Pr

HBa

b

c

d

a

ARR2A

KEESTC

1

tpiDsCyapafcy

0d

Materials Chemistry and Physics 122 (2010) 623–629

Contents lists available at ScienceDirect

Materials Chemistry and Physics

journa l homepage: www.e lsev ier .com/ locate /matchemphys

reparation and structural characterization of SnO2 and GeO2 methanol steameforming thin film model catalysts by (HR)TEM

arald Lorenza, Qian Zhaob, Stuart Turnerc, Oleg I. Lebedevc, Gustaaf Van Tendelooc,ernhard Klötzera, Christoph Rameshana,d, Simon Pennera,∗

Institute of Physical Chemistry, University of Innsbruck, A-6020 Innsbruck, AustriaSchool of Materials Science and Engineering, Tianjin University, Tianjin 300072, People’s Republic of ChinaEMAT, University of Antwerp, B-2020 Antwerp, BelgiumDepartment of Inorganic Chemistry, Fritz-Haber Institute of the Max-Planck Society, D-14195, Berlin, Germany

r t i c l e i n f o

rticle history:eceived 21 October 2009eceived in revised form3 December 2009ccepted 23 March 2010

eywords:lectron microscopylectron energy-loss spectroscopy

a b s t r a c t

Structure, morphology and composition of different tin oxide and germanium oxide thin film catalysts forthe methanol steam reforming (MSR) reaction have been studied by a combination of (high-resolution)transmission electron microscopy, selected area electron diffraction, dark-field imaging and electronenergy-loss spectroscopy. Deposition of the thin films on NaCl(0 0 1) cleavage faces has been carried outby thermal evaporation of the respective SnO2 and GeO2 powders in varying oxygen partial pressuresand at different substrate temperatures. Preparation of tin oxide films in high oxygen pressures (10−1 Pa)exclusively resulted in SnO phases, at and above 473 K substrate temperature epitaxial growth of SnOon NaCl(0 0 1) leads to well-ordered films. For lower oxygen partial pressures (10−3 to 10−2 Pa), mixtures

nOhin filmsatalysis

of SnO and �-Sn are obtained. Well-ordered SnO2 films, as verified by electron diffraction patterns andenergy-loss spectra, are only obtained after post-oxidation of SnO films at temperatures T ≥ 673 K in 105 PaO2. Preparation of GeOx films inevitably results in amorphous films with a composition close to GeO2,which cannot be crystallized by annealing treatments in oxygen or hydrogen at temperatures comparableto SnO/SnO2. Similarities and differences to neighbouring oxides relevant for selective MSR in the thirdgroup of the periodic system (In2O3 and Ga2O3) are also discussed with the aim of cross-correlation in

als, an

formation of nanomateri

. Introduction

SnO2 is one of the most extensively studied and best charac-erized oxides thanks to its outstanding chemical and electronicroperties. A comprehensive review of all aspects especially of

ts surface and materials science has been given by Batzill andiebold [1]. Recent important applications of SnO2 range from

olid-state gas sensors (especially for reducing gases such as H2,O or H2S) over transparent conductors to heterogeneous catal-sis [1,2]. As these applications strongly depend on the structurend morphology of the SnO2 samples, special attention has been

aid to the development of reproducible preparation methods ofvariety of SnO2 samples, especially ordered thin films suitable

or (high-resolution) transmission electron microscopy HR(TEM)haracterization. These methods include sputtering, spray pyrol-sis, electron-beam evaporation, chemical vapour deposition or

∗ Corresponding author. Tel.: +43 5127055056; fax: +43 5125072925.E-mail address: simon.penner@uibk.ac.at (S. Penner).

254-0584/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.matchemphys.2010.03.057

d ultimately, also catalytic properties.© 2010 Elsevier B.V. All rights reserved.

rheotaxial growth and thermal oxidation, among others [3–7]. Aconcise discussion of different preparation techniques for thin filmshas been outlined by Sberveglieri [8]. Regarding heterogeneouscatalysis, the activity and selectivity of SnO2 catalysts can be signif-icantly improved by doping with other metals, such as Cu, Pd, Cr orSb, but even pure SnO2 itself is an active oxidation catalyst, espe-cially in CO oxidation reactions. Since a Mars-van-Krevelen-typereaction mechanism is assumed to be prevalent, a central issue inoxidation catalysis of SnO2 refers to its easy (surface) reducibility[1]. Most important, recent investigations indicate that pure SnO2exhibits a pronounced activity and/or selectivity in methanol steamreforming [9]. The interest to include SnO2 and GeO2 in studieson methanol steam reforming is derived from previously stud-ied oxides such as ZnO [10], Ga2O3 [11] or In2O3 [12]. Since Sn4+

and In3+ represent isoelectronic species a potential functional anal-ogy also of their catalytic action may be anticipated. As the same

4+ 3+ 2+

holds for Ge , Ga and Zn , extension of the studies to oxidesof the fourth main group of the periodic system is straightfor-ward. Contrary to SnO2, structural information on GeO2 thin films isscarce and mainly limited to the preparation of GeO2 nanomaterials[13–15].

6 istry and Physics 122 (2010) 623–629

cemtotptctt(tGsowotaaagm

2

osoevtm(rateaahiawtpgsm

Fs

24 H. Lorenz et al. / Materials Chem

In order to close the materials gap between surface studies espe-ially on tin oxides, where a considerable amount of data alreadyxists [1] and studies on “real world” powder catalysts, the thin filmodel concept will be exploited. From the structural point of view,

hin films grown on well-ordered substrates are usually epitaxialr at least well-ordered, that is, structurally resembling single crys-al studies conducted in UHV. They usually also exhibit a narrowarticle size distribution and well-defined particle sizes. Since athe same time, our thin film dedicated micro-reactor setup allowsorrelation of the catalytic activity/selectivity measured on bothhin film and powder samples [11,12], structure–activity correla-ions of both model systems (thin films) and real-world catalystspowders) can be easily established. The aim of the present con-ribution therefore is – in close correlation to similar studies ona2O3 and In2O3 [11,12] – to use NaCl(0 0 1) cleavage faces as sub-trates to eventually grow single-crystalline tin and germaniumxide films to present an easy and reproducible preparation path-ay to different thin film model systems of tin and germanium

xides and subsequently their thorough structural characteriza-ion, prior to catalytic testing. Adequate tools to fulfil this taskre considered HRTEM, selected area electron diffraction (SAED)nd electron energy-loss spectroscopy (EELS). Special attention willlso be paid to the influence of deposition parameters (O2 back-round pressure, substrate temperature) on film stoichiometry andorphology.

. Experimental

All tin and germanium oxide films were prepared in a high-vacuum chamberperated at a base pressure of 10−4 Pa. Film thicknesses (usually 25 nm) were mea-ured in situ by a quartz crystal microbalance. SnO2 and GeO2 powders (tin(IV)xide, 99.99%, and germanium(IV) oxide, 99.9999%, both Alfa Aesar) were thermallyvaporated from tantalum crucibles onto vacuum-cleaved NaCl(0 0 1) surfaces atarying substrate temperatures (298–603 K) and O2 background pressures (10−3

o 10−1 Pa). The films were floated and rinsed with distilled water, dried and finallyounted on gold grids for electron microscopy. Annealing and oxidative treatments

105 Pa He or O2 for 1 h at 473 K, respectively) were performed in a circulating batcheactor. The structure and morphology of these thin films in the as-deposited statend upon oxidative and annealing treatments were monitored by (high-resolution)-ransmission electron microscopy (HRTEM). The electron diffraction patterns werexternally calibrated with respect to the (1 1 1), (2 0 0) and (2 2 0) Pd reflections ofn as-deposited, untreated Pd/Ga2O3 catalyst. Overview TEM imaging and selectedrea electron diffraction (SAED) were carried out with a ZEISS EM10C microscope,igh-resolution imaging, electron energy-loss spectroscopy (EELS), and dark-field

maging were performed using a JEOL 4000EX high-resolution microscope operatedt 400 kV and a JEOL 3000F FASTEM microscope operated at 300 kV and equipped

ith a GIF 2000 post-column electron energy-loss spectrometer. The film composi-

ion was checked by energy-dispersive X-ray analysis (EDXS), which only showedeaks due the evaporated elemental thin film constituents (Sn, Ge and O) and theold grid (Au). Surface carbon impurities present on the films were removed by Ar-puttering prior to TEM imaging. No influence of the sputtering on structure andorphology of the films has been detected. The purity of the substrate was ensured

ig. 2. Overview bright-field TEM image of a SnO thin film after deposition at 473 K on Natructure of the film.

Fig. 1. Overview bright-field TEM image of a SnO thin film after deposition at 373 Kon NaCl(0 0 1) in 10−1 Pa O2. The SAED pattern is shown as an inset.

by freshly cleaving the NaCl(0 0 1) crystals immediately before deposition of theoxide.

3. Results and discussion

3.1. SnO and SnO2 thin films

3.1.1. Preparation of SnO thin film model catalystsThe structure of the SnO films after deposition in 10−1 Pa

O2 at low substrate temperatures (373 K) is shown in Fig. 1.Although the film exhibits no pronounced contrast in bright-fieldTEM images, its regular grain-size structure is easily visible. Theaverage size of the round-shaped grains is approximately 10 nm.The corresponding SAED pattern (shown as inset) represents analmost amorphous structure and shows a single blurred diffrac-tion ring. Raising the substrate temperature to 473 K convertsthe film into a single-crystalline phase, as revealed by the SAEDpattern shown in Fig. 2b. Fourfold split intense diffraction spots

at 2.7, 1.9 and 1.3 Å correspond to the SnO (1 1 0), (2 0 0) and(2 2 0) lattice spacings of the tetragonal SnO phase (lattice con-stants a = 3.79 Å and c = 4.84 Å, respectively [dtheor(1 1 0) = 2.686 Å,dtheor(2 0 0) = 1.899 Å, dtheor(2 2 0) = 1.343 Å] [16]. As expected for

Cl(0 0 1) in 10−1 Pa O2 (a). The SAED pattern (b) evidences the tetragonal SnO crystal

H. Lorenz et al. / Materials Chemistry and Physics 122 (2010) 623–629 625

Fig. 3. High-resolution electron micrograph of the SnO thin film shown in Fig. 2.Toa

twTaotie(Fd∼sstapatoeSsattgia

tsMScrf

tin oxides at oxygen background pressures higher than 10−1 Pa.

he inset FFT pattern shows a strong ring corresponding to the 3.0 Å {1 0 1} spacingf tetragonal SnO. A single reflection corresponding to the 2.7 Å {1 1 0} spacing islso present.

he tetragonal phase, the (1 1 0)/(2 −2 0) spots form an angle of 45◦

ith the (2 0 0) diffraction spots. The corresponding bright-fieldEM image (Fig. 2a) exhibits much enhanced diffraction contrastnd shows extended dark domains corresponding to differentlyriented SnO grains. These observations are corroborated by addi-ional high-resolution imaging, shown in Fig. 3. The HRTEM imagen Fig. 3 reveals a complex arrangement of individual SnO grainsach exhibiting lattice fringes of ∼3.0 Å, corresponding to the SnO1 0 1) lattice spacings [dtheor(1 0 1) = 2.989 Å] [16]. The inset fastourier transform pattern (i.e. the micro-diffraction pattern) evi-ences these extended SnO(1 0 1) lattice spacings by an ring at3.0 Å, in agreement with the SAED patterns obtained from this

ample. The FFT pattern also shows some reflections at ∼2.7 Åpacing, corresponding to the SnO (1 1 0) lattice spacing. Both elec-ron diffraction and HRTEM therefore unambiguously confirm thatfter deposition at 473 K in 0.1 Pa O2 a single, well-ordered SnOhase is formed. Furthermore, the SAED pattern of Fig. 2b suggestspreferential epitaxial growth of SnO on NaCl(0 0 1) in the rela-

ionship SnO[0 0 1]//NaCl[0 0 1] and SnO[0 1 1]//NaCl[0 1 1]. Thesebservations agree well with previous studies of Choi et al. or Fengt al., where highly oriented tin oxide films (mostly tetragonalnO2) were successfully grown on single-crystalline Al2O3 sub-trates [2,17]. From dark-field experiments (shown in Fig. 4) welso infer that the contrast variations are due to diffraction con-rast of differently oriented SnO grains. The dark-field image wasaken by using the SnO (2 0 0) diffraction spot. Consequently, grainsiving rise to this diffraction spot appear bright in the dark-fieldmage. All other differently oriented grains remain dark (see e.g.reas indicated by arrows in the lower left corner of the image).

A more complex scenario is present if the substrate tempera-ure is raised to 573 K (Fig. 5). The bright-field TEM image in Fig. 5ahows a complete reconstruction of the SnO structure and extendedoire-fringes indicate overlapping lattice of differently oriented

nO grains. However, the SAED patterns (not shown) are identi-al to the one shown in Fig. 2b, indicating that most of the filmemains SnO. Some small areas of the film have however almostully oxidised to SnO2 as a result of the raising of the substrate

Fig. 4. Dark-field image of the SnO film shown in Fig. 2 obtained using the SnO(2 0 0)diffraction reflection. The bright regions correspond to grains giving rise to the 2 0 0reflection. Other differently oriented grains remain dark (examples indicated byarrows).

temperature. The high-resolution image in Fig. 5b shows such anarea. All reflections in the FFT pattern fit with the tetragonal SnO2structure imaged along the [0 0 1] zone axis orientation.

The effect of O2 background pressure variation at constant sub-strate temperature is highlighted in Fig. 6. The left panel showsgrowth of SnO at 473 K in 10−2 Pa O2, on the right panel the samesituation in 10−3 Pa O2 is outlined. In the TEM image of Fig. 6a theformer SnO structure formed after deposition at 473 K in 10−1 PaO2 is still prevailing, but additional round-shaped grains of about20–30 nm size are also observed. The SAED pattern (inset) stillshows strong diffraction spots of the tetragonal SnO phase, but alsoadditional Debye–Scherrer-type diffraction rings. A detailed analy-sis reveals the presence of a second metallic �-Sn phase, with strongreflections at 2.9 and 2.8 Å, corresponding to the �-Sn(2 0 0) and�-Sn(1 0 1) fringes [dtheor(2 0 0) = 2.915 Å and dtheor(1 0 1) = 2.792 Å][18]. It is worth noting, that these sphere-like structure correspondsto the first stages of oxide film preparation by rheotaxial growthand thermal oxidation [7,8], where metallic tin is deposited to formlarge droplets and subsequently oxidized to result in an oxide filmwith a high surface-to-volume ratio. These features are even morepronounced if the deposition is carried out with less O2 (Fig. 6b). At10−3 Pa O2 background pressure the film structure does not resem-ble SnO, but rather shows very large sphere-like aggregates (up to50 nm in size). The SAED patterns still show SnO, but the intensityof the �-Sn phase increased considerably.

The issue of influence of O2 background pressure on phase for-mation has also been addressed by Jimenez et al. upon SnO andSnO2 deposition on MgO and SiO2 [19]. In agreement with thepresent studies, heating of SnO2 and post-exposition to 0.1 Torr(∼13 Pa) O2 for 300 s resulted only in SnO stoichiometry. Completeoxidation to SnO2 was only obtained after exposure to an oxygenplasma at 300 K. In general, SnO thin films are mostly prepared viaoxidation of Sn layers [20] or spray pyrolysis [21]. However, theabove-outlined preparation method is unique since it yields repro-ducible and well-ordered SnO films at low substrate temperatures.

3.1.2. Preparation of SnO2 thin film model catalystsThe question how to reproducibly prepare SnO2 thin films is

intimately connected with the ability to carry out deposition of

Usually, e.g. for Ga2O3 or In2O3, an oxygen background pressure of10−2 Pa is sufficient to fully oxidize the respective thin films duringdeposition. Deposition of tin oxides is different since both of theoxides, SnO and SnO2, are thermodynamically stable oxides and

626 H. Lorenz et al. / Materials Chemistry and Physics 122 (2010) 623–629

Fig. 5. Overview bright-field TEM image of a SnO thin film after deposition at 573 Kon NaCl(0 0 1) in 10−1 Pa O2 (a) and a HRTEM image showing a small SnO2 regionwithin the SnO film (b). The [0 0 1] zone axis orientation and structure of the SnO2

region is evidenced by the inset FFT pattern.

Fig. 6. Overview bright-field TEM image of a SnO thin film after deposition at 473 K on Na

Fig. 7. Overview bright-field TEM image of a SnO thin film after deposition at 573 Kon NaCl(0 0 1) in 10−1 Pa O2 after post-oxidation at 673 K in 105 Pa O2 for 1 h.

deposition in 10−1 Pa O2 only leads to SnO phases. Higher oxygenbackground pressures, even if they are needed for SnO2 deposition,are basically excluded since the deposition rate decreases drasti-cally for very high oxygen background pressures due to oxidativeinhibition of the formation of volatile gas phase species. Hence, theonly preparation pathway to SnO2 thin films, if the deposition iscarried out by thermal evaporation of SnO2 is via formation of SnOand adequate post-oxidation procedures.

Fig. 7 highlights the structure of a previous SnO film preparedat 473 K in 10−1 Pa oxygen background pressure and post-oxidizedat 673 K in 105 Pa O2. The overview TEM image reveals a coarsegrain structure similar to the original SnO film and exhibits strongdiffraction contrast. However, the corresponding HRTEM imagesand SAED patterns indicate formation of SnO2. Fig. 8 shows a high-resolution TEM image with extended SnO2(1 1 0) lattice fringesof 3.30 Å spacing, corroborating the SAED analysis. The inset EDring pattern mainly shows rings corresponding to the SnO2 struc-ture (at 3.4, 2.6, 2.4 and 1.8 Å, corresponding to the (1 1 0), (1 0 1),(2 0 0) and (2 1 1) reflections of the tetragonal SnO structure

2(a = 4.73 Å, c = 3.18 Å; dtheor(1 1 0) = 3.347 Å; dtheor(1 0 1) = 2.643 Å;dtheor(2 0 0) = 2.369 Å; dtheor(2 1 1) = 1.764 Å] [22]). It is worth not-ing, that SnO2 appears to be at least partially ordered, as judgedfrom the elongated diffraction spots overlapping with the diffrac-

Cl(0 0 1) in 10−2 Pa O2 (a) and 10−3 Pa O2 (b). The SAED patterns are shown as insets.

H. Lorenz et al. / Materials Chemistry a

Firi

tvcttt1rtod

FpFab5

ig. 8. High-resolution electron micrograph of the thin film shown in Fig. 7. Thenset SAED pattern evidences the tetragonal SnO2 crystal structure, a faint diffuseing corresponding to the SnO (1 0 1) spacing can also be made out. The arrowsndicate examples of widespread defects (stacking faults) present in the film.

ion rings. A single diffuse ring at 3.0 Å is present, arising from aery small amount of SnO. The arrows in the HRTEM image indi-ate the wide spread defective (mainly stacking faults) structure ofhe SnO2 film. Additional information on the formation of differentin oxide species is provided by EELS. Fig. 9 shows an overview ofhe Sn M4,5 and O–K edges of SnO films after deposition at 473 K in0−1 Pa O (middle), after post-oxidation at 673 K (top) and after

2eduction at 573 K in 105 Pa hydrogen (bottom). The interaction ofhe tin oxides with hydrogen is discussed in part B, but for the sakef clarity, the EELS spectrum is included in this figure. Due to theifferent structural environments in SnO and SnO2 (tetrahedral vs.

ig. 9. EELS spectra (Sn M4,5 edge at 485 eV and O–K edge at 532 eV) of differentlyrepared tin oxide thin films (top: SnO2 corresponding to the thin film shown inig. 7; middle: SnO corresponding to the film shown in Fig. 2; bottom: SnO thin filmfter reduction in H2 at 573 K (105 Pa, 1 h)). The two oxide phases can be discernedy the splitting of the fine structure after 530 eV as well as the elevated peak at50–560 eV in SnO2 (indicated by *) [27].

nd Physics 122 (2010) 623–629 627

trigonal oxygen coordination in the first coordination shell) [23]the electron-energy near-edge structure of the Sn M4,5 and O–Kedges is a reliable tool to distinguish between the different typesof tin oxides. The top-most spectrum is typical for SnO2, with awider splitting of the peaks in the region 530–550 eV and the morepronounced multiple scattering resonances at higher energy losses(around 575 eV). In contrast, the middle spectrum shows two peaksof almost equal intensity with a narrower splitting and a slightlyless pronounced ELNES around 570–575 eV. These differences arebasically due to the different coordination of oxygen in both tinoxides [23]. The bottom spectrum, corresponding to �-Sn accord-ing to the SAED patterns, only shows a broad edge typical for adelayed Sn M4,5 edge [23].

3.2. GeO2 thin films

The preparation of ordered, single-phase GeO2 thin films is ham-pered by the polymorphism of GeO2, that is, it exists not onlya hexagonal �-quartz-type and a tetragonal rutile-type structurebut also a glassy amorphous phase depending on pressure, tem-perature and preparation conditions. Concerning the preparationof thin films using NaCl as growth templates, the solubility of thedifferent polymorphs in water is most crucial. Whereas �-quartz-type GeO2 is slightly soluble in water, rutile-type GeO2 is insoluble[24]. Depending on which phase forms, floating the thin films inwater therefore might pose problems, as was the case for ZnO[25]. However, as discussed below, we succeeded in preparingself-supporting GeO2 films, which are not affected by the floatingprocedure. Generally, GeO2 samples are almost uniquely preparedby oxidative treatments of Ge films or powders [18,26].

For GeO2 we followed a similar strategy as for SnO and SnO2and tried to assess structure, composition and influence of depo-sition parameters. Fig. 10 shows a bright-field overview TEMimage of a GeO2 thin film prepared at 300 K and 10−2 Pa O2 back-ground pressure. Both the image and the SAED pattern (inset)indicate the presence of an amorphous germanium oxide phase.

It is worth to note, that neither raising the substrate tempera-ture to 623 K nor post-oxidation treatments up to 673 K lead tothe formation of single-crystalline phases and only low contrastamorphous structures are obtained. Recent studies have shownthat thin films prepared by evaporation of GeO2 are usually sub-

Fig. 10. Overview bright-field TEM image of a GeO2 thin film after deposition at300 K in 10−2 Pa O2. The SAED pattern is shown as an inset.

628 H. Lorenz et al. / Materials Chemistry a

Fpsb

svspr3i(totatala�tttordso

cicpa

4

wtravgtfi

ig. 11. EELS spectra (O–K edge at 532 eV and Ge L2,3 at 1217 eV) of differentlyrepared germanium oxide thin films (top: GeO2 corresponding to the thin filmhown in Fig. 10; middle: GeO2 film after oxidation in O2 at 673 K (105 Pa, 1 h);ottom: GeO2 thin film after reduction in H2 at 673 K (105 Pa, 1 h).

toichiometric and can only be crystallized by post-annealing atery high temperatures [27,28]. Our observations agree well withtudies of Ardyanin et al. on the structure and photoluminescenceroperties of e-beam evaporated GeOx films. GeO2 was evapo-ated in UHV at 10−8 Torr (∼10−6 Pa; with pressure increase to× 10−6 Torr (∼3 × 10−4 Pa) during deposition) at 373 K on sil-

con substrates and only amorphous GeOx films were obtainedfilm thickness: 200 nm). Only after annealing at 773 K, dispropor-ion of GeOx into crystalline Ge metal and amorphous GeO2 wasbserved [28]. As the method of choice to gain information onhe composition of the amorphous GeO2 films, EELS spectra of thes-deposited thin film, after oxidation at 673 K and after reduc-ion in hydrogen at 673 K were collected. This will also ensuredirect correlation to the properties of SnOx films. Fig. 11 high-

ights a set of O–K edges and Ge L2,3 edges of the films after thebove mentioned treatments. In short, these edges are typical of-quartz-type GeO2 [29] and no substantial differences between

he spectra can be detected. Most important, the spectra allowhe determination of the film stoichiometry by quantification ofhe oxygen-to-germanium ratios. For the as-deposited film and thexidized sample, ratios of O:Ge = 66 at.%:34 at.% and 67 at.%:33 at.%,espectively, have been obtained. This indicates that already aftereposition the film stoichiometry is close to GeO2. For the reducedample, this ratio shifts to O:Ge = 58 at.%:42 at.%, indicating that theverall film stoichiometry changes towards GeO.

We therefore assume that the films are close to a GeO2 stoi-hiometry in the as-deposited state and annealing treatments (alson hydrogen and oxygen) up to 673 K, which are typically used foratalyst regeneration treatments, are not sufficient to induce a com-lete reduction of GeO2, disproportion of GeOx or the formation ofny crystalline phase (metallic or oxidic).

. Conclusions

We successfully demonstrated reproducible preparation path-ays to thin film model catalysts of SnO, SnO2 and GeO2 by simple

hermal evaporation of SnO2 and GeO2 in different oxygen envi-onments and, for SnO2, by post-oxidation of the resulting films

t 673 K. Generally, some differences and similarities to the pre-iously studied neighbouring oxides (In2O3, Ga2O3) in the thirdroup of the periodic system are notable [11,12]. From the struc-ural point of view, SnO/SnO2 more closely resembles In2O3 thinlms in terms of crystallinity and epitaxial growth. Both materials

nd Physics 122 (2010) 623–629

show a strong dependence of the film structure and morphologyon the substrate temperature, although formation of nanomateri-als, such as octahedrons or pyramidal-shaped particles as obtainedfor In2O3, has never been observed. On the contrary, GeO2 filmsresemble Ga2O3 structures, at least at low substrate temperatures.For both materials, amorphous structures are usually observed,even after oxidative and annealing treatments at temperatures ashigh as 673 K. As for SnO/SnO2, no nano-architectures have beenobserved for GeO2 even at high substrate temperatures (623 K). Thisbehaviour might be interpreted in terms of the presence of differentgrowth species during oxide deposition and the resulting alteredwetting properties on NaCl(0 0 1). For Ga2O3 and In2O3, thesegrowth species were basically associated with sub-stoichiometricoxides (e.g. Ga2O or In2O) with a high mobility on NaCl(0 0 1) at highsubstrate temperatures. Therefore, the formation of the sphere-like aggregates for Ga2O3 and pyramidal-shaped particles for In2O3is the result of efficient transport processes preceding the oxida-tion of the volatile sub-stoichiometric oxides and their re-oxidationtowards ordered aggregates (apart from crystallographic matchingof In2O3 and NaCl). For SnO2 we might interpret the growth speciesas SnO and/or Sn metal, as confirmed by the SAED patterns of filmsprepared at different oxygen partial pressures. The highest oxygenpartial pressures that can be used in our HV chamber during depo-sition (10−1 Pa) are sufficient to convert Sn metal into SnO, but nofurther oxidation to SnO2 can be achieved. The same probably holdsfor GeO2, but since only amorphous films are obtained, a thoroughstructural characterization by TEM, apart from EELS, is excluded.

It is clear, that the outlined results represent only a first steptowards the full characterization of these oxide systems also incatalytic reactions. Since partly the aim of this study is the correla-tion of the results to In2O3 and Ga2O3, a thorough catalytic testingof these systems in methanol steam reforming has been carriedout and discussed in detail elsewhere, also in the context of acidicand basic surface sites leading to different activity and selectivitypatterns for SnO, SnO2 and GeO2 [30].

In any case, both oxides are promising candidates as supportsfor small Pd catalyst particles, since they can be reproducibly pre-pared as free-standing, stable oxide support films, into which thenoble metal particles can easily be embedded. Due to their amor-phicity, in TEM experiments GeO2 films especially favour the easydiscrimination of the Pd particles within the oxide matrix due toenhanced contrast, as it has been, e.g. the case for Pd/Ga2O3 [11].

Hence, bimetallic single phases of Pd–Sn and Pd–Ge might beprepared by reduction in hydrogen and catalytically characterizedin methanol steam reforming, in close correlation to Pd supportedon ZnO, Ga2O3 and In2O3.

Acknowledgements

We thank the FWF (Austrian Science Foundation) for financialsupport under project P20892-N19. The authors acknowledge sup-port from the European Union under the Framework 6 programunder a contract from an Integrated Infrastructure Initiative (Ref-erence 026019 ESTEEM).

References

[1] M. Batzill, U. Diebold, Prog. Surf. Sci. 79 (2005) 47–154.[2] Y. Choi, S. Hong, Sens. Actuators B 125 (2007) 504–509.[3] S. Semancik, R. Cavicchi, Thin Solid Films 206 (1991) 81–87.[4] K. Murakami, K. Nakajima, S. Kaneko, Thin Solid films 515 (2007) 8632–8636.

[5] C.F. Wan, R.D. McGrath, W.F. Keenan, S.N. Frank, J. Electrochem. Soc. 136 (1989)

1459–1463.[6] Q. Wu, J. Song, J. Kang, Q. Dong, S. Wu, S. Sun, Mater. Lett. 61 (2007) 3679–3684.[7] G. Sberveglieri, G. Faglia, S. Gropelli, P. Nelli, Proc. 6th Int. Conf. Solid-State

Sensors and Actuators, San Francisco, CA, USA, June 23–27, 1991, pp. 165–168.[8] G. Sberveglieri, Sens. Actuators B 6 (1992) 239–247.

istry a

[[

[

[[

[[[[[[

[[[

[

[

[

[27] M. Ardyanin, H. Rinnert, M. Vergnat, J. Lumin. 129 (2009) 729–733.

H. Lorenz et al. / Materials Chem

[9] T. Mori, S. Hoshino, A. Neramittagapong, J. Kubo, Y. Morikawa, Chem. Lett. 13(2002) 390–391.

10] C. Wöll, Prog. Surf. Sci. 82 (2007) 55–120.11] S. Penner, H. Lorenz, W. Jochum, M. Stöger-Pollach, D. Wang, C. Rameshan, B.

Klötzer, Appl. Catal. A 358 (2009) 193–202 (and references therein).12] H. Lorenz, W. Jochum, B. Klötzer, M. Stöger-Pollach, S. Schwarz, K. Pfaller, S.

Penner, Appl. Catal. A 347 (2008) 34–42 (and references therein).13] C. Jing, J. Hou, Y. Zhang, J. Cryst. Growth 310 (2008) 391–396.14] V.V. Atuchin, T.A. Gavrilova, S.A. Gromilov, V.G. Kostrovsky, L.D. Pokrovsky, I.B.

Troitskaia, R.S. Vemuri, G. Carbajal-Franco, C.V. Ramana, Cryst. Growth Des. 9(2008) 1829–1832.

15] H. Kim, S. Shim, J. Lee, Appl. Surf. Sci. 253 (2007) 7207–7210.16] F. Izumi, J. Solid State Chem. 38 (1981) 381–385.17] X. Feng, J. Ma, F. Yang, F. Li, C. Luan, Mater. Lett. 62 (2008) 1809–1811.18] J.A. Lee, G.V. Raynor, Proc. Phys. Soc., Lond. 67 (1954) 737–747.19] V.M. Jimenez, J.P. Espinos, A.R. Gonzalez-Elipe, Surf. Sci. 366 (1996) 556–563.20] W.E. Boggs, P.S. Trozzo, G.E. Pelissier, J. Electrochem. Soc. 108 (1961) 124–129.

[

[[

nd Physics 122 (2010) 623–629 629

21] L. Sodomka, A. Kleprik, Jemna Mech. Opt. 8 (1963) 43–46.22] W. Baur, Acta Crystallogr. A 9 (1956) 515–520.23] M.S. Moreno, R.F. Egerton, J.J. Rehr, P.A. Midgley, Phys. Rev. B 71 (2005) 035103-

1–035103-6.24] R.J. Meyer, Gmelins Handbuch der Anorganischen Chemie, 8. Auflage, 1958,

Verlag Chemie, GmbH Weinheim/Bergstrasse.25] S. Penner, B. Jenewein, H. Gabasch, B. Klötzer, D. Wang, A. Knop-Gericke, R.

Schlögl, K. Hayek, J. Catal. 241 (2006) 14–19.26] Y. Su, X. Liang, S. Li, Y. Chen, Q. Zhou, S. Yin, X. Meng, M. Kong, Mater. Lett. 62

(2008) 1010–1013.

28] M. Ardyanin, H. Rinnert, X. Devaux, M. Vergnat, Appl. Phys. Lett. 89 (2006)011902-1–011902-3.

29] N. Jiang, J. Spence, J. Non-Cryst. Solids 353 (2007) 2813–2816.30] Q. Zhao, H. Lorenz, S. Turner, O.I. Lebedev, G. van Tendeloo, C. Rameshan, B.

Klötzer, S. Penner, Appl. Catal. A 375 (2010) 188–195.