ce4+ and SC) with a cubic perov-skite structure and a stable porous scaffold (core, SDC) with acubic uorite structure was successfully fabricated by acombined approach (inltration and reactive sintering). Theas-prepared core–shell structured electrode shows high elec-tronic conductivity, chemical and thermal compatibilitybetween components, thus producing excellent electro-chemical activity (1746 mW cm�2 at 750 �C and 0.076 U cm2 at650 �C) and maintaining short-term stability (over 200 h withno degradation at 700 �C) for ORR. The combined method wasdemonstrated to be effective for the development of electrodesfor SOFCs for practical applications, which may accelerate thecommercialization of this attractive technology.

Acknowledgements

This work was supported by the “National Science Foundationfor Distinguished Young Scholars of China” under contract no.51025209. The authors acknowledge an ARC future fellowship.

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3D core–shell architecture from infiltration and beneficial reactive sintering as highly efficient...

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