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Vacuum 85 (2010) 535e540

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Vacuum

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Structure and microstructure of EB-PVD yttria thin films grown on Si (111)substrate

Mária Hartmanová a,*, Matej Jergel a, Juan Pedro Holgado b, Juan Pedro Espinos b

a Institute of Physics, Slovak Academy of Sciences, 84511 Bratislava, Slovakiab Institute of Material Science, Univ.Sevilla e C.S.I.C., 41092 Sevilla, Spain

a r t i c l e i n f o

Article history:Received 9 March 2010Received in revised form27 August 2010Accepted 7 September 2010

Keywords:Yttrium oxideThin filmElectron beam depositionDefect formationX-ray spectroscopyElectron microscopy

* Corresponding author. Tel.: þ421 2 59410 545; faE-mail address: maria.hartmanova@savba.sk (M. H

0042-207X/$ e see front matter � 2010 Elsevier Ltd.doi:10.1016/j.vacuum.2010.09.003

a b s t r a c t

Structure and microstructure of yttria thin films grown by electron beam physical vapour deposition ona stationary Si (111) substrate at room temperature (RT), 500� and 700 �C, were investigated by thegrazing-incidence X-ray diffraction and scanning electron microscopy, respectively. X-ray photoelectronspectroscopy provided information on the surface contamination from the atmosphere and the Yoxidation state. A strong effect of the deposition temperature and the vapour flux incidence angle wasfound. The film deposited at RT is polycrystalline with very fine grains of the body-centered cubic (bcc)crystallographic symmetry. An increase of deposition temperature results in a rapid growth of bcc grainswith an improved crystalline structure. Moreover, the based-centered monoclinic phase appears for thedeposition temperature of 700 �C. Preferred grain orientation (texture) with two main components,(400) and (622), was observed in the films deposited at 500 �C whereas no texture was found for 700 �C.The microstructure exhibits the columnar feather-like structure of different degrees of perfection whichcan be explained by the shadowing effects caused by an oblique vapour flux incidence angle. Surfacemorphology of the films is governed by a combination of the triangular and four-sided (square) columns.All films were found to be dense with a little porosity between the columns.

� 2010 Elsevier Ltd. All rights reserved.

1. Introduction

Yttria (Y2O3) is one of the most important oxides. It is alsoa fundamental ingredient in other more complex oxides. It hasattracted much attention because of several very interesting phys-ical properties, such as structural stability up to 2325 �C (meltingpoint of Y2O3 is 2450 �C) [1], good mechanical properties [2,3], highdielectric constant k ¼ 14e18 [4,5], high electrical resistivity of1011e1012 Um [2,6], a rather high refractive index near to n ¼ 2 anda very good protective behaviour as a coating in reactive severeenvironment [7]. The relatively high dielectric constant k, highconduction band offset ofe2.3 eV and thermal stability on siliconrender Y2O3 a promising material for electronic applications. Forexample, it can replace SiO2 (k ¼ 3.9) as a gate dielectric in metal-oxide-semiconductor (MOS) heterostructures, in dynamic-random-access-memories (DRAM) and in thin film electroluminescentdevices. Furthermore, Y2O3 may serve as a lattice-matched bufferlayer for epitaxial growth of superconducting oxides [8]. Y2O3 is also

x: þ421 2 5477 6085.artmanová).

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a very importantmaterial for optical applications due to its ability tohost rare-earth atoms such as europium or thulium [9]. Eu e dopedY2O3 is awell known redphosphorus. And also somepapers of KwokCh.-K. et al. [10e12] are devoted to the study of process parameter(preparation condition) e property relationships for depositedyttriumeoxygen system with the special emphasis on opticalproperties.

The real structure of a thin film in terms of the crystallographicperfection, microstructure and surface morphology is basicallydetermined by the deposition conditions. These properties controlthe electrical, optical and mechanical behaviour, hence structuralstudies of thin films are of permanent importance for any applica-tion. Actually, structureeproperty relationship is a very broad topic.In this paper, we concentrate on the structure andmicrostructure inconnection with the preparation condition used and these resultswill provide us a guidance for a future targeted optimization ofchosen properties, in particular electrical conductivity, microhard-ness and elastic modulus. We address the effect of the depositiontemperature on the Y2O3 thin films grown by the electron beamphysical vapour deposition (EB-PVD). The grazing-incidence X-raydiffraction (GI-XRD) and scanning electron microscopy (SEM)techniques were employed to study structure and microstructure,

M. Hartmanová et al. / Vacuum 85 (2010) 535e540536

respectively. Oxidation state of yttrium and surface composition ofthe filmswere determined by the X-ray photoelectron spectroscopy(XPS).

2. Experimental

2.1. Sample preparation (EB-PVD)

Yttrium oxide thin films can be prepared by different methods.Recently, there has been a growing interest in the preparation ofsolid electrolyte films by physical vapour deposition (PVD) whichenables us to control precisely the film characteristics such asmicrostructure, porosity and stoichiometry as well as the growthrate during the deposition [13]. In the present study, the electronbeam PVD (EB-PVD), which is one form of the PVD technique, wasemployed. It allows a high deposition rate, large deposition areaand good adherence of the film to the substrate. A low latticemismatch between the lattice parameters of Si (aSi � 2 ¼ 1.086 nm)and Y2O3 (1.0604 nm) [14] is the next advantage. The Y2O3 thinfilmswere deposited on a stationary Si (111) substrates, whichwerecleaned by HF before the deposition, at three temperatures Tdep e

room temperature (RT), 500 �C and 700 �C. The direction of thevapour flux incident on the stationary substrate is usuallyperpendicular. However, in our case it could be partially obliquewith a maximum deviation of w10� from the surface normal (seethe discussion below). The deposition from the crystalline yttria(Y2O3) target was done at a pressure of 3.3 � 10�1 Pa and a depo-sition rate of 0.1 nm/s at each deposition temperature. The filmsthicknesses as determined from the SEMmicrographs were 300 nm(RT), 530 nm (500 �C) and 280 nm (700 �C), with an accuracyof �10 nm (see Section 3.3).

2.2. X-ray photoelectron spectroscopy (XPS)

The surface properties of the Y2O3 thin films and the oxidationstate of Y were studied by the X-ray photoelectron spectroscopy(XPS). The XPS spectra were obtained on ESCALAB 210 spectrom-eter with a hemispherical electron energy analyzer working in theconstant pass energy mode at 50 eV. Non-monochromatized MgKa

radiation (hn ¼ 1253.6 eV) was used as the excitation source. Wehave performed two analyses of each sample. The first one wasdone on the sample surface after its exposition to the laboratoryatmosphere for several months. The second one was done aftera surface cleaning by the soft ion bombardment (sputtering) withArþ ions of 2.5 keV kinetic energy for 10 min. A sputtering rate ofw0.05 nm/min was estimated from the sputtering of a Ta2O5/Tareference sample (sputtered under the same experimental condi-tions) of well-defined thickness and chemical composition,commonly used in the depth-profiling of solid surfaces and thinfilms for the calibration of sputtering depth [15]. The O1s, Y3d andC1s core levels were identified. As the Y2O3 films are bad electronicconductors, all XPS spectra were calibrated against the position ofY3d5/2 peak at 156.7 eV. The intensities of C1s and O1s werenormalized to the area of the Y3d signal.

2.3. X-ray diffractometry (XRD)

The grazing-incidence X-ray diffraction (GI-XRD) was measuredon a Bruker D8 Discover SSS X-ray diffractometer in the parallelbeam geometry. A generator with the Cu rotating anode provideda quasi-parallel primary beam (divergence 0.03�) of 109 photons/sflux at the output of a Göbel mirror. A long Soller slit with 0.35�

divergence was positioned in the diffracted beam in front of a NaI(Tl) scintillation detector. The grazing-incidence geometry with the

angle of incidence of 1� was applied to avoid the diffraction fromthe Si (111) substrate.

2.4. Scanning electron microscopy (SEM)

The SEM images were obtained by a Hitachi S-5200 microscopywith a field emission filament, using an accelerating voltage of4e5 kV, and extraction current of 10 mA. The samples were cut to2 � 2 mm and fixed to the sample holder with a graphite stickytape. Under the working conditions of microscope, the samples didnot need to be covered with a graphite or gold layer but were putdirectly into the analysis chamber, ensuring a realistic image of thesurface. For the cross-sections, the samples were marked witha diamond tip on the back of the Si substrate and cut by traction justbefore the measurement to ensure a “fresh” cut of the sample.

3. Results and discussion

3.1. Oxidation state and surface properties

In order to investigate the oxidation state of yttrium and thesurface composition of the Y2O3 films, a comparison of the XPSspectra of two Y2O3 samples prepared at Tdep ¼ 700 �C was per-formed, one of them being in the as-deposited state and the otheronewith the surface cleaned by the Arþ ion bombardment (see Sec.2.2) (Fig. 1). For the as-deposited films, the oxidation state of Y isalways þ3, as it is expected for samples exposed to the atmospherefor several months (Fig. 1b). The depth of the valley between theY3d5/2 and Y3d3/2 peaks is different for the as-deposited samplesand the sputtered ones, being deeper for the former. This fact iscaused by an increase of the width of the Y3d signals after thesputtering treatment which is presumably due to the presence ofthe Y3þ ions with different coordination numbers with respect tothe oxide ions since it is expected that ion sputtering generatesoxygen vacancies in the top layer.

As it can be seen in Fig. 1c, the as-deposited films show differentdegrees of contamination, e.g. carbonates are clearly visible in theC1s (Fig. 1c) and O1s (Fig. 1a) regions. However, after the surfacecleaning (Fig. 1aec) the amounts of carbonate species stronglydecrease which proves they are surface contaminants.

3.2. Phase composition and crystalline structure

Yttria has a so-called C-type rare-earth sesquioxide structure.This structure belongs to the bixbyite type [VIA2] [IVO3] with thebody-centered cubic structure, space group Ia-3. It can viewed asbeing derived from the cubic fluorite-type structure (CaF2) byremoving one quarter of the oxygen atoms. This structure has largeinterstitial sites comparable with the size of the oxygen ion in theanion sublattice which are ordered in non-intersecting stringsrunning through the structure and may provide relatively openpathways for the interstitial diffusion [16]. This fact is important forthe electrical conductivity. It is known that Y2O3 is predominantlyp-type conductor at the high temperatures and/or low oxygenpartial pressures. The contribution of ionic conductivity increaseswith decreasing temperature. In such case the charge transport hasa mixed character where the charge carriers are the oxygen inter-stitials, the electron holes and the oxygen vacancies, according tothe given conditions. And the mobility of oxygen interstitials ishigher compared to that of its oxygen vacancies [17]. The impor-tance of oxygen interstitials in undoped yttria is supported also bythe measurements of mechanical hardness [2].

Phase composition and crystalline structure of the films werestudied by GI-XRD. The diffraction patterns were evaluated byPawley fitting with fundamental parameters approach (FPA). The

Fig. 1. aec. The XPS spectra of the as-prepared Y2O3 films (full line) and the Y2O3 filmswith cleaned surface (dotted line) deposited on Si (111) substrate at Tdep ¼ 700 �C.Different plots refer to different energy ranges.

Fig. 2. aec. The GI-XRD patterns of the Y2O3 films deposited at different temperatures(a) Tdep ¼ RT, (b) Tdep ¼ 500 �C, (c) Tdep ¼ 700 �C; all diffractions for RT and 500 �Cbelong to the body-centered cubic (bcc) phase; the diffractions in (c) indicated by thearrows belong to the monoclinic phase.

M. Hartmanová et al. / Vacuum 85 (2010) 535e540 537

FPA parameters were determined from a measurement ofa corundum NIST standard under the same experimental condi-tions as the films themselves. In addition to the refined latticeparameters (aref, bref, cref, bref), microstructure was determined interms of the grain size and root-mean square (rms) value ofmicrostrain using the double-Voigt approach. The results areshown in Fig. 2 and Table 1. The ICDD database PDF-2 was used forthe reference lattice parameters (aPDF, bPDF, cPDF, bPDF).

The Y2O3 film deposited at room temperature (Fig. 2a) is poly-crystalline with very fine grains of bcc symmetry and largemicrostrain 30 characterizing random lattice distortions. A z 1%lattice expansion in comparison with the reference phase can be

inferred. With increasing deposition temperature, the grains of thebcc phase grow abruptly, namely z 4� and z16� for 500 �C(Fig. 2b) and 700 �C (Fig. 2c), respectively. Simultaneously, themicrostrain 30 in the grains drops by nearly one order of magnitudefor 500 �C and 700 �C and the lattice expansion/contractiondiminishes to 0.01% which approaches the precision of the deter-mination of the lattice parameters. In addition to the bcc phase,some diffraction peaks appear for 700 �Cwhich can be attributed tothe base-centered monoclinic phase (Fig. 2c). As not all diffractionpeaks of this forming phase emerged from the background, themicrostructure could not be determined completely. In particular,the grain size given in Table 1 represents a minimum estimate

Table 1Deposition temperature Tdep, phase composition, space group, grain size D, microstrain 30 (rms value), refined lattice parameters aref, bref, cref, bref and database latticeparameters aPDF, bPDF, cPDF and bPDF of Y2O3 films.

Tdep (C�) phase space group D (nm) 30 (%) aref (�A) bref (�A) cref (�A) bref (�)

apdf (�A) bpdf (�A) cpdf (�A) bpdf (�)

RT body-centered cubic Ia-3 (206) 6.9 0.612 10.703 e e e

10.604 e e e

500 body-centered cubic Ia-3 (206) 26.8 0.088 10.611 e e e

10.604 e e e

700 body-centered cubic Ia-3 (206) 113.3 0.069 10.592 e e e

10.604 e e e

base-centered monoclinic C2/m (12) 41.2 0.000 13.9120 3.4396 8.6492 100.3513.8992 3.4934 8.6118 100.27

M. Hartmanová et al. / Vacuum 85 (2010) 535e540538

when the microstrain 30 is neglected. The rapid growth of the bccgrains and improvement of their crystalline structure with thedeposition temperature is thus the most prominent featurerevealed by GI-XRD.

A comparison of the measured relative peak intensities withthose in the database suggests a random grain orientation for700 �C (Fig. 2c) while a preferred grain orientation (texture) withtwo main components may be inferred from the diffraction patternfor 500 �C (Fig. 2b), where the 400 and 622 diffractions peaksdominate. This texture evolved presumably due to the large filmthickness which was nearly 2� larger than that for 700 �C (seeSection 2.1.). The results obtained are in a good agreement with theliterature data [18e20].

3.3. Defects influencing the oxide ionic conductivity

The ionic conductivity in ionic and mixed oxide conductorsdepends strongly on the crystal structure. Geometric informationon the ionic conductivity and diffusion is a key to understand theconduction mechanism. The determination of the diffusion path-ways of mobile ions in many ionic and mixed conductors which areinherently connected with the structural defects such as grainboundaries in polycrystalline conductors, is a complex problem[21]. Previous studies of the oxide ceramics indicate that ionmobility is extremely sensitive to the character of grain boundaries.The exact configuration of grain boundaries is problematic toquantify in practice and their influence on the electrical/dielecticproperties is difficult to distinguish from the effect of other defectssuch as porosity, oxygen vacancies, impurities and dislocations[22]. As it was noted by M. Tsuchiya and S. Ramanathan [23], thecationmobility is primarily rate-controlling step for grain boundarymigration at nearly all compositions in the oxides with fluorite-type and related structures [23,24] which is also a case of Y2O3. Thegrain boundaries can act as the sinks for the oxygen vacancies V ::

O(in Kr}oger-Vink notation) due to their small formation energy [23].In pure Y2O3, the grain growth is typically controlled by themobility of Y3þ cations which diffuse via the vacancy mechanismðV 000

Y Þ or ðV ::OV

000Y Þ0 [23]. Moreover, due to C-type bixbyite structure, the

contribution of easymobile yttrium interstitials is also important inY2O3 [16].

The formation of grain boundaries is determined by the graingrowth. In the case of doped oxides, dopant segregation due toa solute-drag effect [25] and a minimization of the size-inducedstrain between the host and dopant cations play an important rolein the grain growth. Moreover, the grain growth can be modified bythe irradiation treatment during the film deposition as it wasshown for the case of yttria-stabilized zirconia in [23].

The abovementioned factors affect crucially electrical propertiesof the oxide films. For instance, the electrical conductivity of the EB-PVD grown films exhibits characteristic anisotropic behaviour [22].

Here, theelectrical conductivity in thedirectionperpendicular to thefilm surface is remarkably higher than that in the direction parallelto it which is attributedmainly to the columnar structure typical forEB-PVD films and the existence of many grain boundaries along thedirection parallel to the film surface.

3.4. Surface and cross-sectional morphology/microstructure

Development of the EB-PVD thin film microstructure is essen-tially determined by four basic processes [26] namely shadowing,surface diffusion, volume diffusion and desorption. They arestrongly influenced by the deposition conditions such as substratetemperature and roughness, rotation speed of the substrate, thevapour flux incidence angle, energy of the vapour particles,chamber pressure, deposition rate.

Under stationary conditions (without rotation of Si (111)substrate), the cross-sectional SEM micrographs of the Y2O3 filmsexhibit the columnar structure with different degree of perfectiondepending on the preparation conditions and the film thickness(Fig. 3). The interface between the Si substrate and the Y2O3 layercan be distinguished clearly which indicates that the interface issharp at the microscopic level with no distinct interlayer even atelevated temperatures. It is known that formation of amorphousSiO2 interfacial layers or silicides such as YSi2 is possible by reac-tions between Y2O3 and silicon substrate during the depositionand/or post-deposition thermal treatment [16]. Due to a lowerpermittivity, such interfacial layer is usually unwanted. In our case,it is obviously very thinwhich together with a low scattering powerof SiO2 explains the absence of the SiO2 diffraction peaks in the GI-XRD pattern.

The columnar microstructure is best visible in the thickest filmdeposited at 500 �C (Fig. 3b), where the strongest texture wasdetected by the GI-XRD, while it is only partly developed at RT(Fig. 3a). The collinear elongated columns are actually elongatedgrains. The average width of the columns increases in the growthdirection from 35 nm to 75 nm for Tdep ¼ 500 �C and from 80 nm to140 nm for Tdep ¼ 700 �C. The difference between the grain sizedetermined by the GI-XRD and SEM is due to the fact that the GI-XRD provides the size of the coherently scattering domains capableof constructive interference in the scattering vector direction whilethe columns visible by SEM are not that structurally perfect,especially close to the boundaries.

The columnar microstructure affects the surface morphology. Amixture of flat triangular (three-sided) grains and grains with therooftopmorphologywas observed on the stationary deposited films[26,27]. In our case the film surface shows a combination of trian-gular columnswith the tips consisting of three triangular facets andsquare (four-sided) columns terminated by four facets (Fig. 3b,c),where the latter are typical for the rotated samples. This fact can beexplained by an oblique incidence angle of the vapour flux on the

Fig. 3. aec. The SEM micrographs of the surfaces (upper) and cross-sections (lower) of the Y2O3 film deposited at different temperatures: (a) Tdep ¼ RT, (b) Tdep ¼ 500 �C, (c)Tdep ¼ 700 �C.

M. Hartmanová et al. / Vacuum 85 (2010) 535e540 539

stationary substratewhichwas reported to give rise to a feather-likeappearance of the columnar structure [26]. Such a feather-likemicrostructure is indeed visible in Fig. 3b,c and is unusual for theconventional stationary EB-PVD deposited films. It evolves due to

the shadowing effect of the columns when the vapour flux is notperpendicular but inclined with respect to the stationary substrate.Since EB-PVD is the line-of-sight deposition process, the substrateadjustment controlling the vapour incidence angle and thus the

M. Hartmanová et al. / Vacuum 85 (2010) 535e540540

shadowing effect is one of the important factors for the micro-structure control. It is known that microstructure and texture of theEB-PVD films are generally strongly influenced by the shadowing ofthe vapour flux by columnar grains. Due to the intrinsic character-istics of EB-PVD, non-equilibrium condensation of the vapour phasein presence of the shadowing effect leads to the development of themicroporous structure [13] which affects strongly the film proper-ties such as electrical conductivity and gas tightness. In comparisonwith the rotationally deposited films, the stationary deposited onesexhibit less porosity between the columns and consequentlya higher density [26]. Indeed, little porosity between the columns isobservable in our films (Fig. 3b,c).

4. Summary

The investigation of the structure and microstructure of theyttria EB-PVD thin films grown on the stationary Si (111) substraterevealed the following effects of the deposition temperature andthe vapour flux incidence angle:

1. Structure of the Y2O3 films deposited at RT, 500� and 700 �C ispolycrystalline with the bcc symmetry; some diffraction peakswhich can be attributed to the base-centeredmonoclinic phaseappear for Tdep ¼ 700 �C

2. Increase of the deposition temperature results in a rapidgrowth of the bcc grain and improvement of their crystallinestructure; the preferred grain orientation (texture) with twomain (400) and (622) components evolves in the thickest filmdeposited at 500 �C whereas a random grain orientation ispresent in the thinner film deposited at 700 �C which indicatesthe role of the film thickness in the texture build-up

3. Microstructure of all yttria films is governed by columns withdifferent degrees of perfection depending on the depositiontemperature; surface morphology is controlled by triangularcolumns with a three-sided column tip and four-sided squarecolumns terminated by four facets

4. Feather-like structure of columns, clearly observable for thefilm deposited at 500 �C, could be explained by the shadowingeffect due to the oblique incidence angle of the vapour fluxwith respect to the stationary substrate

5. All Y2O3 films were found to be dense with little porositybetween columns.

Acknowledgements

The work was partially supported by the research grants No. 2/0053/10 and 2/0047/08 of Slovak Grant Agency (VEGA), Project No.2005 SK0001, Slovak Academy of Sciences, Bratislava, Slovakia e

C.S.I.C. Sevilla, Spain (y. 2006e2007) and SAV-FM-EHP-2008-01-01project. The authors would like to thank Dr. �S. Chromik from theInstitute of Electrical Engineering, SAS, Bratislava, for the prepara-tion of yttria thin films.

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