Atomic Resolution Coordination Mapping in Ca2FeCoO5Brownmillerite by Spatially Resolved Electron Energy-LossSpectroscopyStuart Turner,*,†,‡ Johan Verbeeck,‡ Farshid Ramezanipour,§,⊥ John E. Greedan,§,⊥

Gustaaf Van Tendeloo,‡ and Gianluigi A. Botton†,⊥,∥

†Canadian Centre for Electron Microscopy, §Department of Chemistry, ⊥Brockhouse Institute for Materials Research,and ∥Department of Materials Science and Engineering, McMaster University, Hamilton, Canada, L8S 4M1‡EMAT, University of Antwerp, B-2020 Antwerp, Belgium

ABSTRACT: Using a combination of high-angle annular dark field scanningtransmission electron microscopy and atomically resolved electron energy-lossspectroscopy at high energy resolution in an aberration-corrected electronmicroscope, we demonstrate the capability of coordination mapping in complexoxides. Brownmillerite compound Ca2FeCoO5, consisting of repetitiveoctahedral and tetrahedral coordination layers with Fe and Co in a fixed 3+valency, is selected to demonstrate the principle of atomic resolutioncoordination mapping. Analysis of the Co-L2,3 and the Fe-L2,3 edges showssmall variations in the fine structure that can be specifically attributed to Co/Fein tetrahedral or in octahedral coordination. Using internal reference spectra,we show that the coordination of the Fe and Co atoms in the compound can bemapped at atomic resolution.

KEYWORDS: oxygen coordination, valency, electron energy-loss spectroscopy, EELS, energy-loss near edge structure, ELNES,Fe-L2,3 edge, Co-L2,3 edge, brownmillerite

■ INTRODUCTIONIn complex transition metal oxides, the coordination andoxidation state of the transition metals is of fundamentalimportance. Both are known to be related to the structural,electronic, magnetic, catalytic, and ionic transport properties of(complex) oxides.1−3 The coordination and valency of transitionmetals in complex oxides can be determined by bulk methodslike Mossbauer spectroscopy and in some conditions bydiffraction techniques. However, local changes in coordinationcan be the key factors in materials applications.4,5 For example,coordination changes at surfaces are of vital importance formany catalytic processes while changes at defects like twinboundaries, grain boundaries, and interfaces can greatly affectthe electronic, optical, and transport properties of bulkmaterials and thin films. The spatial resolution needed for thedetermination of coordination for these important problems ishowever not offered by classical bulk techniques like Mossbaueror diffraction methods.Atomic resolution elemental mapping has become feasible

over the past years by means of spatially resolved electronenergy-loss spectroscopy (EELS) and, more recently, energydispersive X-ray spectroscopy in a scanning transmission electronmicroscope (STEM-EELS and STEM-EDX)6−11 and has beenused for local structure refinement.12 Even though STEM-EDXappears to allow more straightforward data acquisition andinterpretation as compared to EELS, fine structure informationlike bonding or coordination cannot be obtained from X-ray

spectra in the electron microscope with current spectrometers.13

Combining the atomic resolution capabilities of a STEM withthe exquisite spectroscopic sensitivity to bonding and valenceinformation from EELS fine structure is therefore a highlyattractive prospect that would draw significant attention in thefield of solid state chemistry, physics, and materials science. Fortransition metals, the shape of the L2;3 edges,

14 chemical shifts,15

and the L3/L2 ratio16 have all been used as a sensitive fingerprint

for valence.17 A large correlation between the energy-loss near-edge structure (ELNES) and valence in transition metal oxideshas been confirmed, even at atomic resolution, by experiments inliterature. Using this correlation between ELNES and valency,oxidation state mapping at atomic resolution in CeOx nano-particles and bulk Mn3O4 was recently demonstrated in some ofour work.18,19 The main issues hindering atomic resolutionvalency and bonding measurements remain poor EELS signal-to-noise ratios because of the need for simultaneous high spatial andenergy resolution of the instrument. This problem is exacerbatedfurther as, in general, changes in the fine structure of the L2,3edge of transition metals because of bonding or coordinationare far more subtle than changes related to valency.20 Studieswhere bonding and coordination have been mapped at atomicresolution are therefore rare, and usually the change in bonding

Received: February 27, 2012Revised: April 23, 2012Published: April 25, 2012

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coincides with a change in valency which also effects the EELSfine structure.8,21−23

In this work, we set out to map the coordination of Fe andCo in a Ca2FeCoO5 brownmillerite sample with constant Feand Co valency but varying coordination.24 Ca2FeCoO5 is abrownmillerite compound with an orthorhombic crystalstructure (Pbcm) with lattice parameters a = 5.3652(3) Å,b = 11.0995(5) Å, c = 14.7982(7) Å, and is structurally similarto the well-documented Ca2Fe2O5.

25 However, in Ca2FeCoO5one unit cell axis is doubled compared to Ca2Fe2O5 and otherregular brownmillerites. Also, unlike typical brownmillerites,Ca2FeCoO5 has two octahedral and two tetrahedral sites thatare all crystallographically distinct, making the structure evenmore complex. The Fe and Co cations are stacked along thec-direction in alternating octahedral and tetrahedral stackinglayers in a constant 3+ valency (confirmed by our EELS data).The fact that Ca2FeCoO5 is a mixed coordination/single-valency compound makes it an ideal candidate to demonstratepure coordination mapping at atomic resolution; EELS finestructure variations over the tetrahedral and octahedral layerscan unequivocally be assigned to coordination changes, and notvalency variation. The presence of two different cation species(Fe, Co) in the sample can provide insight into the effect ofcoordination on the L2,3 edge ELNES of both Co and Fe atatomic resolution.

■ EXPERIMENTAL SECTIONFor experimental investigation, samples were crushed and deposited dryonto a lacey carbon grid. The samples were investigated at theCanadian Centre for Electron Microscopy (McMaster University,Canada) using a FEI Titan “cubed” microscope equipped with a probecorrector, operated at 200 kV for Figures 1, 3, 4, and 5 and a similar

instrument at the EMAT laboratory (University of Antwerp, Belgium)operated at its optimal acceleration voltage of 120 kV, for the data inFigure 2. For the data in Figures 1, 3, 4, and 5 the microscope was

operated in STEM mode using a convergence semiangle α of18.5 mrad, yielding a probe size of ∼1.0 Å. The ADF inner collectionsemiangle was 60 mrad for imaging, and the spectrometer acceptancesemiangle β was 70 mrad. The beam current was kept close to 40 pAfor imaging and between 80 and 140 pA for spectroscopy. Suitableacquisition times (in the range of 10 ms/pixel to 20 ms/pixel) werechosen to avoid beam damage and provide the best combinationof signal-to-noise ratio, spatial and energy resolution. Where needed,the monochromator was excited to provide an energy resolution of250 meV.

The elemental maps in Figure 3 were generated in the case of Caand O by mapping the intensity under the background subtractededges. For Fe and Co, to improve the S/N ratio, the maps weregenerated through the so-called jump-ratio method. A post edge imagewas divided by a pre-edge image, yielding a qualitative result.13 Whenfiltered, a 7×7 light low-pass filter was used in the Digital Micrographsoftware package. The data in Figure 5 was fitted to an octahedral anda tetrahedral internal reference component, being the summed finestructure spectrum of the top octahedral and the tetrahedral layer(which corresponds well to literature fine structure for 4 and 6coordinated cations) using the EELS model software according to amethod described in previous publications.18,19

For the data in Figure 2 the microscope was operated in diffractionmode with excited monochromator providing an energy resolutionof 250 meV. The convergence semiangle α was ∼1 mrad, and theacceptance semiangle β was ∼2 mrad.

Image simulations were performed with the STEMSIM softwarepackage, a MATLAB based image simulation program capable ofhandling the complex interplay between elastic and inelastic scatteringin the double channeling approximation. Elastic scattering wassimulated with a Bloch wave approach (max wavevector 2.5 Å−1),in 10 thickness steps to a total thickness of 20 nm using a unit cellsampling: 72×116 pixels per unit cell and a supercell of 3×1 unit cells.Source size broadening was taken into account using a Gaussian with0.7 Å fwhm. Proper agreement between simulations and experimentswas only obtained using a random distribution of Fe and Co over theoctahedral and tetrahedral sites, as discussed below. Inelastic scattering

Figure 1. HAADF-STEM images of Ca2FeCoO5 along the [100] (a)and [010] (b) zone axes orientation. The images show the repetitionof octahedral and tetrahedral layers along the c direction. An HAADF-STEM image simulation is inserted to the [100] zone axis image.Projected structural models for the [100] and [010] zone axisorientations are displayed in (c) and (d) respectively (Red = Co,Green = Fe, Brown = Ca, Pink = O). The unit cell is indicated by theblack rectangles.

Figure 2. High resolution EELS data of Ca2FeCoO5. (a) EELSspectrum of Ca2FeCoO5 showing the Ca-L2,3 edge at 352 eV, the O−Kedge at 529 eV, the Fe-L2,3 edge at 708 eV, and the Co-L2,3 edge at 779 eVsimultaneously; (b) the Fe-L2,3 edge in high resolution showing a shoulderto the L3 edge at 709 eV, typical for Fe in a brownmillerite-type structure;(c) the Co-L2,3 edge showing a far less-pronounced shoulder to the L3edge, and a small postpeak shoulder at ∼782 eV.

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was simulated using a relativistically corrected Dipole approximationwith a hard Bethe ridge cutoff.

■ RESULTS AND DISCUSSIONFigure 1 shows two high-angle annular dark field (HAADF)STEM images of the Ca2FeCoO5 sample along the [100] andthe [010] zone axis orientation. In both images, the c-axis isvertical, as indicated by the arrow in Figure 1b. The repetitionof octahedral (example indicated by the red arrow) andtetrahedral (example indicated by the blue arrow) Co/Felayers is apparent from the different structure of the layers inthe [100] zone axis orientation and the contrast changes in the[010] zone axis orientation. Projected structural models for the[100] and [010] zone axis orientations are displayed in (c) and

(d) respectively and correspond well to the HAADF-STEMimages. A Bloch wave HAADF image simulation for a 20 nmthick crystal is inserted into Figure 1a, and agrees remarkablywell with the experimental image.Bulk Ca2FeCoO5 EELS data recorded over areas of several

hundred nm are plotted in Figure 2. The spectrum in Figure 2ashows the Ca-L2,3, the O−K, the Fe-L2,3, and the Co-L2,3 edgesimultaneously. The Fe-L2,3 and Co-L2,3 are plotted in moredetail (at a lower energy dispersion) in Figure 2b and 2c. Toacquire these data, a monochromator was used to provide anenergy resolution of approximately 250 meV. At a first glance,the fine structure of the Fe-L2,3 and Co-L2,3 edges is similar.This is to be expected, as both transition metals occupy similarpositions within the Ca2FeCoO5 structure. Comparing the

Figure 3. Atomic resolution elemental maps for Ca, O, Fe, and Co along the [100] zone axis. (a) Raw Data. (b) Smoothed data. (c) Doublechannelling, inelastic Bloch wave image simulations for Ca, O, Fe and Co. The maps clearly show that Co and Fe are intermixed in each Co/Fecolumn, and demonstrate the repetition of the octahedral and tetrahedral coordination layers. (d) Averaged EELS spectra over the entire regionshowing the Fe/Co fine structure in the octahedral and tetrahedral layers at an energy resolution of ∼1 eV.

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Fe-L2,3 edge to reference spectra in the literature reveals thatthe fine structure is identical to the literature fine structure forFe3+ in Ca2Fe2O5.

20 Moreover, literature data for Fe3+ inoctahedral and in tetrahedral coordination show that theshoulder of the L3 edge at ∼709 eV can be used as a marker forthe coordination of Fe; in compounds with octahedral Fe3+,the shoulder is well-pronounced because of a t2g-eg splitting ofthe L3 peak. In compounds with solely tetrahedral Fe3+, thisshoulder is hardly visible.20 In Ca2Fe2O5, the shoulder is lesspronounced as Fe is present in both an octahedral and atetrahedral coordination. A similar shoulder is visible near theCo-L3 onset around 779 eV, although it is much less prominentthan for the Fe-L3 edge. This is in accordance with calculationsand experiments in literature showing that both octahedral andtetrahedral Co3+ (especially in the expected high spin state ofCo3+ in this compound) only display a small L3 prepeak.

24,26−28

To acquire the spectroscopic data needed for coordinationmapping, the so-called spectrum imaging technique wasadopted.13 In this technique, the electron probe (in our setupsthe probe has a diameter of ∼1.0 Å) is scanned over the sampleand an EELS spectrum is acquired in each point together withan annular dark-field signal as image reference. A first set ofatomic resolution elemental maps, acquired at an energyresolution of ∼1 eV with the sample oriented along the [100]zone axis, is displayed in Figure 3. Elemental maps for Ca, O,Fe, and Co (top) are displayed together with low-pass filtereddata (middle) and image simulations (bottom). Although some

instabilities are present because of sample charging, the mapsclearly show an unexpected intermixing of Fe and Coat both the octahedral and the tetrahedral sites, which aresandwiched in between layers of CaO. Indeed, a properagreement between simulations and experiments could only beobtained using a random distribution of Fe and Co over theoctahedral and tetrahedral sites. This appears to be at least inpartial disagreement with a previously published neutrondiffraction study24 where site ordering was found for Fe/Coon the tetrahedral sites, and warrants further investigation inthe future. The contrast in the oxygen map is influenced by thepresence of the nearby “heavy” Fe and Co columns, whichcommonly occurs for atomic resolution oxygen maps and isreproduced well in the simulation.29 Averaged Fe-L2,3 andCo-L2,3 edge spectra at ∼1 eV energy resolution from theoctahedral and tetrahedral layers are plotted below the elementalmaps and show minimal discernible differences. This is to beexpected, as most large fine structure changes like edge-onsetenergy and L2,3 white lines ratio are related to valency, notcoordination. It is clear that a better energy resolution is neededto discern the small variations in the EELS fine structure that arerelated to coordination.Figure 4 shows atomic resolution data acquired with an

excited monochromator, providing an energy resolution of250 meV. The overview image (Figure 4b), taken simulta-neously with the EELS data, shows that working with an excitedmonochromator leads to only a limited loss in spatial resolution.

Figure 4. EELS data acquired at high energy resolution and atomic spatial resolution. (a) Fe-L2,3 edge summed over the first octahedral (red) andtetrahedral (blue) layer. (b) Survey image indicating the area used for octahedral (red) and tetrahedral (blue) data summation (3 pixel width).(c) Co-L2,3 edge summed over the first octahedral (red) and tetrahedral layer (blue).

Figure 5. Octahedral and tetrahedral EELS maps generated by fitting the average components to the atomic resolution data. The octahedralcomponent (b) and tetrahedral component (c) peak at the corresponding Fe/Co intermixed layers, as clearly evidenced by the intensity profiles overthe full image width.

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In Figure 4a and 4c data from a single layer of Co/Fe inoctahedral and tetrahedral coordination (both indicated byarrows in the overview image, data averaged over 3 verticalpixels) is displayed. Contrary to the data at 1 eV energyresolution, clear differences in the Fe-L2,3 and the Co-L2,3 finestructure can be resolved by simple inspection. In the octahedrallayer (red) the shoulder to the L3 edge for Fe, and to a lesserextent also for Co, is present in good agreement with the datafrom literature.20 In the case of octahedral Co, a small shoulderis also present behind the L3 maximum (at ∼782 eV), inagreement with literature.27 The data acquired from a singletetrahedral layer on the other hand (blue) shows a far lessdistinct shoulder in the case of Fe. Concomitant with thischange, the L3 signature for Fe and Co-L3 in the tetrahedralcoordination peaks higher and slightly narrower than the L3 forthe octahedral sites (after alignment of the data to the integralunder the L2,3 edge). This is at least in part because of the whiteline ratio; as the L3/L2 white lines is known to remain constantfor constant valency, the intensity “lost” under the L3 shoulder iscompensated by a higher and narrower L3 peak. For tetrahedralCo, the small shoulder at 782 eV is also subdued.26 In all, clearexquisite differences in fine structure are visible between theoctahedral and the tetrahedral spectra for both Fe and Co atconstant valence. These differences can be used to map thecoordination of the Fe and Co cations in the structure usingthe averaged EELS signals from Figure 4a and 4c as internalreferences.In Figure 5, the internal reference data for Fe and Co in the

two distinct coordinations are fitted to the entire acquiredEELS datacube, following the procedure from earlier work onatomic resolution valency mapping.18,19,30,31 Fitting the internal“octahedral” and “tetrahedral” reference data to each acquiredspectrum allows coordination maps with the weight of eachcomponent to be generated. Prior to fitting, the data cube wasrebinned in X and Y direction by a factor of 2 to improve theS/N ratio in the individual spectra. The results of the fit aredisplayed in Figure 5b,c. The octahedral and tetrahedralcomponent maps of the transition metals clearly follow theatomic contrast of the separate coordination layers in the ADFsurvey image. Moreover, each component map peaks at thecorrect coordination layer. These results clearly demonstratethat coordination can be mapped in these structures down toatomic resolution. Because of the small variations in the finestructure of the Fe L2,3 and Co L2,3 edges used for mapping,lateral resolution is not clearly obtained.The results obtained in Figures 3 and 5 clearly indicate that

using the fine structure of EELS data acquired at atomicresolution with a high energy resolution, advanced electronmicroscopy can surpass elemental mapping, and provide atomicresolution information on, for example, oxygen coordination.The data presented in this paper also suggests that mapping ofvalence and coordination independently, along with theelemental map, can be done in the future. As well, this studyshows that that Fe3+ and Co3+ are distributed in a roughlyrandom manner over both sets of the octahedral and tetrahedralsites in Ca2FeCoO5. This is at least in partial disagreementwith the results of a published refinement of neutron diffractiondata where Fe/Co site order was found on the tetrahedral sites.The origin of this discrepancy is unclear at the moment, but itmight be related to the inherent difference between the twomethods: neutron diffraction being a bulk probe as opposed tothe electron microscopy which is a local method. A new neutron

data set has been obtained and refinements are underway toreconcile this discrepancy.

■ CONCLUSIONWe have demonstrated the possibility of coordination mappingby atomically resolved high resolution EELS in an aberration-corrected electron microscope, through the example of singlevalency/mixed coordination cation compound Ca2FeCoO5.Detailed analysis of the fine structure of the Co-L2,3 and theFe-L2,3 edges showed small variations in the fine structure thatcould be attributed to Co/Fe in tetrahedral or Co/Fe inoctahedral coordination and that were in excellent agreementwith literature spectra for octahedral and tetrahedral coordi-nated Fe3+/Co3+. Using the spectra from octahedral/tetrahedral Fe/Co layers as an internal reference (and in thefuture possibly theoretical spectra to identify unknownmaterials), the coordination of the Fe and Co atoms in thecompound could be mapped at atomic resolution. Theseexperiments open the gate for coordination determination atsurfaces, defects, and grains boundaries in a plethora of complexoxide materials that have been out of reach in the past becauseof the lack of suitable analysis techniques for the study ofvalence and coordination as input for solving their structure.

■ AUTHOR INFORMATIONCorresponding Author*E-mail: stuart.turner@ua.ac.be.NotesThe authors declare no competing financial interest.

■ ACKNOWLEDGMENTSA.M. Abakumov is thanked for fruitful discussions. S.T. gratefullyacknowledges the Fund for Scientific Research Flanders (FWO).J.E.G. and G.A.B. acknowledge the support of the NSERC ofCanada through Discovery Grants. The Canadian Centre forElectron Microscopy is a National Facility supported by NSERCand McMaster University and was funded by the CanadaFoundation for Innovation and the Ontario Government. Part ofthis work was supported by funding from the European ResearchCouncil under the FP7, ERC Grant N 246791 COUNTATOMSand ERC Starting Grant N 278510 VORTEX. The EMATmicroscope is partially funded by the Hercules fund of the FlemishGovernment.

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