Solid State Ionics 227 (2012) 46–56

Contents lists available at SciVerse ScienceDirect

Solid State Ionics

j ourna l homepage: www.e lsev ie r .com/ locate /ss i

Chemical stability of La0.6Sr0.4CoO3−δ in oxygen permeation applications underexposure to N2 and CO2

Vincenzo Esposito ⁎, Martin Søgaard, Peter Vang HendriksenDepartment of Energy Conversion and Storage, Technical University of Denmark—DTU, Roskilde, DTU Risø Campus, DK-4000, Denmark

⁎ Corresponding author.E-mail address: vies@risoe.dtu.dk (V. Esposito).

0167-2738/$ – see front matter © 2012 Elsevier B.V. Allhttp://dx.doi.org/10.1016/j.ssi.2012.08.015

a b s t r a c t

a r t i c l e i n f o

Article history:Received 29 February 2012Received in revised form 6 August 2012Accepted 19 August 2012Available online 23 September 2012

Keywords:PerovskitesLa0.6Sr0.4CoO3−δ

Oxygen Transport MembranesSrCO3

Phase stability and chemical reactivity of (La0.6Sr0.4)0.99CoO3−δ (LSC64) was tested in oxidative (pO2=0.21 atm) and slightly reducing conditions (pO2~10−5 atm), as well as in carbon dioxide (pO2~10−4 atm)to evaluate the material performance for oxygen separation technologies. Thin film LSC64 oxygen separationmembranes (20–30 μm) were manufactured and electrochemical performance was evaluated at a range oftemperatures with either nitrogen or CO2 purged on the permeate side of the membrane. Material stabilitywas also investigated by high temperature X-ray diffraction, TGA and conductivity measurements in air, N2

and CO2. Under mild reduction LSC64 partly decomposes to a K2NiF4-type phase (i.e. (La,Sr)2CoO4), andCo-oxide, and under high pCO2 forms SrCO3. The latter is found to impair membrane performance. Electricalproperties and oxygen permeation (jO2) in thin membranes depend on the thermal and chemical history ofthe samples. A flux of 4–6 Nml min−1 cm−2 in the temperature range of 800–900 °C was demonstrated foroptimized membranes and conditions.

© 2012 Elsevier B.V. All rights reserved.

1. Introduction

Lanthanum strontium cobaltite, La1−xSrxCoO3−δ (LSC) perovskitematerials (i.e. ABO3 type materials where A=La, Sr and B=Co) havefor decades been the subject of significant interest in many researchfields such as in fuel cells, chemical reactors, oxygen separation mem-branes and for magnetic applications [1–24]. For oxygen transportmembrane (OTM) applications, the benefits of using LSC are strictlyrelated to the material performance and durability [22]. Although im-pressive performances have been demonstrated with LSC basedmembranes, the current technology is still just at a level of scientificresearch and there is still some uncertainty on whether the materialpossesses sufficient chemical stability for technical application inthe area [23].

Performance degradation of LSC has been reported both at low oxy-gen activity and high temperatures [24–28]. The phase decompositionof the La1−xSrxCoO3−δ perovskite phase has been extensively studiedin a range of oxygen activities and in the presence of CO2. It has beendemonstrated that under strong reduction at high temperatures com-plete decomposition into La2O3, SrO and CoO occurs [15,25,27,28].However, other forms of phase instability are possible. Generally, in-crease of alkaline earth content on the A site (e.g. Sr, for x≥0.4) leadsto an increased instability of the perovskite towards reductionwith pos-sible formation of related phases such as Brownmillerite (A2B2O5)[25,26] or K2NiF4-type structure (n=1 of the Ruddlesden–Popper

rights reserved.

series An+1BnO3n+1) [24–27]. The decomposition of LSC has also beenstudied in oxygen gradients. For an LSC73 perovskite membrane itwas observed that SrO preferentially segregates at the low pO2 side[29]. Thiswas ascribed to de-mixing processes inwhich, under the elec-trochemical potential gradient, heterogeneous diffusion of the cationstakes place [30].

The stability of the perovskite phase is further affected by presenceof reactive gases. Particularly, exposure to CO2 can lead to formationof SrCO3 at the surface modifying the stoichiometry of the material.Strontium carbonate decomposes in air at a temperature above 900 °C[31–35] but its stability highly increases if the CO2 activity (pCO2) inthe environment increases [31,35]. Use of thematerial in oxyfuel routescouldwell involve exposure to high CO2 concentrations if one considers“direct integration routes”, where exit gas from the boiler is purgedback over the oxygen membrane [22,23]. Formation of a surface cover-ing layer of Sr-carbonate will impair the oxygen sorption–desorption atthe permeate side and consequently limit the oxygen permeation[36–38]. All these phenomena are even further complicated as the ki-netics of phase transformations in membrane applications are affectedby the oxygen flux through the membrane. The permeation can thussupply oxygen reducing the de-mixing driving forces or stabilizing theperovskite phase also at the low pO2 side. Therefore, although LSC ma-terials have been intensively studied, a clarification of whether thesechemical transformations are prohibitive for use of the material in oxy-genmembranes is necessary. In this study, the La0.6Sr0.4CoO3−δ (LSC64)stability was investigated for selected thermal cycles under slightly re-ducing conditions (pO2=10−5 atm) and under CO2 with the aims toestablish the phase stability window as well as the consequences of a

47V. Esposito et al. / Solid State Ionics 227 (2012) 46–56

partial decomposition on transport properties and finally to correlateobserved oxygen permeation measurement with phase stability.

2. Experimental

2.1. Sample preparation

LSC powder with composition (La0.6Sr0.4)0.99CoO3−δ (LSC64) wassupplied by EMPA (Switzerland). This was calcined at 1200 °C for 4 hand milled in a planetary mill for 5 min to reach a specific surfacearea of ~7 m2/g and average agglomerate size of about 1.5 μm. Tape‐casting slurries for the preparation of dense membranes and poroussupports were manufactured from the powder as described in otherpapers [39–41]. The slurries were tape-cast with a blade clearancefrom 70 to 300 μm for membranes and 300–500 μm for the support.Green tapes were then dried at room temperature for 4 days and co-laminated at 130 °C to form green body components for dense thinmembranes (b30 μm) supported on porous LSC64 support (500 μm)and self-standing dense “thick” samples (135 μm) in several sizes andshapes. Samples were sintered in a single step thermal treatment at1230 °C for 1.5 h in stagnant air. Gas tightness of the dense membranewas verified by a vacuum-leak test developed in house [40]. For thesupported membranes a porous activation layer of LSC64 was screen‐printed on the dense membrane and subsequently sintered at 1000 °Cin stagnant air. The LSC64 ink used for screen printing was preparedfrom fine LSC64 powder as described in reference [41]. This additionalporous layer serves as a catalyst that reduces the resistance associatedwith the oxygen desorption reaction.

2.2. Samples characterization

Porosity, membrane thickness and the microstructures of the layerswere investigated by SEM. Low resolution imaging was performedusing a Hitachi TM-1000. A Carl Zeiss SUPRA 35 FEG-SEM was usedfor high resolution imaging. Elemental analysis was carried out usingenergy-dispersive X-ray spectroscopy (EDS) with a Thermo ElectronCorporation detector. Pore size and porosity of the support were evalu-ated by Hg intrusion (Micromeritics, AutoPore IV 9510) and by SEMusing image analysis software (ImageJ).

The phase stability of the perovskite phase of the calcined powderand of sintered samples was analyzed by X-ray diffraction. Phase purityand lattice parameters of the perovskites were determined using aBragg–Brentano θ−2θ X-ray diffractometer with a Cu Kα source(1.54056 Å), over an angular range of 2θ=15–85°. XRD phase analysiswas performed on a Bruker AXS D8Discover, both for the calcined pow-ders and the sintered tape-cast samples. XRD at high temperatures(HT-XRD) was performed in a hot chamber (MRI, hot-humidity cham-ber) placing the powder on a Pt heating filament. Diffractograms wereobtained in air at 30, 400, 600, 800, and 1000 °C in synthetic air, N2

and CO2. A step size of 0.02° and an acquisition time of 2 s at eachstep were used. Pt was used as an internal standard. Also time resolvedXRD analysis was carried out. Stability was followed over long term(hours or days) to ensure that stable phases had formed. For the tempo-ral XRD a step size of 0.049° and a total acquisition time of 176 s wereapplied,with 30 min intermissions between each scan. The oxygen par-tial pressure at the outlet was measured by an yttria stabilized zirconiapO2 sensor.

Simultaneous thermogravimetry and differential thermal analysis(TG/DTA, Seiko TG/DTA 320U and Netzsch STA 409CD) were performedon powders with open Pt crucibles between room temperature and1000 °C with treatments in synthetic air, N2 and CO2. To obtain a higherprecision and verify reproducibility on themass losses, an enhanced pre-cision thermo-microbalance (Netzsch TG 439) was used with open Ptcrucibles using the same conditions of temperature and atmospheres.Alumina powder and YSZ powder (of equivalent volume to the samples,measured under identical conditions) were used as a reference, in order

to correct the measured weight loss of the perovskites for buoyancyeffects. Multistep temperature-gas sweep cycleswere carried out involv-ing: 1) RT-1000 °C ramp in air, N2 or CO2; 2) dwell at 1000 °C for 1 h inair, N2 or CO2; 3) 1000–400 °C ramp in CO2; 4) dwell at 400 °C for 3 h inCO2; 5) 400–1000 °C ramp in CO2 or air; 6) dwell at 1000 °C for 1 h inair; 7) 1000 °C-RT ramp in air. Heating and cooling rates were 2, 2.5, 5,7.5 and 10 K/min. Total flows were between 100 and 300 Nml min−1.The oxygen concentration at the outlet stream with different gases wasmonitored by a pO2 sensor. Activation energy calculation for carbonatedecomposition in LSC64 powders pre-treated in CO2 was estimated by“dynamic thermal analysis of the solid state reaction” [34]. The methodwas applied on LSC64 powders previously treated 1 h in CO2 at1000 °C, cooled to 400 °C and treated at 400 °C for 4 h in CO2 to promotecarbonate formation. The carbonates decomposition was then carriedout on a continuous heating from 400 °C to 1000 °C in air, using 4 differ-ent heating rates (2.5, 5, 7.5 and 10 K/min).

2.3. Electrochemical characterization

Conductivity was measured by using the van der Pauw method(4-probes), on a dense disk sample with a diameter of 2.5 cm and athickness of 135 μm. Pt paste and Pt-wires were used for contacting.Van der Pauw experiments were carried out using slow heating andcooling ramps (i.e. 0.2–0.5 K/min) between RT and 950 °C under expo-sure to various gases: synthetic air, N2 and CO2. Fig. 5a shows a schematicrepresentation of the experimental setup on a square sample.

Permeation experiments were carried out using a three chambersetup. The membrane was placed between two alumina tubes with aninner diameter of 10.25 mm resulting in an area exposed to the gasesof 0.8 cm2. Silver rings (100 μm thickness) were used for sealingbetween the alumina tubes and the membrane. Gas was delivered toboth the low pO2 and high pO2 sides of the membrane via thosetubes. Thermocouples were placed within the inner alumina tubesuch that they were in contact with the membrane. Fig. 6a shows aschematic representation of the experimental setup. The gas fed onthe permeate side (either N2 or synthetic air or CO2/N2/O2 mixtures)passed through an oxygen sensor upstream to the membrane. The per-meate outlet gaswas then directed to a secondpO2-sensor. Synthetic airor pure O2 was fed to the high pO2 side (feed side) of the membranewith a flow of 100 Nml min−1. The gas flow to the permeate side ofthe membrane was controlled and measured using mass-flow control-lers. Nitrogen was flowed through a sweep gas compartment aroundthe sample perimeter which ensures that no oxygen is supplied fromthe perimeter of the membrane. The oxygen flux through the mem-brane was calculated from the difference between the oxygen partialpressure and gas flows passing over the membrane (see details on theexperimental setup in Fig. 6a and in reference [40]; details about fluxcalculations in references [42] and [43]).

3. Results and discussion

3.1. Phase characterization

Phase characterization was carried out by X-ray diffraction anal-ysis. Although XRD techniques are reliable for such a purpose, it isa problem that the threshold for phase detection is rather high (sev-eral volume %) especially for compounds, like cobalt oxide, whichcan absorb Cu Kα X-rays. Moreover, in the dynamic process ofphase transformation, it can be difficult to distinguish the equilibriafrom metastable phases particularly at low temperatures and formassive samples where kinetics can be slow and transformationscan be also limited by diffusion processes. Due to these consider-ations, characterization of cobalt oxide phases was also carried outby other elemental analysis techniques and stabilization of thephase, at different temperatures and chemical conditions, wastracked over “long term” (hours or days). Fig. 1a shows selected

Fig. 1. (a) XRD patterns for LSC64 powders at 30 °C in air, at 1000 °C in air, and at 1000 °C in N2; (b) XRD patterns for LSC64 collected at 1000 °C in 10 h time frame registering thephase transformation when pO2 is changed from 0.21 to 10−5 atm at t=0 h.

48 V. Esposito et al. / Solid State Ionics 227 (2012) 46–56

49V. Esposito et al. / Solid State Ionics 227 (2012) 46–56

XRD patterns for LSC64 powders at 30 °C in air, at 1000 °C in air(pO2=0.21 atm), and at 1000 °C in N2 (pO2=10−5 atm). Fig. 1bshows a dynamic phase transformation at 1000 °C, where at timezero the atmosphere was changed from air to N2. The diffractogramsshow that the perovskite phase is stable from RT to 1000 °C in air butthat it partly decomposes in N2. The LSC64 perovskite (P) showsrhombohedral symmetry at room temperature as also reported inreference [44] (Millers indices are given in Fig. 1a, lower curve). At1000 °C peaks have shifted to lower angles and some of the peakshave disappeared consistent with a change from rhombohedral tocubic symmetry. The shift of the peaks is attributed thermal andchemical expansion related to oxygen vacancy formation [45]:

“Oxygen loss reaction”

La0:6Sr0:4CoO3−δ→La0:6Sr0:4CoO3−δ′ þ12

δ′−δ� �

O2 gð Þ ð1Þ

La1−xSrxCoO3−δ is at room temperature and high pO2 close to stoi-chiometric (δ=0) [28]. Fig. 1a shows that under reduced oxygen partialpressure in N2 (pO2=10−5 atm) the perovskite (P) transforms to a Ki2-NiO4-type phase (La,Sr)2CoO4 (Ruddlesden–Popper structure: R–P):

“P↔R–P reaction”

2 La; Srð ÞCoO3−δ′→pO2¼10−5atm

←pO2¼0:21 atm

La; Srð Þ2CoO4 þ CoO

þ 0:5−δ′� �

O2 gasð Þ

ð2Þ

In principle both (La,Sr)2CoO4 and CoO may be non-stoichiometric(2) but for simplicity reasons this is not considered here. As the perov-skite decomposes according to reaction (2) CoO is formed, however,this is not observed in the diffractograms due to low counting or the pos-sible X-ray sorption of this phase. Fig. 1a shows that at 1000 °C in N2, theperovskite phase was not completely transformed to the R–P over2 days indicating that a limitation in the decomposition kinetic or equi-librium between the three phases was achieved in our experiment. Thedynamic analysis in Fig. 1b showed that the R–P phase in the materialwas formed by a gradual transformation of the starting P phase at hightemperatures and low pO2 over 10 h whereafter no further changeswere observed. This indicates that the three phases; R–P, P, CoO are inequilibrium, which is in agreement with recent calculated phase dia-grams for the La,Sr,Co,O system [53]. Contrary to the slow P to R–Pdecomposition, the inverse R–P to P transformation at 1000 °C wasvery quick and took only a few minutes when the pO2 was changedfrom 10−5 back to 0.21 atm.

The XRD analysis further showed that the R–P phase formed at1000 °C in N2 could be preserved in the material at reduced tempera-ture; 800 °C>T>400 °C when the pO2 remained constantly low at~10−5 atm. In air the R–P phase prevails below 400 °C. This observationillustrates that phases created at high temperatures and low oxygen ac-tivity, can be maintained as metastable phases even in the typical tem-perature stability domain of the perovskite as long as oxygen is notsupplied or if the cation mobility is low. In the perspective of usingLSC64 in membrane technology, it is noteworthy that a certain amountof R–P phase can thus be “frozen” in thematerial when the temperatureis reduced and the material is maintained under reducing conditions.

Similar considerations can be extended to the discussion of the reac-tivity of LSC64 toward CO2. Fig. 2 shows XRD patterns of LSC64 powdersat 30 °C in air, at 800 °C in CO2 (pO2=10−4 atm), and at 1000 °C inCO2. As expected, LSC64 reacts with CO2 at temperatures around 800 °Cforming an orthorhombic SrCO3 phase. Peaks ascribable to a cubic CoOphase were also observed in the diffractograms at 800 °C. Hence, in CO2

the perovskite decomposes according to:

“P carbonatisation”

La0:6Sr0:4CoO3−δ′ þ CO2 gasð Þ→T¼800 °C La0:6Sr0:4−xCo1−xO3−δ″

þ xSrCO3 þ12

δ″−δ′� �

O2 gasð Þ þ xCoO ð3Þ

The observation of CoO together with the carbonate is consistentwith the finding in literature that the LaSrCoO3 perovskites only toler-ate very little deviation from perfect A/B-site stoichiometry withoutforming secondary phases [51], i.e. when more than ~2% of Sr isextracted from the perovskite by carbonate formation a rebalancingamount of CoO will also be formed.

Moreover, as the pO2 in the CO2 stream is ~10−4 atm the treatmentalso leads to amild reduction as discussed above for the treatment in N2

and it is likely that oxygen vacancies are also formed in thematerials. InCO2 the overall products were: R–P formed by the “P↔R–P reaction”(2), CoO by the “P↔R–P reaction” reaction (2) and “P carbonatisation”reaction (3), SrCO3 by the “P carbonatisation” reaction (3). Also traces ofother unidentified extra-phase/s were observed. Strontium carbonateand R–P phases were also observed at 1000 °C in CO2. The SrCO3 ispresent in rhombohedral form [46]. In air SrCO3 is usually reported todecompose at temperatures around 850–950 °C [31–35]. The findinghere, that it is stable up to at least 1000 °C in pure CO2, is consistentwith earlier studies [31,35] reporting that at high partial pressure ofCO2 the compound stability domain is extended by formation of therhombohedral form. The transformations were fully reversible; theperovskite reformed when the pO2 was increased to 0.21 atm (andthe pCO2 reduced) at high temperatures. Reversibility of SrCO3 forma-tion in perovskites with similar composition to LSC64, such as LSCF,has been reported in several studies [21,38,47–50]. However, it is notclear from these studies whether the presence of carbonates was detri-mental for membrane performances [21,29].

3.2. Oxygen loss and SrCO3 stability

Fig. 3 shows the weight change of the LSC64 powder in differentgases and under different thermal conditions; Fig. 3a shows TGplots as relative weight percentage and as equivalent change in oxy-gen non‐stoichiometry, from 200 to 1000 °C, in N2, and in CO2. Inthe calculation of the oxygen non‐stoichiometry it has been assumedthat δ=0 at room temperature in air. This assumption has previouslybeen verified by Dalslet et al. [51]. Fig. 3b shows “carbonate decompo-sition curves” from 400 to 1000 ° C obtained with different heatingrates for powders previously treated in CO2 (relative weights above100% of the weight at 400 °C are due to the SrCO3 formation). Theinset in Fig. 3b shows differential reaction rates (here expressed asdα/dt, where α is the ratio between the mass change during the reac-tion and the initial mass of the sample) as a function of the reciprocalabsolute temperature for the 4 scan rates in air.

The presented TGA confirmed the XRD results. Particularly, in Fig. 3a,the relative weight loss in air and in N2 can be related to oxygen vacancyformation, as presented in the “Oxygen loss” reaction (1). The oxygenvacancy concentration at 1000 °C was estimated to be around 0.1 and0.25 in air and N2, respectively. In N2, from approximately 900 °C andupwards, the vacancy concentration most likely includes P-phase turn-ing into non-stoichiometric R–P. However, the vacancy concentrationgiven in the figure does not reflect an equilibrium vacancy concentrationin the P phase. From Fig. 3a it is also observed that in air LSC64 starts los-ing oxygen at around 600 °C, while in CO2 or N2 the onset temperature isas low as 350 °C. Such results are comparable with those reported byMizusaki et al. [28] and confirm that the P phase does not decomposefully at high temperatures in N2, as also seen in Fig. 1b. Fig. 3a showsthat the samples treated in CO2 formed SrCO3 by gas uptake above800 °C.

50 V. Esposito et al. / Solid State Ionics 227 (2012) 46–56

The temperatures of formation and amount of carbonate dependedon the heating rate used for the treatment. A certain kinetic limitationin carbonate formation was also observed by XRD, and the TG resultsin Fig. 3a confirmed that slow processes were involved in the reaction.Carbonate formation could likely be limited by SrO segregation and Codiffusion (reaction (3)). Stability ranges of strontium carbonate are usu-ally identified by decomposition experiments where strontium carbon-ate bulk usually decomposes in a single step reaction [34] and whererate and yield of the decomposition depend on the pCO2:

“SrO-CO2 reaction”

SrCO3⇆SrOþ CO2 gasð Þ ð4Þ

In Fig. 3b, the samples that have previously been treated at 1000 °Cand cooled at different rates to 400 °C in CO2 have formed carbonate indifferent amounts and at 400 °C constitutes between 1% and 2.5% of thetotal weight of the samples, depending on the cooling rate. Fig. 3bshows that the strontium carbonate formed is relatively stable all theway up to 1000 °C at pCO2=1 atm, and no large loss due to theSrO-CO2 reaction (4) is registered. The slight mass loss observed at750 °C can be attributed to partial decomposition of the carbonate orformation of other phases. Oxygen loss from the perovskite and R–Pphase formation is also expected in the same temperature range, asidentified byXRD analysis at 1000 °C in CO2 (Fig. 2). In air the carbonateis observed to start decomposing around 700 °C. The process is

Fig. 2. XRD patterns for LSC64 powders at 30 °C in air, at 800 °C in C

completed at temperatures around 850 °C. Relativeweights at this tem-perature of ~99.5% indicate that carbonate decomposition occurs simul-taneously with other oxygen loss reactions (e.g. loss of oxygen from theperovskite).

The inset in Fig. 3b shows the reaction rates (dα/dt) versus 1000/Tfor different heating rates. The presence of more than one peak foreach curve can be attributed to amultistep reaction. Themain peak rep-resents the SrCO3 reaction advancement with temperature and an acti-vation energy of 288±12 kJ/mol can be calculated from the data. Thisvalue is in good agreement with the results of Li et al. and other valuesreported in literature for SrCO3 decomposition [31–36] although differ-ences in the reaction of bulk SrCO3 reported in literature, and SrCO3

formed in LSC64 powders can be expected because of the different reac-tion mechanisms and material morphology.

3.3. Microstructure and chemical analysis

Fig. 4 shows SEM pictures of a dense LSC64 sample sintered in airat 1230 °C (Fig. 4a: surface), and then treated at 950° for 3 days andcooled in CO2 in 12 h. Numbers marked in Fig. 4b, c, d indicate thephases indentified by EDS analysis to be ① Sr-rich, ② Co-rich, and③ containing La, and Co with traces of Sr. Fig. 4a and b show a directcomparison at the same magnification in order to highlight the effectof long term treatment of the sintered material surface in CO2. Thematerial sintered in air showed typical polycrystalline morphology

O2 (pO2=10−4 atm), and at 1000 °C in CO2 (pO2=10−4 atm).

Fig. 3. (a) TG plots for LSC64 perovskite powder from 200 to 1000 °C in air, in N2 and in CO2 at three heating rates (2.5, 5, 7.5 and 10 K/min); (b) TG plots for LSC64 perovskitepowders previously treated in CO2 from 1000 °C to 400 °C and heated in air and in CO2 from 400 to 1000 °C at different rates (2.5, 5, 7.5 and 10 K/min); (b, inset) differential re-action rates dα/dt in CO2 as function of temperature for the different decompositions rates (2.5, 5, 7.5 and 10 K/min).

51V. Esposito et al. / Solid State Ionics 227 (2012) 46–56

with 1% molar excess of cobalt oxide at the grain boundaries and tri-ple points (arrows in Fig. 4a). Fig. 4b shows that long term treatmentand cooling in CO2 led to a transformation at the surface with severemicrostructural rearrangement and formation of additional phases.

Fig. 4c shows details of the surface at higher magnification. Threedifferent phases were identified by EDS: ① A Sr-rich phase, which isabundant on the material surface and from the XRD results it wasidentified as SrCO3; ② A Co and O rich phase, which is especiallyobserved at the grain boundaries. From the XRD this phase is recog-nized as cubic CoO; ③ A La, Sr and Co‐rich phase being the matrixperovskite phase. Fig. 4c shows a fracture cross-section and parts ofthe surface. Clearly, the additional phases ① and ② are formed

mainly at the surface. Elemental scanning in the bulk at the cross‐sec-tion revealed no presence of phases ① and ②. Segregation of largeamounts of CoO in the material is expected considering that besidesthe excess used in the synthesis, both the P↔R–P reaction (2) andthe carbonatisation reaction (3) are sources of CoO. Particularly, thecarbonatisation reaction occurs preferentially at the surface. Thepresented results are in agreement with what is presented in the lit-erature, particularly, the results on the formation of R–P phases andcarbonates at the surface are comparable with results reported byvan Doorn et al. on LSC73, where SrO enrichment and SrCO3 forma-tion were observed at the material surface as consequence of longterm membrane permeation and exposure to CO2 [29].

Fig. 4. SEM picture of a LSC64 dense sample as sintered in air at 1230 °C (a), treated at 950° for 12 h and cooled in CO2 (b: surface; c: surface detail; d: fracture section). Numbersindicate phases indentified by EDS analysis as Sr-rich (1), Co-rich (2), and containing La, Sr and Co (3).

52 V. Esposito et al. / Solid State Ionics 227 (2012) 46–56

3.4. Electrical characterization

LSC64 is a p-type conductor and thus the electrical conductivity isintimately connected with the oxygen vacancy concentration as chargeneutrality is maintained at all times. Therefore, the above describedchanges in the different thermal and chemical conditions can affectthe electrical properties of the material. The evolution of the total con-ductivity for the different chemical and thermal conditions is describedin Fig. 5.

Fig. 5a shows the experimental setup used to measure the conduc-tivity and SEM fracture cross‐sections of the sample (inset). Fig. 5b il-lustrates cycles in air at 830 °C, until the equilibrium conductivity hasbeen obtained, followed by cooling in air, heating in N2, cooling in N2,and heating in air. The conductivity of the LSC phase was found to beslightly higher than that reported in literature by Petric et al. [50] andothers [27,52,53] (dashed line). However, the first cycle reproducedthe effects of reaction (2) on the conductivity: Heating the materialin N2 leads to a decrease in conductivity from above 600 °C comparedto the conductivity in air. At 975 °C the conductivity decreased from1200 S cm−1 in air to 700 S cm−1 in N2 which is attributed mainlyto loss of oxygen which decreases the electron hole concentrationbut possibly also to the formation of a lower conductivity R–P phasesince (La,Sr)2CoO4 in air and N2 below 1000 °C shows a total conduc-tivity between 10−1 and 102 S cm−1 [52]. The cycle in N2 shows thatthe conductivity was not recovered by simply cooling the sample toRT (segment 2) in N2. In this condition the sample is in a metastablemultiple phase form (perovskite/R–P phase/CoO) as the nitrogenstream does not contain sufficient oxygen for a re-oxidation. A highpO2 treatment is necessary to restore the material properties as evi-dent from the subsequent heating curve in air (segment 3) whereconductivity reaches its equilibrium value at ~600 °C.

Fig. 5c shows a heating and cooling cycle obtained in CO2. Such atreatment reproduces the effects of reaction (3) on the conductivity.

Segment 1 shows that heating in CO2 leads to a decrease of the con-ductivity above 600 °C; Segment 2 shows that the decrease can bepreserved for low temperatures when cooling the material in CO2.The variations in electrical conductivity are comparable to the onesfound in nitrogen. It is thus clear that even though strontium carbon-ate has formed it does not affect the conductivity to a measurabledegree, suggesting that the carbonates only form at the sample sur-face. It should be noted that a sample with a very high surface tobulk ratio has been used in these experiments, wherefore changesin the bulk conductivity should be more easily identifiable.

In Fig. 5d, the second cycle is completed with a further heating inCO2, cooling in CO2, and heating in air for re‐oxidation. The overall treat-ments in CO2 showed that thematerial was reactive and the conductiv-ity decreased both in heating conditions (segment 3) and in coolingconditions (segment 4). Reversibility of the reaction (3) was also con-firmed in segment 5 by the final treatment in air where at the temper-ature range of 300–600 °C the initial properties were fully recovered.

3.5. Oxygen permeation

For the data collected in “static” conditions, the phase formationsshowed dependency only on specific conditions of temperature andpO2. In the case of an operating membrane things are different; the per-meation flux can induce phase changes because of the altered pO2 condi-tions which can be generated locally in the membrane, at the permeateside and at the membrane/active layer interface [43]. Moreover, cationde‐mixing can take place when an electrochemical potential gradient isapplied. To clarify the material stability under membrane applicationpermeation experiments were carried out under varying conditionsand stability was also followed over time.

Fig. 6a shows an illustration of the experimental setup and SEMcross‐section of the sample (inset) used for the oxygen permeationexperiments. For dense membranes of cobalt based perovskites the

Fig. 5. (a) Sample and van der Pauw experimental setup descriptions; (b, c, d) conductivity vs temperature plots for dense LSC64. (b) Sample equilibrated in air (broken line), heating in N2 (segment-1), cooling in N2 (segment-2), heating inair (segment-3); (c) treated in air, heating in CO2 (segment-1), cooling in CO2 (segment-2); (d) heated in CO2 (segment-3), cooling in CO2 (segment-4), and heating in air (segment-5). Insets show schematic representation of the cycles.Dashed lines correspond to conductivity values reported by Petric et al. [50]. 53

V.Esposito

etal./

SolidState

Ionics227

(2012)46

–56

Fig. 6. (a) Sample and setup descriptions for oxygen permeation experimental; (b) oxygen permeation for dense LSC64 20 μm thick sample supported on porous LSC64 substrate(feed side) and porous LSC64 screen-printed electrode (active layer) treated in N2 and CO2 with temperatures.

54 V. Esposito et al. / Solid State Ionics 227 (2012) 46–56

characteristic length is usually estimated to be in the 100 μm range[50]. The present oxygen permeation experiments have been carriedout on a dense LSC64 membrane (20 μm thickness) supported by aporous LSC64 substrate feed side with a 50 μm LSC64 screen-printed porous electrode on the permeate side. Although the mem-brane is equipped with activation layers/electrodes that improvesthe sorption–desorption and thus decreases the characteristic length[51], typical polarization resistance of LSC based electrodes at the per-meate side of the membrane is the main loss in driving force [41,53].Therefore, the oxygen flux through the membrane does heavilydepend on the sorption–desorption processes at the material surface.Fig. 6b shows the measured oxygen permeation during heating in N2

and CO2 with an intermediate recovering treatment in air. The mem-brane treated for 150 h in N2 and then 100 h in CO2 at the permeateside and pure O2 at the feed side, showed comparable performances

with a maximum oxygen flux of 4–6 Nml min−1 cm−2 at 900 °C.Use of N2 with pO2=10−5 atm at the permeate side is observed tolead to slightly higher fluxes than with CO2 (pO2=10−4 atm)which is consistent with the expectations according to Wagnertheory.

Although the membrane in CO2 showed high flux values at tem-peratures higher than 850 °C, a fast degradation was observed at tem-peratures lower than 800 °C. This is illustrated in Fig. 7a that shows along term experiment at 780 ° C with CO2 at the permeate side (aftera test in N2). The performance slowly decayed over four days with adrastic reduction of the oxygen flux from 2.2 to 0.5 Nml min−1 cm−2.Such effect on the membrane performance is consistent with TGA re-sults and can be attributed to a stable carbonate forming on the surfaceat high pCO2. In principle the decaying flux values could also be due todetrimental effects of a cation de‐mixing. This is however unlikely as

Fig. 7. Permeation experiments for dense LSC64 20 μm thick sample supported on porous LSC64 substrate (feed side) and porous LSC64 screen-printed electrode (active layer)treated in N2 and CO2 with time at 800 °C (a) and at 915 °C (b).

55V. Esposito et al. / Solid State Ionics 227 (2012) 46–56

the driving force for this is even stronger in N2 and here nodecay in per-meation fluxes was observed over time.

The reversibility of the “P↔R–P reaction” (2) and “P carbonatisation”(3) was also observed and verified in the membrane test. Fig. 7b showsoxygen fluxes at 915 °C where the permeate atmosphere was changedfrom CO2 (step 1), to air (step 2), to N2 (step 3), and back to CO2 (step4). Consistent with the results in Fig. 7a, at the first stage the membranetreated in CO2 at 650 °C showed a reduced performance with a flux ofaround 4 Nml/min cm−2 at 950 °C. The steady conditions of the mem-branes in CO2 at 915 °C indicates that the carbonates formed at 650 °Cdid not decompose even with a significant oxygen supply through the

membrane. This test is a further demonstration that, as indicated in reac-tion (5), carbonates are stable at high pCO2 and the membrane perme-ation depends on the carbonate stability. After a treatment at 950 °C inair at the permeate side for 12 h the flux dramatically increases in bothN2 and CO2, illustrating that the formed carbonate previously impairingthe flux has now decomposed regenerating the P phase. Here the oxygenflux in CO2 at 915 °C after the perovskite phase had been restored wasfound to be fairly stable. A direct comparison with the first step showshow the carbonate at high temperature and higher oxygen flux cannotbe reformed. If oxygen was not permeating through the membrane oneshould see carbonate formation at 915 °C demonstrating a substantial

56 V. Esposito et al. / Solid State Ionics 227 (2012) 46–56

influence of the dynamic permeation conditions on the carbonate forma-tion for high oxygen fluxes and high operation temperatures (T>850–900 °C).

The typical time needed to recover the performance in air was lessthan 12 h. This suggests that the phenomenon is related to the surfaceand is not due to cation de‐mixing. Moreover, electrical measurementson the different phase formed also points out that different phase havesimilar conductivities and any limitations in the permeation can bemainly attributed to the increase of oxygen desorption polarization.On the other hand, a decay of the membrane performance is observedduring step 3 in N2 resulting in a slow degradation at around 915 °Cand pO2=10−5 atm of about 1% in 10 h. However, the effect of N2 onthe oxygen permeated by the membrane is difficult to estimate andsuch effect can be either due to cation de‐mixing or slow degradationof the P phase as pointed out in the XRD analysis. These results are con-sistent with other papers on degradation which showed membraneperformance stability of LSCF6428 and LSC73 at 900 °C under similarconditions [21,29].

The post-analysis observation carried out on themembrane showedthat the samples maintained a remarkable microstructural integrity aswell as mechanical robustness although the severe chemical and ther-mal cycles performed in the permeation tests.

4. Conclusions

Chemical stability and oxygen permeation performance of LSC64perovskite phase (P) were tested in slightly reducing conditions andin CO2 to verify the reliability of the material in oxygen permeationapplications. LSC is unstable at high temperature and under reduced ox-ygen activity. The P phase decomposes with formation of Ruddlesden–Popper-type phases (R–P) and consequent segregation of CoO. SrCO3 isalso formed as a result of SrO segregation and reaction with CO2. Thisphenomenon leads to sequestration of Sr from the P phase and furtherformation of CoO. The additional CoO and SrCO3 are mainly localized atthe surface of the material. SrCO3 can easily be decomposed at temper-atures above 750 °C in air. However, SrCO3 is stable at pCO2=1 atm atlow temperatures and R–P phase persisted asmetastable phases at lowtemperatures under low oxygen activity. As result, LSC is relativelyunstable in the operative conditions of membranes. Its use in these ap-plications should consider the risk of a drastic decay of the performancewhen the perovskite phase decomposes into R–P and/or carbonate isformed under CO2 exposure. Use of CO2 is particularly detrimentalbelow 800 °C. The thermal and chemical history of the material canalso affect the performance of the membrane because the decomposi-tion product can be present as metastable phases.

A careful use of the material allows a long term integrity reachingremarkable performance also after several chemical and thermal cycles.In this work, 20 μm LSC thin membranes supported on porous LSCshowed permeation fluxes between 4 and 6 Nml min−1 cm−2 at900 °C at very lowdriving force (permeation side pO2~10−5 atm). Per-formance was stable over the periods examined as long as the perov-skite phase is preserved.

Acknowledgements

S.P.V. Foghmoes is acknowledged for tape-casting and slurriespreparation, K. Tydén for EDS analysis and N. Bonanos for discussionand suggestions on van der Pauw experimental setup.

References

[1] B.C.H. Steele, Solid State Ionics 134 (2000) 3.[2] P. Hjalmarsson, M. Søgaard, M. Mogensen, Solid State Ionics 179 (2008) 1422.[3] Y. Teraoka, T. Nobunaga, N. Yamazoe, Chem. Lett. 17 (1988) 503.[4] H.J.M. Bouwmeester, A.J. Burggraaf, in: P.J. Gellings, H.J.M. Bouwmeester (Eds.), CRC

Handbook of Solid State Electrochemistry, CRC Press, Boca Raton, USA, 1997, p. 481.[5] H. Okamoto, H. Obayashi, T. Kudo, Solid State Ionics 1 (1980) 319.[6] P.J. Gellings, H.J.M. Bouwmeester, Catal. Today 12 (1992) 1.[7] W. Zhu, J.M. van de Graaf, L.P.J. van de Broeke, F. Kapteijn, J.A. Moulijn, Ind. Eng.

Chem. Res. 37 (1998) 1934.[8] H.J.M. Bouwmeester, Catal. Today 82 (2003) 141.[9] K. Tsutsui, J. Inoue, S. Maekawa, Phys. Rev. B 59 (1999) 4549.

[10] H.W. Hsu, Y.H. Chang, G.J. Chen, K.J. Lin, Mater. Sci. Eng. B 64 (1999) 180.[11] A.N. Petrov, O.F. Kononchuk, A.V. Andreev, V.A. Cherepanov, P. Kofstad, Solid State

Ionics 80 (1995) 189.[12] A. Thursfield, I.S. Metcalfe, J. Mater. Chem. 14 (2004) 2475.[13] Y. Teraoka, H.-M. Zhang, S. Furukawa, N. Yamazoe, Chem. Lett. 14 (1985) 1743.[14] M. Søgaard, P.V. Hendriksen, M. Mogensen, F.W. Poulsen, E. Skou, Solid State

Ionics 177 (2006) 3285.[15] Y. Teraoka, H.-M. Zhang, N. Yamazoe, Chem. Lett. 14 (1985) 1367.[16] N.P. Bansal, Z. Zong, J. Power Sources 158 (2006) 148.[17] H. Kurokawa, C.P. Jacobson, L.C. DeJonghe, S.J. Visco, Solid State Ionics 178 (2007) 287.[18] M. Sase, J. Suzuki, K. Yashiro, T. Otake, A. Kaimai, T. Kawada, J. Mizusaki, H. Yugami,

Solid State Ionics 177 (2006) 1961.[19] V.V. Kharton, E.V. Tsipis, A.A. Yaremchenko, I.P. Marozau, A.P. Viskup, J.R. Frade,

E.N. Naumovich, Mater. Sci. Eng. B 134 (2006) 80.[20] Y. Liu, L. Hong, J. Membr. Sci. 224 (2003) 137.[21] X. Tan, K. Li, A. Thursfield, I.S. Metcalfe, Catal. Today 131 (2008) 292.[22] R. Bredesen, J. Sogge, Seminar on the ecological applications of innovative mem-

brane technology in the chemical industry, Chem/Se., Cetraro, Calabria, Italy,1996. [cited in]. G. Iaquaniello, F. Giacobbe, B. Morico, S. Cosenza, A. Farace, Int.J. Hydrogen Eng. 33 (22) (2008) 6595–6601.

[23] A.J. Ellett, in: Schriften des Forschungszentrums Jülich Reihe Energie & Umwelt /Energy & Environment Band /1866-1793vol. 43, ISBN: 978-3-89336-581-4, 2009.

[24] K. Watanabe, D. Takauchi, M. Yuasa, T. Kida, K. Shimanoe, Y. Teraoka, N. Yamazoe,J. Electrochem. Soc. 156 (2009) E81.

[25] J. Ovenstone, J.S. White, S.T. Misture, J. Power Sources 181 (2008) 56.[26] L.M. Liu, T.H. Lee, L. Qiu, Y.L. Yang, A.J. Jacobson, Mater. Res. Bull. 31 (1996) 29.[27] A.N. Petrov, V.A. Cherepanov, A.Y. Zuev, J. Solid State Electrochem. 10 (2006) 517.[28] J. Mizusaki, Y. Mima, S. Yamauchi, K. Fueki, H. Tagawa, J. Solid State Chem. 80 (1989)

102.[29] R.H.E. van Doorn, H.J.M. Bouwmeester, A.J. Burggraaf, Solid State Ionics 111 (1998)

263.[30] H. Schmalzried, J.Chem. Soc. Faraday Trans. 86 (1990) 1273.[31] M.J. Scholten, J. Schoonman, J.C. van Miltenburg, H.A.J. Oonk, Solid State Ionics 61

(1993) 83.[32] B.H. Erné, A.J. van der Weijden, A.M. van der Eerden, J.B.H. Jansen, J.C. van

Miltenburg, H.A.J. Oonk, Calphad 16 (1992) 63.[33] A. Reller, R. Emmenegger, C. Padeste, H.-R. Oswald, Chimia 45 (1991) 262.[34] C.-R. Li, T.B. Tang, J. Therm. Anal. 49 (1997) 1243.[35] D.W. McKee, Fuel 59 (1980) 308.[36] H. Yokokawa, N. Sakai, T. Kawada, M. Dokiya, Solid State Ionics 52 (1992) 43.[37] M.A. Penã, J.L.G. Fierro, Chem. Rev. 101 (2001) 1981.[38] G. Juhász, Z. Homonnay, K. Nomura, T. Hayakawa, S. Hamakawa, A. Vértes, Solid

State Ionics 139 (2001) 219.[39] T. Klemens, M. Mogensen, J. Am. Ceram. Soc. 90 (2007) 3582.[40] A. Kaiser, S. Foghmoes, C. Chatzichristodoulou, M. Søgaard, J.A. Glasscock, H.L.

Frandsen, P.V. Hendriksen, J. Membr. Sci. 378 (2011) 51.[41] J. Hjelm, M. Søgaard, R. Knibbe, A. Hagen, M.B. Mogensen, in: ECS transactions, The

Electrochemical Society1938-5862vol: 13, issue: 26, 2008, pp. 285–299, http://dx.doi.org/10.1149/1.3050400.

[42] J. Sunarso, S. Baumann, J.M. Serra,W.A.Meulenberg, S. Liu, Y.S. Lin, J.C. Diniz da Costa,J. Membr. Sci. 320 (2008) 13.

[43] P.V. Hendriksen, P.H. Larsen, M. Mogensen, F.W. Poulsen, K. Wiik, Catal. Today 56(2000) 283.

[44] R. Sonntag, S. Neov, V. Kozhukarov, V. Neov, J.E. ten Elshof, Physica B 24 (1998) 393.[45] N.M.L.N.P. Closset, R. van Doorn, R. Kruidhof, J. Boeijsma, Powder Diffr. 11 (1996) 31.[46] K.O. Strømme, Acta Chem. Scand. Ser. A 29 (1975) 105.[47] Q. Yang, J.Y.S. Lin, Sep. Purif. Technol. 49 (2006) 27.[48] Q. Yang, Y.S. Lin, M. Bülow, AICHE J. 52 (2006) 574.[49] K. Nomura, Y. Ujihira, T. Hayakawa, K. Takehira, Appl. Catal., A 137 (1996) 25.[50] A. Petric, P. Huang, F. Tietz, Solid State Ionics 135 (2000) 719.[51] B.T. Dalslet, M. Søgaard, P.V. Hendriksen, Solid State Ionics 180 (2009) 1050.[52] T. Matsuura, J. Tabuchi, J. Mizusaki, S. Yamauchi, K. Fueki, J. Phys. Chem. Solids 49

(1988) 1403.[53] Weiwei Zhang, Internal communication (2011).