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CXXXX American Chemical Society

Direct Evidence of Stacking Disorder inthe Mixed Ionic-Electronic ConductorSr4Fe6O12þδMarta D. Rossell,†,* Artem M. Abakumov,‡ Quentin M. Ramasse,§ and Rolf Erni†

†Electron Microscopy Center, Empa, Swiss Federal Laboratories for Materials Science and Technology, U.berlandstrasse 129, 8600 Dübendorf, Switzerland,‡EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium, and §SuperSTEM Laboratory, STFC Daresbury Campus, Keckwick Lane,Warrington, WA4 4AD, United Kingdom

Perovskites exhibiting mixed ionic-electronic conducting properties arebeing extensively researched as key

components in advanced energy applications.Especially interesting for high-temperature ap-plications, such as cathodes in solid oxide fuelcells, are perovskite-related oxides exhibitingdifferent structural blocks due to their superiorphase stability. Among them, the oxygen-deficient perovskite compounds with the gen-eral formula A4B6O12þδ have been extensivelystudied over the past decade on account oftheir complex mechanism of anion nonstoi-chiometry. Up to now, two representatives ofthis family are known: the Sr4Fe6O12þδ and theBa4In6-x-MgxO12þδ solid solutions.1�7

The A4B6O12þδ structure is built up of AOand BO2 perovskite-type layers alternating

along the b-direction with double B2O2þδ

layers of 5-fold B polyhedra (Figure 1). TheBO5 polyhedra are arranged within the dou-ble layers in bands consisting of double edge-sharing trigonal bipyramids and bands ofedge-sharing tetragonal pyramids with vari-able thickness. These two types of bands areconnected to each other by sharing corners.Because of a strong deviation of the shapeof these polyhedra from perfect trigonalbipyramidal or tetragonal pyramidal, theassignment of a specific geometry hasbeen the focus of extensive discussions.1,4,5

Nevertheless, it is widely accepted that withvarying oxygen content the coordinationnumber of the B-cations does not change.Instead, variation of the oxygen content isaccommodated by changing the thickness

* Address correspondence tomarta.rossell@empa.ch.

Received for review October 24, 2012and accepted March 4, 2013.

Published online10.1021/nn3058449

ABSTRACT Determining the structure-to-property relationship

of materials becomes particularly challenging when the material

under investigation is dominated by defects and structural disorder.

Knowledge on the exact atomic arrangement at the defective struc-

ture is required to understand its influence on the functional prop-

erties. However, standard diffraction techniques deliver structural

information that is averaged over many unit cells. In particular,

information about defects and order�disorder phenomena is contained in the coherent diffuse scattering intensity which often is difficult to uniquely

interpret. Thus, the examination of the local disorder in materials requires a direct method to study their structure on the atomic level with chemical

sensitivity. Using aberration-corrected scanning transmission electron microscopy in combination with atomic-resolution electron energy-loss spectroscopy,

we show that the controversial structural arrangement of the Fe2O2þδ layers in the mixed ionic-electronic conducting Sr4Fe6O12þδ perovskite can be

unambiguously resolved. Our results provide direct experimental evidence for the presence of a nanomixture of “ordered” and “disordered” domains in an

epitaxial Sr4Fe6O12þδ thin film. The most favorable arrangement is the disordered structure and is interpreted as a randomly occurring but well-defined

local shift of the Fe�O chains in the Fe2O2þδ layers. By analyzing the electron energy-loss near-edge structure of the different building blocks in the

Sr4Fe6O12þδ unit cell we find that the mobile holes in this mixed ionic-electronic conducting oxide are highly localized in the Fe2O2þδ layers, which are

responsible for the oxide-ion conductivity. A possible link between disorder and oxygen-ion transport along the Fe2O2þδ layers is proposed by arguing that

the disorder can effectively break the oxygen diffusion pathways.

KEYWORDS: modulated structure . oxygen-deficient perovskite . mixed ionic-electronic conducting oxide . ferrite .electron energy-loss spectroscopy . scanning transmission electron microscopy

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of the bands consisting of tetragonal pyramids, that is,by changing the ratio between edge- and corner-sharing BO5 polyhedra. Thus, varying the oxygen con-tent in these phases changes the polyhedral sequencein the B2O2þδ layers along the a-direction resulting inthe formation of incommensurately modulated struc-tures. These incommensurate modulations were firstidentified in Sr4Fe6O12þδ thin films using electrondiffraction.2 By applying a combination of electrondiffraction data and high-resolution transmission elec-tron microscopy (HRTEM) imaging, structural modelsfor the Sr4Fe6O12þδ phases were proposed in which amodulation vector q = Ra* was identified to changewith δ, so that δ = 2R.2,4 The exact configuration of themodulated Fe2O2þδ double layers was later deter-mined from single-crystal X-ray diffraction data.5 How-ever, X-ray diffraction data also evidenced the pres-ence of diffuse scattering along the stacking directionb*. This phenomenonwas suggested to be caused by astacking disorder of the Fe2O2þδ double layers.5 In-deed, different stacking sequences were also observed

in the Sr4Fe6O12þδ solid solution by HRTEM imaging,4

but the exact stacking scheme of the Fe2O2þδ layerswas difficult to derive due to the possibility of imagedelocalization affecting phase contrast micrographs inconventional high-resolution electron microscopes.Thus, despite the fact that there is clear evidence oflocal ordering in the Fe2O2þδ layers, the exact atomicstacking structure of the Sr4Fe6O12þδ phases is not yetfully understood.Here, we investigate the local crystal and electronic

structure, and the structural defects of a Sr4Fe6O12þδ

film employing high-resolution aberration correctedscanning transmission electronmicroscopy (STEM) andelectron energy-loss spectroscopy (EELS). Because ofthe incoherent nature of the signal, high-angle annulardark-field (HAADF) STEM imaging provides directlyinterpretable and chemically sensitive images. Besides,as STEM can be combined with EELS to map theelemental distribution at atomic resolution, STEM-EELShas become a powerful tool to study the structure ofmaterials,8�10 and to characterize defects11�15 andinterfaces16�18 on the atomic scale.

RESULTS AND DISCUSSION

The (010)-oriented epitaxial Sr4Fe6O12þδ (SFO) filmwas deposited on a SrTiO3(001) substrate by pulsedlaser deposition.19 Details about the deposition condi-tions can be found in the Methods section. Figure 2ashows a HAADF-STEM image of the film/substrateinterface. The HAADF signal (95�195 mrad) scalesapproximately with the square of the atomic numberZ. Thus, the HAADF image of Figure 2a is dominated bythe Sr, Fe, and Ti cations, and to a lesser degree, by theO atoms. At the SFO/SrTiO3 interface a ∼2.5-nm-thicklayer with perovskite structure and a HAADF imageintensity very close to the one of the SFO film can beobserved (Figure 2a). Therefore, its composition can

Figure 1. Sr4Fe6O12þδ (δ = 0.74) crystal structure viewedalong [001]. The Sr atoms are shown as red spheres and theFe atoms are found in the shaded polyhedra. Perovskite-type layers of gray FeO6 octahedra alternate along theb-direction with double Fe2O2þδ layers of FeO5 polyhedra.Double trigonal bipyramids are shown in dark green anddouble and triple tetragonal pyramids in light green.

Figure 2. Antiphase boundary in a SFO film on a SrTiO3 substrate. The presence of a SrFeO3‑x buffer layer of higher inten-sity promotes the formation of an antiphase boundary. (a, b) Cross-sectional HAADF- and ADF-STEM images of theSFO/SrTiO3(001) interface, respectively. (c) Simultaneous HAADF image and atomic-resolution maps of the Sr M4,5, Fe L2,3and O K edges calculated from the EELS images obtained from the area outlined with a yellow square in panels a and b.

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safely be ascribed to SrFeO3‑x. Similar perovskite-typebuffer layers were previously observed in SFO filmsgrown on SrTiO3 and NdGaO3 substrates.20 It wassuggested that during the growth process the firstdeposited layers mimic the perovskite stacking of thesubstrate and a perovskite buffer layer with limitedthickness is formed instead of the stable SFO phase.In the present case the buffer layer thickness is notconstant and it oscillates between 6 and 7 perovskiteunit cells. Therefore, steps on the SrFeO3‑x layer sur-face are created promoting the formation of anti-phase boundaries (APBs) in the SFO film. Thesedefects extending to the film surface with an inclina-tion of ∼53� are frequently observed. In the particularcase of Figure 2, the APB acts also as a permutationboundary, that is, a twin boundary with a supplemen-tary translation vector. Across the permutation bound-ary the a- and c-axes are swapped leaving the b-axisunchanged. In Figure 2a, the visible repeat periodalong the out-of-plane b-axis corresponds to the -SrO-FeO2-SrO-Fe2O2þδ-SrO- sequence of alternatinglayers. However, from the HAADF image of Figure 2ait is hard to distinguish the [001] domain from the [100]domain at each side of the APB. This becomes clearwhen inspecting the annular dark field (ADF) STEMimage of Figure 2b, which was recorded simulta-neously with the HAADF-STEM image. The ADF signal(41�95 mrad) is more sensitive to diffraction contrastand to any effect that leads to a dechannelling of theincident electron beam.21,22 Therefore, strain fieldsfrom oxygen vacancies are readily detected by ADFimaging through the associated displacements of theadjacent cation sites and the subsequent dechannel-ling of the electron beam from the cation columns.23

Indeed, at the left side and at the lower right side ofthe APB, the characteristic modulation of the Fe2O2þδ

layers along the a-axis is clearly visible as a hexagonalpattern (this pattern reproduces the polyhedral se-quence shown in the model of Figure 1). The satellitereflections generated by this modulation are used todetermine the oxygen composition of the film. The experi-mentallymeasuredR=0.37componentof themodulationvector allows us to derive a Sr4Fe6O12.74 film compositionfollowing the relation δ = 2R given in reference 2.The formation of APBs may impact the oxide-ion

conductivity of the films which is assumed to be strictlytwo-dimensional, confined to the Fe2O2þδ doublelayers.24�26 To understand the structure of the APBsand their influence on the conducting properties of epi-taxial SFO thin films, we performed atomic-resolutionEELS mapping of the elements at the vicinity of theAPB. Such an EELS image was collected from themarked yellow region in Figure 2a,b by using a dedi-cated aberration-corrected scanning transmissionelectron microscope operated at 100 kV (for detailssee Methods section). After noise reduction by principalcomponent analysis (PCA) using multivariate statistical

analysis,27 the background in all spectra was subtractedby fitting a decaying power-law function to an energywindow just in frontof thedifferent core-loss edge onsets.Figure 2c shows the elemental atomic-resolutionmapsof the SrM4,5, Fe L2,3, and O K edges extracted from thespectrum image together with the simultaneouslyacquired HAADF image. It is noteworthy that theresolution of the EELS map even enables to resolvethe adjacent Fe atomic columns of the Fe2O2þδ doublelayers separated by a projected distance of ∼1.5 Åalong the b-axis. The EELS data reveal that the APBinterrupts the lateral continuity of only one-half ofthe SrO layers. On the other hand, all Fe2O2þδ doublelayers are laterally interrupted and connected to FeO2

perovskite-type layers and vice versa. Therefore, we cananticipate that the presence of APBs with a translationvector R≈ 1/5[010]bwill have a negative impact on theoxide-ion conductivity of the SFO films. It should benoted that APBs with a similar atomic arrangementare present in the bulk Tl- and Pb-doped SFO sys-tems, forming in the latter case an unusual terrace-like structure.28,29

Another striking feature can be observed in the [100]domain displayed in the ADF image of Figure 2b:successive double Fe2O2þδ layers show substantialintensity variations. Figure 3 panels a and b show aHAADF- and the corresponding ADF-STEM image offour double layers imaged along the [100] direction.The Fe atoms of the three upper double layers arearranged in a zigzag pattern delimited by brighter Sratomic columns (Figure 3a) while in the lower doublelayer the zigzag pattern is blurred. This layer appearsbrighter in the ADF image of Figure 3b indicatingdifferent channeling conditions caused by a distinctatomic reconstruction. Elemental maps extracted froman EELS image obtained from the same area (Figure 3c)prove that the Fe atomic columns in the lower layer aredisordered when compared to the three upper doublelayers.

Figure 3. Atomic-resolution chemical imaging of SFO alongthe [100] direction. (a, b) High-resolution HAADF- and ADF-STEM images, respectively. (c) Simultaneous HAADF andADF images and elemental atomic-resolution maps calcu-lated from the EELS images. The RGB map is generatedusing red for Sr, green for Fe, and blue for O.

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Tobuildpossiblemodels for theorderedanddisorderedvariants, we adopted the local atomic arrangements of theFe2O2þδ layer derived from the incommensurately modu-lated Sr4Fe6O12þδ structure refined from single-crystalX-ray diffraction data.5 Each Fe2O2þδ layer can be con-sidered as a pair of adjacent FeO1þδ/2 layers. The FeO1þδ/2

layers are in turn composed of Fe�O columns runningalong the a-axis (Figure 4). Along the c-axis the chainsare separated by short (s) and long (l) distances.Figure 4 panels a and b show identical configurationsof the adjacent FeO1þδ/2 layers but with the orderedpattern of alternating s and l separations shifted by c/2(this can be seen from the orange rectangles outliningthe unit cells in both figures). Being viewed along [100],such ordering results in a zigzag arrangement of theprojected Fe�Ocolumns. The disordered arrangementarises if the ...s-l-s-l-s... pattern is locally violated byinsertion of a fragment shifted by c/2 and mirroredwith respect to the (001) plane, as is shown, for ex-ample, in Figure 4c. This mixed configuration causesthe apparent disorder in the projected Fe�O columnsobserved in the [100] STEM images. The HAADF imagesshown in Figure 5a,b are obtained from two different

areas exhibiting respectively a disordered and a zigzagarrangement of the Fe2O2þδ layers. The crystal struc-ture models of both variants were obtained from thecommensurate q = Ra* = 3/7a*≈ 0.43a* approximantsof the modulated Sr4Fe6O12þδ structure,

5 adopting thecorresponding subgroup allowing for an ordered ...s-l-s-l-s... alternation of the separations between the Fe�Ochains. The excellent match between the experimentaland the calculated HAADF images (Figure 5a,b) con-firms the proposed structure models of the orderedand disordered Fe2O2þδ layers. Furthermore, consecu-tive ordered Fe2O2þδ layers are seen to pile up follow-ing two different stacking configurations as shown inFigures 3a and 5b. While no shift along the c-axis isobserved between the consecutive zigzag layers ofFigure 3a, the zigzag layers of Figure 5b are shifted byc/2 making the zigzag layers appear mirrored withrespect to a (010) plane located at the center of theperovskite-type layer.For bulk Sr4Fe6O12þδ ceramics the oxygen ionic con-

ductivity at 1173 K varies between ∼6� 10�4 and ∼3�10�4 S 3 cm

�1 for partial oxygen pressures p(O2) = 10�1�105 Pa.30 Lowoxygen ionic conductivity was also foundfor epitaxial Sr4Fe6O12þδ thin films on NdGaO3, LaAlO3

and SrTiO3 substrates, being below the measurablelimit.31 It was noticed that such low ionic conductivityis not in line with ab initio computations and with theionic conductivity in other similar structures such asthe Sr2Fe2O5 brownmillerite32�35 admitting that theorigin of this apparent discrepancy is not yet clear andmay be associated with blocking effects of the grainboundaries in polycrystalline ceramic materials. Theobserved mixed configuration of the Fe2O2þδ layers

Figure 4. Structure of the Fe2O2þδ double layer: the FeO1þδ/2

sublayers. (a) Model of an “ordered” FeO1þδ/2 sublayerconsisting of Fe�O atoms linked through extra oxygens.The chains run along the a-direction and are marked withdashed lines. The Fe atoms are shown as larger greenspheres and the O atoms as smaller blue spheres. The shortand long separations between the chains along the c-axisare marked as s and l, respectively. The unit cell of the com-mensurate approximant is outlined with an orange rectan-gle. (b) The -s-l-s-l-s- pattern in this “ordered” sublayer isshifted by c/2 with respect to the one in panel a (see orangerectangle). (c) FeO1þδ/2 sublayer representing a mixed “dis-ordered” configuration, where a fragment of the layer inpanel a is inserted into the matrix of the layer in panel b.At the right side of the models, simulated HAADF-STEMimages of the corresponding Fe2O2þδ double layers areincluded. The atomic columns highlighted in green in theHAADF-STEM images are obtained by projecting all the Feand O atoms shown in the models along the a-direction. The-s-l-s- and the -s-l-s-l-s- patterns of Figure 2a,b are readilyobservable in thecorrespondingHAADF-STEM images,while thesequence is lost for the “disordered” configuration of Figure 4c.

Figure 5. Atomic reconstructions of the Fe2O2þδ layers inSFO along [100]. (a) Experimental HAADF-STEM image ofthe disordered Sr4Fe6O12þδ structure (left) with the mixedFe2O2þδ layers (as in Figure 4c) and calculated image for a12-nm-thick crystal (right). (b) Experimental HAADF-STEMimage of the Sr4Fe6O12þδ structure (left) with the orderedFe2O2þδ layers (as in Figure 4a,b) and calculated image for a12-nm-thick crystal (right). The corresponding structuralmodels are overlaid on the experimental images (Sr = red,Fe = green, O =blue). The originalmicrographs are shown inthe Supporting Information Figures S1 and S2.

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provides an alternative possible explanation for the lowoxygen-ion transport through the Sr4Fe6O12þδ structure.It is assumed that the migration of the ion charge

carriers in Sr4Fe6O12þδ, both by oxygen vacancies andinterstitials, occurs along the Fe2O2þδ layers.

24�26 Bothmigration of oxygen vacancies and interstitial diffusionare believed to contribute to the total ionic conductiv-ity of this material. At low partial oxygen pressures[p(O2) < 105 Pa] the ionic conductivity is dominated bythe migration of oxygen vacancies, whereas at higherp(O2) > 105 Pa the oxygen transport occurs throughoxygen interstitials.30 The atomic arrangement withinthe FeO1þδ/2 sublayers provides the native paths forthe migration of the oxygen interstitials: it consists of ajump of the oxygen atom along the l channel into aneighboring interstitial position, coordinated by twoFe atoms (marked with green arrows in Figure 6). Inthe uniform FeO1þδ/2 sublayers infinite chains of inter-stitial positions are available and the migration canoccur at a relatively low energy cost. However, intro-ducing a mixed configuration with a fragment shiftedby c/2 effectively breaks this diffusion pathway. Thedistance between the O atom and the next availableinterstitial position becomes ∼

√2 times larger (red

arrows in Figure 6) and a bottleneck of two closelypositioned oxygen atoms additionally prevents theoxygen transfer. This qualitative picture, essentiallybased on geometrical considerations, implies that thedisorder in the Fe2O2þδ layers could deteriorate theoxide ion transport. Careful control of the synthesisconditions might diminish the degree of this disorderand its potential negative effect.Information about the bonding and coordination of

the Fe atoms was obtained by analyzing the electronenergy-loss near-edge structure (ELNES) of the differ-ent layers in the SFO unit cell. Figure 7 shows the OK-edge and Fe L2,3-edge EELS spectra acquired fromthe two different building blocks present in SFO, thatis, the perovskite-type layers and the double Fe2O2þδ

layers. Smart Acquisition linescans36,37 were run alongthe stacking direction crossing several FeO2 andFe2O2þδ layers (for details see the Methods section).Thus, the experimental results reflect the integral of allpossible oxygen and iron environments in each layer.The EELS spectra were background subtracted for eachcore-loss edge and no PCA processing was applied.

The ELNES of the O K-edge spectra can be dividedinto two regions: the prepeak region from ∼528to ∼534 eV and the postedge region from ∼534 to∼548 eV. The prepeak region is characterized by thepresence of a well-defined peak at 531.2 eV (labeled A)in both spectra and by an additional low-energy peakat 530.4 eV (labeled A0) in the FeO2 spectrum; see theinset in Figure 7. The postedge region of the FeO2

layers shows three well-separated features at 536.5 eV(labeled B), 541 eV (labeled C), and 544 eV (labeled D),while the Fe2O2þδ layer spectrum is characterized by aB peak maximum followed by a much lower D peak.As EELS is comparable to X-ray absorption spectros-

copy (XAS), in that X-rays produce similar ionizationprocesses to those inducedby the electron beam in theTEM, we compared our experimental spectra with XASand EELS data from SrFeO3,

38 and other iron oxides.39�41

We found that peak A stems from transitions tohybridized O 2p�Fe 3d orbitals, while the postedgeregion of the spectra is attributed to hybridizationbetween the O 2p and Fe 4sp states. The observeddifferences in the postedge fine structure of the FeO2

and Fe2O2þδ spectra reveal the different site coordina-tion of the Fe atoms which is determined by thenearest oxygen neighbors. Besides, the splitting ofthe prepeak region of the perovskite-type FeO2 layerspectrum by 0.8 eV (see the inset in Figure 7) indicatesthat the 3d orbital states further split into t2g and eglevels by octahedral ligand-field splitting. A strik-ing difference between the spectra of the FeO2 andFe2O2þδ layers is also found on the relative intensity ofthe A and B peaks. While the intensity of these peaks isvery similar for the Fe2O2þδ layer spectrum (red), the Bpeak intensity is almost twice the intensity of the Apeak for the FeO2-layer spectrum (black). Previousstudies showed that the prepeak can be used toestimate the density of hole states.42,43 Therefore, thehigher prepeak in the Fe2O2þδ-layer spectrum indi-cates a higher concentration of mobile holes in thedouble-layers. This is in good agreement with pre-viously reported electron�hole conductivity measure-ments in epitaxial SFO film which showed that the

Figure 7. Oxygen K-edge (left) and iron L-edge (right)spectra of the perovkite-type FeO2 layers, and the doubleFe2O2þδ layers. The FeO2 spectra are averaged from 30spectra while the Fe2O2þδ spectra are averaged from 45spectra. Inset: enlarged view of the prepeak region A.

Figure 6. Possible oxygen migration pathways along theFeO1þδ/2 sublayers. Green arrows indicate the oxygen jumps tothe closest interstitial position in the ordered layer; red arrowsindicate the pathway in the layer with mixed configuration.

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conductivity of the films can be tailored by changingthe oxygen content in the double layers.44

In comparison, changes in the fine structure of the FeL2,3-edgeare farmore subtle (Figure 7). The spectra consistof two “white line” peaks which are separated by∼13 eVdue to the spin�orbit splitting of the Fe 2p core hole.Besides, the FeO2-layer spectrum (black) shows a shoulder(at 710 eV) on the left side of the L3 maximum due to thet2g�eg crystal field splitting of the octahedral coordina-tion. In the data acquired from the nonperovskite Fe2O2þδ

layers (red) on the other hand, the shoulder is hardlyvisible in the L3 peak,whichmakes it appear narrower. Thelack of fine structures in the Fe L2,3-edge of the doublelayersmightbecausedby theoverlappingbetween the3dstates arising from the various distorted pentacoordinatedgeometries of the iron atoms. No significant relativeintensity changes of the L3 and L2 white lines are obser-vable in the experiments confirming that the oxidationstate of Fe is equivalent in both types of layers.

CONCLUSIONS

By using aberration-corrected HAADF- and ADF-STEM imaging combined with EELS at atomic resolu-tion, we have demonstrated that the double Fe2O2þδ

layers in the SFO structure may adopt different atomicarrangements. Our results provide direct experimentalevidence for the simultaneous presence of areas exhibit-ingdifferent atomic reconstructions of the Fe2O2þδ layers.In the most favorable arrangement a “disordered”

structure is ascertained. The spatial arrangement of the FeandO atoms in the “disordered” layers is derived from theincommensuratelymodulatedSr4Fe6O12þδ structure

5 as alocally occurring shift of the Fe�O chains in the Fe2O2þδ

layers. “Ordered” atomic reconstructions of the Fe2O2þδ

layers are alsoobservedover relatively large local areas. Bycomparing our high-resolution STEM experimentaldatawithmultislice image calculations, we corroboratethe proposed structural models. We have also shownthat the different coordination environment for the Featoms in the perovskite-type layers and the doubleFe2O2þδ layers results in clear differences in the near-edge fine structure of the O K-edge. Thus, the observeddifferences in the relative intensity of the O K-edgeprepeak indicate a higher mobile-hole concentrationin the double layers. Therefore the presence of APBs,which laterally transform Fe2O2þδ layers into FeO2

perovskite-type layers and vice versa, is likely to havea deleterious effect on both the electron�hole and theoxide�ion conduction of the films. Even a more strik-ing negative effect on the ionic conductivity can beexpected from themixed configuration of the Fe2O2þδ

layers which effectively breaks the pathways for theinterstitial diffusion. The results we present clarify theexact atomic stacking structure of the Sr4Fe6O12þδ

phases and evidence STEM-EELS is a valuable comple-mentary technique to X-ray diffraction when determi-nation of a crystal structure is hinderedby the presenceof disorder-related diffuse scattering.

METHODSSample Preparation. The epitaxial Sr4Fe6O12þδ (SFO) thin film

was grown on a polished (001) oriented SrTiO3 substrate bypulsed laser deposition using a Nd:YAG laser with tripledfrequency (355 nm wavelength), 9 ns pulse length, 10 Hzrepetition rate, and 2�3 J/cm2 energy density per pulse. Thetarget for the ablation experiment was prepared by a solid-state reaction from a stoichiometric mixture of SrCO3 andFe2O3. The resultant powder was milled, uniaxially pressedinto pellets, and finally sintered in air at 1150 �C. X-raydiffraction measurements confirmed that the target wasthe single phase Sr4Fe6O13(δ, as determined by Yoshiasaet al.1 During deposition, the substrate temperature waskept at 750 �C and the oxygen pressure in the chamber wassettled at 3 � 10�2 mbar. After the sample was deposited,it was cooled at 10 �C/min under the same atmosphere.Details of the X-ray diffraction studies and scanning elec-tron microscopy are found in Pardo et al.19 Samples forTEM were prepared in cross-section geometry; the samples werecut parallel to the (110) planes of the SrTiO3 substrate andmechanically groundwith a tripod polisher to a thickness of about20 μm, followed by final liquid nitrogen-cooled ion milling undergrazing incidence until electron transparency.

Instrumentation. The EELS data in this study was obtained at100 kV on a Nion UltraSTEM 100 microscope equipped with aC5 Nion quadrupole�octupole-type probe aberration correctorand an ultrahigh vacuum Gatan Enfina spectrometer. Theconvergence and collection semiangles were set to 30.5 and33 mrad, respectively. For these values, the energy resolutionmeasured as the full width at half-maximum of the zero-losspeak is ∼0.35 eV. Yet, in order to cover a large energy range inEELS, the dispersion was set to 1 eV/channel for the spectrum

images in Figures 2 and 3. No drift correction during acquisitionwas used in this work. Energy windows of 50 eV were used toform the EELS images of the Sr-M4,5, Fe-L2,3, and O-K edges. Toresolve the near-edge fine structure of the core-loss edges ofFigure 7 the dispersion was set to 0.2 eV/channel. In typicaloperating conditions for the experiments described in thispaper, the microscope provides an estimated spatial resolu-tion of 0.85 Å. The annular semidetection range of the annulardark-field detector was set to collect electrons scatteredbetween 95 and 195mrad for the HAADF images and between41 and 95 mrad for the ADF images (when used at the sametime as the HAADF detector). HAADF-STEM in Figure 5 wascarried out using the aberration-corrected TEAM 0.5 micro-scope operated at 300 kV. The TEAM 0.5 microscope45 is amodified FEI Titan 80�300 equipped with a Schottky-typehigh-brightness field-emission electron source and an improvedhexapole-type illumination spherical aberration (CS) corrector.

46 Aprobe convergence angle of ∼19 mrad was set yielding a calcu-lated probe size of about 65 pm. The annular semidetection rangeof the annular dark-field detector was calibrated at 80�300 mrad.

Image Simulations. High-resolution STEM image simulationswere carried out using the xHREM simulation package.

Conflict of Interest: The authors declare no competingfinancial interest.

Acknowledgment. The authors would like to thank J. Santiso,J. A. Pardo, and C. Solís for providing the sample. We acknowl-edge the financial support of the Swiss COST office under theSBF project number C10.0089. The SuperSTEM Laboratory issupported by the U.K. Engineering and Physical Sciences Re-search Council (EPSRC). This research used resources of theNational Center for ElectronMicroscopy (LBNL) which is supported

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by the Office of Science, Office of Basic Energy Sciences of the U.S.Department of Energy under Contract No. DE-AC02-05CH11231.

Supporting Information Available: Additional figures show-ing experimental HAADF-STEM images of the disordered andordered Sr4Fe6O12þδ structure. This material is available free ofcharge via the Internet at http://pubs.acs.org.

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of Sr4Fe6O13, A New Perovskite-Derivative in the Sr�Fe�OSystem. Mater. Res. Bull. 1986, 21, 175–181.

2. Rossell, M. D.; Abakumov, A. M.; Van Tendeloo, G.; Pardo, J. A.;Santiso, J. StructureandMicrostructureofEpitaxial Sr4Fe6O12þδ

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