IOP PUBLISHING NANOTECHNOLOGY

Nanotechnology 19 (2008) 395601 (7pp) doi:10.1088/0957-4484/19/39/395601

Nanostructural control in solution-derivedepitaxial Ce1−xGdxO2−y films

M Coll1, J Gazquez1, F Sandiumenge1,4, T Puig1, X Obradors1,J P Espinos2 and R Huhne3

1 Institut de Ciencia de Materials de Barcelona, CSIC Campus de la UAB, E-08193 Bellaterra,Catalonia, Spain2 Instituto de Ciencia de Materiales de Sevilla, CSIC Avenida Americo Vespucio s/n Isla de laCartuja, E-41092 Sevilla, Spain3 IFW Dresden, Leibniz Institute for Solid State and Materials Research, Dresden,Helmholtzstraße 20, D-01069, Germany

E-mail: felip@icmab.cat

Received 6 June 2008Published 11 August 2008Online at stacks.iop.org/Nano/19/395601

AbstractA novel mechanism based on aliovalent doping, allowing fine tuning of the nanostructure andsurface topography of solution-derived ceria films, is reported. While under reducingatmospheric conditions, non-doped ceria films are inherently polycrystalline due to aninterstitial amorphous Ce2C3 phase that inhibits grain growth, a high quality epitaxial film canbe achieved simply by doping with Gd3+ cations. Gd3+ ↔ Ce4+ substitutions within the latticeare accompanied by charge-compensating oxygen vacancies throughout the volume of thecrystallites acting as an efficient vehicle to reduce the barrier for grain boundary motion causedby interstitial Ce2C3. In this way, the original nanostructure is self-purified by pushing theamorphous Ce2C3 phase towards the free surface of the film. Once a full epitaxial cube-on-cubeoriented ceria film is obtained, its surface morphology is dictated by the interplay betweenfaceting on low energy {110} and/or {111} pyramidal planes and truncation of those pyramidsby (001) ones. The development of the latter requires the suppression of their polar characterwhich is thought to be achieved by charge compensation between the dopand and oxygen along〈100〉 directions.

(Some figures in this article are in colour only in the electronic version)

1. Introduction

Ceria-based oxides are widely investigated because of theirmultiple applications as catalysts [1, 2], electrolyte material ofsolid oxide fuel cells [3, 4], solar-cell materials [5] or as bufferlayers in coated conductor architectures [6–9]. Such diverseapplications require different nano- and mesostructures,ranging from nanocrystalline high surface area materials todense epitaxial films. For instance, it has been reportedthat the ionic conductivity of Ce1−x GdxO2−y (CGO) dependson grain size [10]. On the other hand, the use of CGOas a buffer layer requires a high epitaxial quality andlow surface roughness. In this context, chemical solutiondeposition (CSD) [11, 12] emerges as a scalable and cost-

4 Author to whom any correspondence should be addressed.

competitive deposition technique, with a high potential forthe creation of specific nanostructures [13, 14]. Whatmakes CSD unique is that nucleation occurs under very highundercoolings [15, 11, 12]. Under such a large thermodynamicdriving force, homogeneous and heterogeneous nucleationevents occur with similar probability. Typically, grainsheterogeneously nucleated on the substrate exhibit a lowenergy epitaxial orientation, which upon thermal annealingexperience secondary grain growth at the expense of thedisordered precursor [16]. In this stage, the large interfacialenergy difference between the nanostructured film and theunderlying epitaxial grains provides the driving force forrecrystallization into an epitaxial structure. It appears thatthe microstructural trajectory from the polycrystalline tothe epitaxial forms of the film is fully governed by grain

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Nanotechnology 19 (2008) 395601 M Coll et al

boundary (GB) motion and therefore it is extremely sensitiveto impurity interactions. Interestingly, this particular featureof CSD growth offers a unique opportunity for tailoringspecific nanostructures by properly tuning the microstructuraltrajectory from a nanocrystalline metastable state to a fullyepitaxial one.

We have recently shown that, after the pyrolysis step,residual C-impurities from the metal–organic precursor remaintrapped at grain boundaries and the interstitial space withinthe ceria nanocrystalline film. It is well known that grainboundary motion can either be blocked or enhanced by variousmechanisms, including doping or the formation of grainboundary layers of segregated phases [17, 18]. In the presentcase, the drag force exerted by interstitial carbon inhibits graingrowth, thereby precluding the development of an epitaxialfilm. In that previous report, C-impurities were eliminated byoxidation and then the film underwent a rapid transformationinto an epitaxial structure. However, the use of oxidizingatmospheres may be elusive for processes involving oxidation-sensitive materials. This circumstance constitutes a majordrawback for coated conductor development, where metallicsubstrates are used [8].

In the present work we demonstrate a new mechanism forthe elimination of grain boundary and interstitial C-impuritiesthat operates even under reducing atmospheric conditions.The methodology uses aliovalent doping by Gd3+ cations,which is well known to depress the activation energy for grainboundary motion [19], and therefore appears as a suitablecandidate to promote epitaxial growth. It is shown that, in thequenched nanocrystalline state, C-impurities are in the form ofamorphous Ce2C3 filling interstitial space. The enhanced grainboundary mobility is sufficient to unblock grain growth againstthe drag force exerted by the amorphous Ce2C3 phase, whichis pushed by the growth front towards the film surface where itis eliminated through the flowing Ar/H2 gas.

2. Experimental details

Precursor solutions were prepared by dissolution of stoi-chiometric amounts of gadolinium acetylacetonate (Gd(CH3

COCHCOCH3)3, 99.9% from Aldrich) in 1 ml of propionicacid (Aldrich) and 1 ml of anhydrous isopropanol (Aldrich).The solution was stirred and heated up to 40 ◦C for 20 min,followed by the incorporation of cerium (III) acetylacetonatehydrate (Ce(CH3COCHCOCH3)3xH2O, 99.9%, Alfa Aesar),and then stirred for another 20 min at 60 ◦C obtaining a yellowtransparent solution. The concentration of the solution was ad-justed to 0.25 M (total metal ion concentration). Samples withnominal composition Ce1−x Gdx Oy (x = 0, 0.01, 0.1, 0.2, 0.4)were prepared. Optimum results in terms of surface morphol-ogy and epitaxy were obtained for a 10 at% Gd concentration.The final precursor solution was kept in sealed vials for severalmonths. The solution lifetime was controlled by controlling itsviscosity at room temperature using a rheometer (3 mPa s).

15 μl of the precursor solution was deposited onto5 mm × 5 mm (001)-cut Y-stabilized ZrO2 (YSZ) single-crystal substrates, previously cleaned for 10 min in acetone and10 min in methanol using an ultrasonic bath. Deposition was

performed with a rotation speed of 6000 rpm, rotation time of2 min and an acceleration of 3000 rpm s−1. After spin-coating,samples were heated in a tubular furnace at 1500 ◦C h−1 upto 700–900 ◦C, then held for 8 h to allow completion of theepitaxial transformation of the precursor, and finally cooleddown to room temperature. Film growth was carried out bothunder reducing (Ar/5%H2) and oxidizing (O2) atmospheres.These deposition and growth conditions yielded films with afinal thickness of ∼30 nm.

X-ray photoelectron spectroscopy (XPS) measurementswere carried out on an ESCALAB 210 spectrometer at apressure of 10−10 Torr (1 Torr = 133.322 Pa). Data for theCe(3d) and C(1s) core levels were acquired using an Mg Kα x-ray source (photon energy 1253.6 eV). All spectra were takenwith the hemispherical electron energy analyser working inthe constant-pass energy mode, at �E = 50 eV and using astep size of 100 meV. The Ce(3d) spectrum of a film grownunder a pure oxygen atmosphere was selected as a pure 100%Ce4+ ceria film reference, for Ce3+ quantification. Bindingenergies were normalized using the satellite peak at 917 eVof the reference spectrum, which is exclusively contributed byCe4+.

Structural properties and surface morphology of thefilms were investigated by x-ray diffraction (XRD), reflectionhigh energy electron diffraction (RHEED) and atomicforce microscopy (AFM). XRD texture measurements wereperformed in a Philips mrd system using Cu Kα radiationand parallel beam optics. RHEED patterns were obtainedwith a Staib Instruments system using an electron beam of30 kV, a beam current of about 50 μA and an incidence angleof less than 2◦ with respect to the substrate surface. Theincidence directions were parallel to the 〈100〉 of the CGO. Themicrostructure and local chemical composition of the sampleswas characterized by cross-sectional low magnification andhigh resolution transmission electron microscopy (TEM andHRTEM) using a JEOL 2010 FEG electron microscopeoperated at 200 kV (point-to-point resolution 0.19 nm)equipped with a Gatan Image Filter 2000 EELS spectrometerwith an energy resolution of 0.8 eV. Thin foils for TEMobservation were prepared by the conventional cutting, gluingand grinding procedures, followed by a final milling step withAr ions down to perforation.

3. Results and discussion

The effect of Gd doping on the surface morphology of ceriafilms is illustrated in figure 1. The AFM image shown in(a) clearly reveals that non-doped ceria films processed in anAr/5%H2 atmosphere are composed by nanometric roundedgrains. The RHEED pattern depicted in (b) shows diffractionrings, revealing that surface grains are randomly oriented. Incontrast with this morphology, a 10%Gd-doped CeO2 film (thisfilm composition will be hereafter referred to as CGO) exhibitsa terraced topography though, strikingly, its rms roughnesshas increased from 3.2 to 4.7 nm. The RHEED patterncorresponding to this sample exhibits spots characteristic ofa single epitaxial orientation (d). Finally, it is observed that,when doped films are processed in an O2 atmosphere (e), the

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Figure 1. AFM images along with representative line scans, corresponding to the CeO2 (a) and CGO (c) films grown in an Ar/5%H2

atmosphere, and the CGO film post-annealed under O2 (e) and corresponding RHEED patterns (b), (d) and (f), respectively. AFM images areaccompanied by representative line scans. The vertical z-axis coordinate is drawn to the same scale in the three plots.

amount of terraced surface is further increased and the rmsroughness is reduced down to 3.0 nm, bearing witness to anenhanced atomic mobility. Diffuse streaks appearing in theRHEED pattern corresponding to this sample, see (f), can beassociated with the increased flat surface area.

Comparing the line scans shown in (a) and (c), itmay be inferred that, accompanying the slight increase ofsurface roughness mentioned above, the granular to terracedmorphological transformation also favours the formationof pinholes, which tend to smooth out under oxidizingatmospheric conditions (e). This slight increase in roughnesswhen the film becomes epitaxial can be understood consideringthe mechanisms controlling the topographical features offilms in the polycrystalline and epitaxial regimes. In aprevious investigation [20] we have shown that, in thegranular regime, the surface roughness of ceria films isgoverned by the grain size effect (i.e. height variationswithin individual grains) rather than grain boundary energycontributions (i.e. grain boundary grooving) [21]. Therefore,for small grain sizes the surface roughness is kept relativelylow. However, in the epitaxial quasi-single-crystalline film,surface energy anisotropy brings into competition differentexposed crystalline surfaces, namely {111}, {110} and {100},as ordered from the lower to the higher surface energies.As a consequence the surface roughness becomes dominatedby faceting on crystallographic planes inclined to the film

Figure 2. X-ray diffraction θ–2θ scans of the (200) peakscorresponding to the CeO2 and CGO films grown in an Ar/5%H2

atmosphere.

surface. The formation of pinholes is very likely favoured bysuch a strong tendency to facet on planes inclined to the filmsurface [13].

Figure 2 compares the θ–2θ scans corresponding to apure ceria film and a CGO one. An increase by a factor

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Figure 3. XPS characterization corresponding to films 1: CeO2

grown under an oxygen atmosphere (reference), 2: CGO grownunder an oxygen atmosphere, 3: CGO grown under Ar/5%H2, and4: CeO2 grown under Ar/5%H2; (a) C(1s) and (b) Ce(3d) spectra.Inset shows the Ce3+ contributions to samples 2, 3 and 4, obtained bysubtracting the reference spectrum (1).

of two of the intensity of the CGO (002) peak relative tothe CeO2 one is clearly observed, thus signalling that theobserved morphological transformation is accompanied bya dramatic increase of the (001) oriented fraction of thematerial. This enhanced epitaxy necessarily involves a severerecrystallization of the film, which takes place independentlyof the growth atmosphere, temperature and duration.

Since carbon impurities appear at the origin of therecrystallization inhibition mechanism [13, 20], we haveconducted a comparative XPS analysis between CeO2 andCGO films in order to identify the carbon species and thusget insights into their basic interaction mechanisms with themicrostructure. Figure 3(a) compares the C(1s) XPS spectra

corresponding to the CeO2 films grown in O2 (referencesample) and Ar/5%H2 with the CGO films processed underthe same conditions. All the samples show a main peakat 285.4 eV which can be assigned to carbon adsorptionfrom environmental contamination, especially hydrocarbons.However, the Ar/5%H2CeO2 film exhibits a low bindingenergy feature on the main C(1s) peak that corresponds tothe values reported for carbide compounds [22], and it hasto be associated with carbon residues remaining from anincomplete elimination of the metal–organic precursor materialunder reducing atmosphere. Finally, the weak shoulder centredabout 289.5 eV corresponds to carbonate species, a commoncontaminant of oxide surfaces.

On the other hand, figure 3(b) shows the Ce(3d) spectraobtained from the same films. We observe that Ce4+ is the mainoxidation state in all the samples (see, e.g., [23]). However,careful inspection of the spectrum corresponding to the CeO2

film grown in Ar/5%H2 reveals the presence of Ce3+ cations,which manifests as two shoulders centred approximately at 904and 884 eV, as indicated by arrowheads. The inset showsthe Ce3+ contribution obtained after subtracting the Ce4+reference spectrum. We clearly observe that, in both CGOfilms, the Ce3+ contribution is almost absent. The Ce3+ atomicratio derived from these spectra is 10% for the Ar/5%H2 pureceria film, 3% for the Ar/5%H2 GCO film and <2% for the O2

CGO one. As a major result, it appears that the introduction ofGd3+ cations appears almost as efficient as a post-oxygenationstep in removing the carbon residue from the film.

In order to shed light on the microstructural mechanismunderlying the self-purification process induced by Gd3+doping, we have conducted a detailed microstructural study ofthe samples. Figure 4(a) is an HRTEM of the ceria film grownin Ar/5%H2, showing misoriented CeO2 nanocrystals near thesurface. Careful inspection of the image did not reveal anystructural signal to be associated with any crystalline carbidephase, while interstitial space appears filled with amorphousmaterial. An example is arrowed in the image. On the otherhand, an analysis of the lattice parameters indicates that thegrains are essentially stoichiometric ceria [20], thus suggestingthat Ce3+ cations segregate to grain boundaries and interstitialspace as reported by other authors [24, 25]. Indeed, a detailedEELS analysis of CeO2−x nanoparticles has demonstrated that,for sizes in the range 10–20 nm, the valence reduction ofcerium ions occurs mainly at the surface, leaving the core asessentially CeO2 [26]. The observed nanostructure, along withthe scenario outlined above, points to the reaction of peripheralCe3+ cations with the interstitial carbon residue to form anamorphous Ce2C3 phase. On the other hand, figure 4(b) showsa low magnification cross-section TEM micrograph of a CGOfilm viewed along [100] grown under identical conditions. Inagreement with the AFM image shown in figure 1(c), thisimage reveals a terraced topography, consisting of {111}/{110}pyramids truncated by wide (001) planes. Furthermore, thecorresponding SAD pattern shown in (c) reveals a sharpcube-on-cube epitaxial relationship with the YSZ substrate,revealing a severe microstructural transformation throughoutthe film volume.

In order to understand the dramatic influence of Gd3+doping on the microstructural development of the films, we

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Figure 4. (a) HRTEM image of the near-surface region of a CeO2 film grown in an Ar/5%H2 atmosphere; (b) low magnification TEM imageof a CGO film exhibiting a wide (001) terrace; (c) SAD pattern corresponding to the image shown in (b); (d) HRTEM image of a residual grainboundary in a CGO film. The magnification is as in (a); (e) EELS spectra taken from the grain boundary region and bulk matrix shown in (e).

are faced with the basic mechanisms of GB motion. Asdescribed elsewhere [13, 20], as-pyrolyzed films consist of afirst layer of epitaxial crystallites on the substrate, underlying adisordered matrix of nanometric rounded crystallites. Withinthis scenario, we can discern two clear contributions to thedriving force for epitaxial recrystallization: one arising fromthe interfacial energy difference between epitaxial grains anddisordered ones [16] and a second one related to the variationof the chemical potential across a curved interface, forcingGBs to move towards their centre of curvature [17]. While theformer selects the epitaxial orientation to expand throughoutthe film by secondary grain growth, the latter favours theconsumption of smaller spherical grains lying on top of thefirst layer of epitaxial grains. Significantly, both contributionsare independent of doping while appearing intimately relatedto the topological features of the microstructure. Thus, therole of oxygen vacancies is to facilitate diffusion across theGB, i.e. provided the driving force for recrystallization islarge enough to overcome the drag force exerted by theamorphous Ce2C3, oxygen vacancies just constitute an efficientvehicle to facilitate the evolution of the microstructure inthe direction marked by the driving force. Studies on theinfluence of aliovalent doping of ceria ceramics clearly revealthat the GB mobility (Mb) increases with the concentration

of charge-compensating oxygen vacancies (V··O) according

to Mb α [V··O]2 [19]. According to this mechanism, and

at variance from our observations for pure ceria, thermalannealing of the nanostructured precursor under a reducingatmosphere would increase the GB mobility owing to theoxygen vacancies accompanying the partial reduction of Ce4+cations. This discrepancy can be understood in the lightof the present XPS and microstructural results, pointing toa segregation of Ce3+ cations towards the grains peripherywhere they become consumed to form Ce2C3. This behaviourhas two negative and unavoidable effects, namely, Ce3+cations accompanying the formation of oxygen vacanciesfeed themselves on the amorphous carbide that inhibits GBmotion, and as a consequence segregated Ce3+ cations becomeinefficient to generate oxygen vacancies within the interior ofthe nanocrystallites. In this regard, aliovalent Gd3+ ↔ Ce4+substitutions provide an alternative means for the generationof charge-compensating oxygen vacancies within the interiorof the nanocrystallites independently of the atmosphericconditions, according to the reaction

Gd2O3 = 2Gd′Ce + V··

O + 3OO

which indicates that each oxygen vacancy (V··O) is generated

by the substitution of two Ce4+ cation sites by Gd3+ cations

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Nanotechnology 19 (2008) 395601 M Coll et al

(denoted as 2Gd′Ce). Therefore, substitution by Gd3+

constitutes an efficient and tunable way to generate oxygenvacancies within the grains, thus potentially enhancing Mb ,regardless of the atmospheric conditions.

The proposed mechanism necessarily involves the pushingand elimination of the carbide phase by the growth frontthrough the surface of the film. Different reactions then makepossible the release of the carbide phase from the film surface.As the growing atmosphere is Ar + H2, H2 gas may react withcarbides leading to methane gas (CH4). Alternatively, carbidescan also be released from the film through the formation of COgas as long as oxygen is available in CeO2 [27].

Supporting the proposed pushing mechanism and therole of carbon-containing impurities in blocking GB motion,figure 4(d) shows a residual GB in a CGO film. Comparison ofthe EELS spectra obtained from the GB and the grain interior(e) reveals that the former exhibits a C–K peak while thelatter does not, hence indicating that the only GBs remainingin the final microstructure are those for which the carbon-containing impurities become accidentally trapped after therecrystallization process.

As far as the surface morphology is concerned, it isworth noting that, in contrast with the granular film, wherethe major contribution to surface roughness comes from grainsize effects [20], an important contribution in (001) orientedepitaxial films comes from the strong tendency of such filmsto facet on pyramidal {111} and {110} facets [20, 28], whichcorrespond to the sharpest minima in the Wulff plot [29].Surprisingly, in this scenario the faceting contribution tosurface roughness is only balanced by truncation of thepyramidal hillocks by the otherwise energetically prohibitive(001) polar surfaces [30]. As a matter of fact, the surfaceroughness is even lower in the granular film (rms = 3.2 nm)than in the epitaxial one (rms = 4.7 nm). Thus, the surfacemorphology is dictated by the competition between faceting onlow energy {111} and/or {110} pyramidal planes and truncationof those pyramids through the stabilization of (001) surfaces.An example is presented in figure 4(b), showing a large (001)terrace bounded by {111} or {110} facets (note that a traceinclined by 45◦ in [100] projection may correspond either to{111} or {110} planes). The latter mechanism involves thesuppression of the dipole moment associated with the (001)surface, which in pure ceria has been reported to be achievedby substantial reconstruction of the oxygen sublattice [31, 32].In the present case charge compensation among Gd3+, Ce4+and oxygen along 〈100〉 directions may constitute an additionalresource and therefore CGO may be more prone to developlarge area (001) surfaces than pure ceria. Besides this complexinterplay among surface energies, a significant contribution tosurface roughness in CGO films may result from grooving atresidual C-containing grain boundaries. This is evidenced bythe fact that, after annealing in an O2 atmosphere, the rmssurface roughness of the CGO films is reduced from 4.7 nmdown to 3.0 nm (compare figures 1(c) and (e)).

4. Conclusion

In summary, a new microstructural mechanism promotingepitaxial recrystallization in an as-pyrolyzed solution-derived

nanocrystalline ceria film has been investigated. The presentmechanism manifests as efficient as the conventional oxidationprocess for the elimination of C-containing impurities, butworks under reducing atmospheric conditions, thus beingspecially suited for processing samples containing oxidation-sensitive materials like metallic tapes in coated conductorarchitectures. Epitaxial recrystallization consists of thesecondary growth of epitaxial crystallites at the expenseof the bulk disordered ones, which appear frozen byan amorphous Ce2C3 phase decorating and filling grainboundary and interstitial space, exerting a drag force thatcannot be compensated by the driving force for epitaxialrecrystallization. The formation of this phase results fromthe combination of a C-residue from the metal–organicsolution and Ce3+ cations segregated at the surface ofthe ceria nanocrystallites. The alternative microstructuralmechanism presented here is based on an enhanced grainboundary mobility achieved by Gd3+ doping, favouring thepushing of the carbide phase by the growth front untilits elimination through the surface of the film. It turnsout that the microstructural process, by itself, constitutesa self-purification mechanism analogue to those known fordirectional solidification in which the growth medium is a solidnanocrystalline disordered precursor.

Acknowledgments

We acknowledge financial support from MEC (NANOARTIS,MAT2005-02047); NANOFUNCIONA, NAN2004-09133-CO3-01; CSIC (PIF-CANNAMUS); EU (HIPERCHEM,NMP4-CT2005-516858; SUPER3C, SES6-CT-2004-502615)and Generalitat de Catalunya (Catalan Pla de RecercaSGR-0029 and CeRMAE). Serveis Cientıfic-Tecnics de laUniversitat de Barcelona is acknowledged for TEM facilities.

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