Structural and optical properties of polycrystalline thin films of rare earth oxides grown on fused quartz by low pressure MOCVD

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Journal of Crystal Growth 276 (2005) 148–157 www.elsevier.com/locate/jcrysgro

Structural and optical properties of polycrystalline thin films of rare earth oxides grown on fused quartz by low pressure MOCVD M.P. Singh, S.A. Shivashankar Materials Research Centre, Indian Institute of Science, Bangalore-560 012, India Received 3 September 2004; accepted 8 November 2004 Communicated by M. Kawasaki

Abstract We report the structural and optical properties of oriented polycrystalline thin films of rare earth oxides (REO), namely Er2O3, Gd2O3, Eu2O3, and Yb2O3 grown on fused quartz by low-pressure metalorganic chemical vapour deposition (MOCVD) in the temperature range of 450–800 1C, using adducted b-diketonate complexes of rare earth metals as precursors. The films were characterized by X-ray diffraction (XRD) and UV–visible spectrophotometry. While the films grown at low temperatures (
500 1C) are poorly crystalline, those grown on or above 525 1C display a significant (1 1 1) texture. Growth of textured cubic REO in (1 1 1) direction on amorphous substrate is interpreted in terms of minimization of surface energy at the film–substrate interface. The degree of misfit between the interatomic distances in the disordered fused quartz substrate and those in the crystalline REO is invoked to explain the presence (absence) of low-intensity reflections other than (1 1 1). The incorporation of carbon (from the precursor) into the films affects the optical properties of REO, as examined by UV–visible spectroscopy. r 2004 Elsevier B.V. All rights reserved. PACS: 68.55.Jk; 81.15.Gh; 77.84.Bw Keywords: A1. Disordered substrate; A3. Metalorganic chemical vapor deposition; A3. Texture; A3. Thin films; B1. Rare earth oxides

1. Introduction

Corresponding author. Tel.: +91 80 2293 2782; fax: +91 80 2360 0683. E-mail address: [email protected] (S.A. Shivashankar).

The oxides of rare earth metals exhibit interesting physical properties, such as high transparency from the ultraviolet to the infrared, high refractive index, good dielectric and insulating properties because of a relatively high dielectric constant

0022-0248/$ - see front matter r 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2004.11.325

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(
12–14), low dielectric losses, as well as high dielectric breakdown field strength [1–3]. Therefore, thin films and coatings of rare earth oxides (REO) have potential for application in various fields of science and technology, such as optical devices, semiconductor devices, and telecommunication. Recently, REO films have also been studied as an alternative gate dielectric to replace SiO2 in complementary-metal–oxide–semiconductor (CMOS) devices of the forthcoming generations [4–6]. This possibility has prompted the synthesis and characterization of thin films of REO and studies of their properties [1–6]. In principle, either a physical vapour deposition (PVD) or a chemical vapour deposition (CVD) may be employed to deposit the films of REO. To date, thin films of REO have generally been deposited by a PVD process such as thermal evaporation and electron beam evaporation of the oxide powder, on a variety of substrates, including single crystal silicon, fused quartz, gallium arsenide, and glass [1,4–6]. Generally, these films have been found to be amorphous or polycrystalline in nature. In some cases, oriented REO films with preferred crystallographic orientation have also been deposited by PVD. It has been found that any preferred orientation in the films depends on the deposition technique and the processing parameters. Recently, efforts have also been made to deposit epitaxial films of rare earth oxides, Gd2O3 in particular, on single crystal silicon and GaAs substrates [6]. The lattice match between Si and Gd2O3 is excellent. This presents the prospect of an epitaxially grown high-k oxide as the gate dielectric in CMOS devices. It is well known that properties of thin films depend on their structure and composition, which are governed directly by the growth process and the substrates employed. Of the various film deposition processes, metalorganic chemical vapour deposition (MOCVD) has become an attractive process for the deposition of thin films of oxides and other coatings; because the technique offers uniform growth over large areas at relatively low growth temperatures and the conformal coverage of high aspect ratio features. In such a chemical process for thin film deposition, the composition and structure of the films are

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determined by the chemical precursor used, deposition conditions employed, and the nature of the substrate. However, MOCVD process can lead to the incorporation of heteroatoms (i.e., C and H) into the film, which derives from the precursor itself [7,8]. For this reason, perhaps, there are only a few reports on the deposition of REO films by the MOCVD process [9–11]. This may also be due to the unavailability of suitable CVD precursors. Thus, very little information is available on the nucleation and growth of REO films on amorphous and single crystal substrates by low-pressure MOCVD and on the properties of MOCVD-grown REO films. It may also be noted that, when REO films are grown by MOCVD on Si(1 0 0), the growth surface is likely to be SiO2. This is because a native oxide layer (which is amorphous) may be formed on Si due to the presence of oxygen in the molecules of MOCVD precursors such as b-diketonates. Hence, a study of REO films grown on fused quartz substrates by the MOCVD process is warranted. Further, an appreciation of the factors that determine the growth in MOCVD is essential in understanding the structure and properties of the resulting films and the performance of devices employing them. In this paper, we report an investigation of the structural and optical properties of MOCVDgrown REO films on fused quartz substrate—an amorphous surface, using (powder) X-ray diffraction (XRD) and UV–visible spectrophotometry. The effects on the optical properties of the incorporation of heteroatoms into the REO films have also been studied.

2. Experimental procedure Thin films of the various REO were grown on 15 mm  15 mm fused quartz substrates (polished), using the (1,10-phenanthroline) adduct of the acetylacetonate (acac) complex of the different rare earth metals, namely Er, Gd, Eu, and Yb. These adducted complexes, which are subliming solids may be designated as M(acac)3.Phen, where (M¼Er, Gd, Eu, Yb), were synthesized and characterized in house. It may be noted that such complexes have direct metal–oxygen bonds,

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making them suitable for the CVD metal oxide films. Their thermal characteristics, i.e., their vapourization and decomposition behaviour, were analysed by simultaneous thermogravimetry/differential thermal analysis (TG/DTA). Films were prepared in a home-made, horizontal, hot-wall, low-pressure MOCVD reactor. Films were deposited at temperatures ranging 450–800 1C, using ultrahigh purity argon (flow rate 20–150 sccm) and oxygen (50–500 sccm) as carrier and reactive gas, respectively. Gas flow was regulated using electronic mass flow controllers. A capacitance manometer was employed for monitoring the total reactor pressure, which was maintained at 2 Torr. The precursor, which is a crystalline solid at room temperature, was taken in a finely powdered form, and placed in a glass boat inside a vapourizer maintained in the range of 150–250 1C, depending on the rare earth metal. Prior to deposition, the fused quartz substrates were cleaned ultrasonically in different solvents, and subsequently boiled in acetone, methanol, and trichloroethylene. The thickness of the films was measured by stylus profilometry. The surface morphology of the films was examined by scanning electron microscopy (SEM). Powder X-ray diffractometry was employed to characterize the crystallinity and the texture of the films. A UV–visible spectrophotometer in the transmission mode (wavelength range of 190–800 nm) was employed to characterize the optical properties of the films.

3. Results and discussion 3.1. Experimental results of X-ray diffraction The surfaces of the as-grown REO films were found to be uniform, pore-free, and essentially featureless, as examined by SEM. Films deposited at relatively higher growth temperatures (
550 1C) were transparent, whereas films grown at lower temperatures (
500 1C) were blackish in colour. The blackish appearance is probably due to the presence of carbon in the film, deriving from the precursor itself [7]. Irrespective of the growth temperature and the REO, films were found to adhere to the substrate well, as revealed by the

adhesive tape peel test. The thickness of the films was in the range 100–800 nm (as measured by stylus profilometry), depending on the REO and the CVD conditions employed. Fig. 1 shows the powder XRD patterns of erbium oxide films grown at various temperatures. The patterns were indexed on the basis of the standard powder pattern for polycrystalline erbium oxide (JCPDS card no. 43-1007). The XRD patterns show that the films comprise the cubic phase of Er2O3 (c-Er2O3) at all temperatures of growth. Fig. 1a reveals that the Er2O3 film grown at 500 1C have two broad XRD peaks of low intensity, which is indicative of poor crystallinity. But XRD pattern of the Er2O3 film grown at 525 1C (Fig. 1b) reveals clearly that film crystallinity has improved. At temperatures above 550 1C, films grow preferentially with the (2 2 2) crystallographic orientation, at the expense of growth along the (2 1 1) direction, as seen in Figs. 1c and d. Thus, the XRD patterns of Er2O3 films establish clearly that films grown at lower temperatures are poorly crystalline, whereas films grown at higher temperatures tend to have the (1 1 1) orientation. Despite the developing (1 1 1) texture, ‘‘minor’’ peaks corresponding to other crystallographic orientations are present. But the intensity of these peaks is significantly less than that of the (2 2 2) peak of erbium oxide. As the growth temperature is raised, the relative intensity of the (2 2 2) peak increases, indicating a stronger tendency for the (1 1 1) texture. For example, the intensity ratio of (2 2 2) and (4 0 0) reflections in the film grown at 700 1C is 40, in contrast with a ratio of 3 in polycrystalline Er2O3. Thus, thin films of (cubic) Er2O3 on fused quartz, grown by MOCVD from a b-diketonate precursor, have a significant (1 1 1) texture. A similar study was carried out on gadolinium oxide films. Fig. 2 shows the powder XRD patterns of films of gadolinium oxide deposited on fused quartz grown at various temperatures from an adducted acetylacetonate complex. Identification of XRD peaks is based on the standard powder pattern for polycrystalline gadolinium oxide (JCPDS card nos. 43-1014, 43-1015), revealing that the films comprise the cubic phase of Gd2O3 (c-Gd2O3) at all temperatures of growth

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(222)

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0 20

40

60 2θ

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80

(d)

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60

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Fig. 1. XRD patterns of thin films of Er2O3 grown at various temperatures; the intensities of the different patterns are on the same relative scale.

750°C

550

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M(-401)

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0 (b)

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40 (d) 30 20 10 0 1100 (c)

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525°C

0 20

30

40

50

60

70

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2θ Fig. 2. XRD patterns of thin films of Gd2O3 grown at various temperatures; the intensities of the different patterns are on the same relative scale. ‘‘C’’ stands for cubic and ‘‘M’’ stands for monoclinic.

employed in this study. However, at the higher growth temperatures, the films comprise also the monoclinic phase, as indicated by the low intensity peaks corresponding to it. Fig. 2a shows that the film grown at 525 1C is more crystalline, with a preference for the (1 1 1) orientation. Films grown at temperatures higher than 525 1C show a stronger preference for the (1 1 1) crystallographic orientation, as seen in Fig. 2. Thus, the XRD patterns of Gd2O3 films establish that even films grown at relatively low temperatures are crystalline, while those grown at higher temperatures tend to have the (1 1 1) texture. This tendency increases up to a growth temperature of 600 1C, beyond which the monoclinic phase begins to form, together with the cubic phase. With increasing growth temperature, it is found that the intensity of the peaks due to the monoclinic phase increases. This is consistent with the (binary) phase diagram for ceramic Gd2O3, which indicates that the phase stable at lower temperature is cubic

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3.2. Growth mechanism of oriented thin films of REO It is noteworthy that the fused quartz is amorphous in nature, i.e., it has no long-range ordering. Strongly oriented growth of REO films on fused quartz may be understood on the basis of the minimization of surface energy of the interface through such growth [13], and verified experimentally for thin oxide films [14–17]. In such a growth process, i.e., growth of oriented film on an amorphous substrate; that orientation of the crystalline film is preferred which has the highest atomic density. This happens to be the (1 1 1) orientation for cubic materials. It is also possible that such oriented growth is aided by the presence of M–O bonds in the precursor. It may also be promoted by the presence of oxygen in the substrate, as indicated by the strong adhesion between the films and substrate. As discussed above, the minimization of surface energy at the interface plays a dominant role in establishing the textured growth of the cubic phase

200

C(444)

Intensity

whereas, at higher temperatures the phase formed is monoclinic [12]. Fig. 2d shows the XRD pattern of a film grown at 750 1C, which clearly indicates that the film comprises both the monoclinic and the cubic phases. The intensity of the (2 2 2) reflection due to the cubic phase is larger than that of the (4¯ 0 1) reflection corresponding to the monoclinic phase. This indicates that even the film grown at 750 1C has a (1 1 1) texture for the cubic phase present in it. (The peak intensities in this pattern are smaller than in the other three because of the smaller film thickness, the deposition rate decreases as the temperature is raised.) Thus, thin films of Gd2O3 on fused quartz, grown by MOCVD from an adducted b-diketonate precursor, have a (1 1 1) texture. Similar trends were observed in the case of europium oxide and ytterbium oxide films. Fig. 3 illustrates the XRD pattern of Eu2O3 and Yb2O3 films grown at 600 1C, revealing clearly that the films have a significant (1 1 1) texture. Thus, we conclude that, irrespective of the REO, films grown on fused quartz by the low-pressure MOCVD process have the (1 1 1) texture.

Intensity

152

0 20

(b)

30

60

2θ

Fig. 3. XRD patterns of thin films of (a) Eu2O3 and (b) Yb2O3 grown at 600 1C; the intensities of the different patterns are on the same relative scale.

of these oxides on a smooth amorphous substrate. There are other factors as well, such as misfit between film and substrate, surface roughness in the substrate, etc., which play a role in the film growth process and the development of the microstructure [13]. Even though the MOCVDgrown REO films display a prominent (1 1 1) texture, the presence of relatively weak reflections due to other orientations of the cubic phase of REO depends on the specific oxide and the CVD conditions. A detailed examination of the XRD data for Er2O3 films and Gd2O3 films reveals that, in the case of cubic gadolinium oxide (c-Gd2O3) films, weak XRD peaks corresponding to other

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crystallographic directions of c-Gd2O3, such as (4 0 0) and (6 2 2), effectively disappear when the growth temperature is raised. However, in the case of Er2O3 films, the (4 0 0) and (6 2 2) peaks, though weak, are present irrespective of the growth temperature. Furthermore, the XRD data show that Eu2O3 films deposited on fused quartz behave in a manner similar to Gd2O3 films on fused quartz, as far as the temperature dependence of the peak intensities is concerned. It may thus be inferred from the present work that, in REO films grown under identical CVD conditions, the extent of deviation from the (1 1 1) orientation depends on the specific oxide material, whereas the propensity for the (1 1 1) texture of the cubic phase in the REO film (even when a second oxide phase is present) is independent of the identity of the REO. The presence of these weak XRD peaks in Er2O3 films [e.g., (4 0 0), (6 2 2)], corresponding to the other allowed reflections due to the cubic phase, and the absence of such peaks in the XRD patterns of the other REO, may be understood on the basis of interfacial misfit between the lattice of the specific REO and of fused quartz. To interpret the growth process on a fused quartz (SiO2) substrate and to estimate the misfit between the REO lattice and the substrate, an understanding of the Si and O arrangements in crystalline quartz and in its amorphous form is necessary. A detailed structural study of SiO2 reveals that the arrangement of SiO4 tetrahedrons oriented in a particular way is the basis for all the structural modifications of silicon dioxide (cristobalite, tridymite, and others) [18]. In fused quartz, this arrangement is nearly coincident with that in crystalline quartz [18]. Each atom of silicon is enclosed in a

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tetrahedron formed by four atoms of oxygen, whereas each oxygen atom is linked with two silicon atoms, forming a flexible bond, which exhibits some spread in the Si–O–Si angle, a; ranging from 120–1801, and having the maximum at a ¼ 1441: The possibility of the variability in the bond angle a in fused (amorphous) quartz results in some elasticity of the Si–Si distance which, for a ¼ 1441; is about 3.1 A˚. The interatomic distances for the different crystallographic forms of SiO2 and its amorphous modification are given in Table 1. The growth of oriented thin films on a crystalline substrate is primarily determined by the lattice mismatch between the film and substrate. Such interfacial misfit between film and substrate is quantitatively expressed by the formula   ai  bi %misfit ¼ 2 ; (1) ai þ bi where ai and bi are the lattice parameters of film and substrate, respectively. If the value of interfacial misfit is very close to zero (a few percent), one expects strongly oriented (epitaxial) growth of films on that substrate. This concept can be extended and employed to understand the growth of strongly oriented thin films on amorphous substrates, where the texture of the films is governed by the minimization of surface energy. In this regard, Andreeva and Kasumov have made an attempt to understand the growth of vacuum-evaporated REO films on fused quartz [19]. To understand the observed variations in the texture of such REO films deposited on fused quartz, these authors have made an effort to estimate the misfit between the REO lattice and fused quartz. The misfit between

Table 1 The interatomic distances (A˚) for crystalline, amorphous fused quartz, and amorphous SiO2 films (adapted from Refs. [18,19]) Bonds

a-Cristobalite

b-Tridymite

Amorphous fused quartz

Amorphous SiO2 films obtained by sputtering

Si1–O1 O1–O2 Si1–Si2 Si1–O2 O1–O3

1.54 2.52 3.08 3.88 4.36

1.54 2.52 3.08 3.88 4.36

1.62 2.65 3.10 4.00 4.50

1.60 2.65 3.10 3.88 4.36

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the interatomic distances in the crystal lattices of the different REOs lattices and the interatomic distances in SiO2 (fused quartz) nearest to them was estimated by the equation   ai  bi %misfit ¼ 2 ; ai þ bi

(2)

where ai is the interatomic distance in the ith direction in the REO lattice, and bi is the nearest interatomic distance in SiO2. The results of this estimation for several REOs are given in Table 2. The data in Table 2 shows clearly that the interfacial misfit between fused quartz and erbium oxide in the (2 2 2), (4 0 0), and (6 2 2) crystallographic direction of c-Er2O3 are 1.9%, 0.6%, and 0.6%, respectively. This indicates that, although surface energy minimization would favour (1 1 1) texture in c-Er2O3 films, the smaller interfacial misfit would favour the growth of c-Er2O3 crystallites bearing the (4 0 0) and (6 2 2) crystallographic orientation. Thus, in practice, even though the CVD-grown film has (1 1 1) texture, crystallites of Er2O3 with the (4 0 0) and (6 2 2) orientations are also present, because of their interfacial misfit being significantly smaller than for the (2 2 2) direction. However, in the case of Gd2O3, the interfacial misfit between Gd2O3 and fused quartz in the (2 2 2), (4 0 0), and (6 2 2) crystallographic directions of c-Gd2O3 are 0.5%, 1.7%, and 1.7%, respectively. This clearly indicates that surface energy minimization, and interfacial misfit, together favour the growth of Gd2O3 films with the (1 1 1) texture. This is the reason why, in the case of Gd2O3

films, the weaker XRD peaks, corresponding to (4 0 0) and (6 2 2) crystallographic planes, disappear as the substrate temperature is increased. By contrast, in the case of Er2O3, the (4 0 0) and (6 2 2) peaks continue to be present even when the growth temperature is raised. That is, the greater mobility of the growth species at a higher growth temperature is not sufficient to ‘‘eliminate’’ the weaker reflections and to give exclusively the texture dictated by surface energy considerations alone. Furthermore, in the case of Eu2O3, the percent interfacial misfit between the c-Eu2O3 lattice and fused quartz in the (2 2 2), (4 0 0), and (6 2 2) crystallographic directions is 1.2, 2.0, and 2.3, respectively. Thus, here again, the tendency for surface energy minimization, along with the low misfit between the c-Eu2O3 lattice and fused quartz would favour the (1 1 1) texture for films of europium sesquioxide. The considerably larger misfit between Eu2O3 and fused quartz in other directions i.e., (4 0 0) and (6 2 2), is a much weaker tendency for Eu2O3 crystallites to grow with these particular orientations. To verify the validity of this analysis further, the misfits between cubic ytterbium oxide (c-Yb2O3) and cubic samarium oxide (c-Sm2O3) and fused quartz were calculated (Table 2). Subsequently, these results were tallied with the XRD patterns of Yb2O3 and Sm2O3 films grown by MOCVD under identical deposition conditions. The CVD precursor used for the deposition of thin films of Sm2O3 was Sm(acac)3.Phen, which is exactly analogous to those used in the present study. Fig. 4 shows the XRD patterns of Sm2O3 film grown on fused

Table 2 %misfit between the interatomic distances in different directions in cubic REO lattices and the nearest onces in fused quartz Material

/2 2 2S Si1–Si2

/4 0 0S O1–O2

/6 2 2S Si1–O2

Er2O3 Gd2O3 Eu2O3 Yb2O3 Sm2O3

1.9 0.5 1.2 0.6 1.7

0.6 1.7 2.0 1.1 3.7

0.6 1.7 2.3 1.2 3.0

Fig. 4. XRD patterns of Sm2O3 films grown at different temperatures (from Ref. [20]).

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quartz at nearly the same temperature as employed in the present work [20]. These XRD data reveal clearly that both oxides are strongly oriented in the (1 1 1) crystallographic direction and that there are no ‘‘minor’’ XRD peaks present, e.g., due to the (4 0 0) plane. Thus, the data on Yb2O3 and Sm2O3 clearly support the hypothesis that surface energy minimization, together with interfacial misfit, decides the ‘‘strength’’ of the crystallographic texture, as well as the presence of weaker reflections due to the non-favoured crystallographic orientations. It is, however, to be noted that, apart from these two factors, i.e., surface energy minimization and interfacial misfit with the substrate, there are other parameters, such as impurity incorporation, the deposition technique employed, etc. which may also be important in the determination of the orientation of a film and its resulting microstructure [13]. Hence, based on the present study, it may be concluded that the crystal structure and the microstructure of MOCVD-grown REO films depends on the crystal structure of the specific oxide. In particular, the degree of preferred orientation of a cubic REO grown on amorphous SiO2 surface depends on the exact lattice parameter of the REO. In this context, it is to be pointed out that the need for an elevated substrate temperature in a CVD process renders texturing due to surface energy minimization more probable than in a PVD process, wherein the substrate temperature need not be high.

50

3.3. Optical properties Optical properties of the REO films were studied by UV–visible spectrophotometry (transmission mode) in the wavelength range of 190–800 nm. Fig. 5 shows transmittance spectrum in UV–visible range for the Er2O3 films grown at 525 and 600 1C. What is noteworthy is that the Er2O3 films grown at 525 1C was blackish in appearance—ostensibly due to the presence of carbon, deriving from the precursor itself—as clearly seen from the optical transmittance spectrum of the film grown at 525 1C, in which the maximum transmittance is only
10%. By contrast, the maximum transmittance in the UV–visible range of the films grown above 525 1C is as high as 90%, which indicates clearly that the film is carbon-free. The transmittance data show, furthermore, that there is a sharp absorption edge below 250 nm. In addition, the transmission spectra of erbium oxide films are composed of a series of relatively sharp features in the visible region. The sharp bands may be assigned to the intraconfigurational f–f transition from the 4I15/2 ground state to the 4F9/2, 2H11/2, 2G9/2, and 2G11/2 excited states, while sharp absorption edge below 250 nm represents the intrinsic band absorption edge. Further, the transmission spectra of the carboncontaining erbium oxide film (Fig. 5a) is characterized by a weaker f–f transition, which is due to the low concentration of erbium oxide nanocrystalline grains in the film. This is because the

100 (i): 4I15/2->4F9/2

40 % Transmission

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Fig. 5. UV–visible transmittance versus wavelength of as-grown Er2O3 film at (a) 525 1C and (b) 600 1C. Inset shows various f–f transitions.

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presence of carbon alters the concentration and local coordination of the erbium oxide crystallites present in the film [21,22]. A detailed study of the UV–visible transmission spectra of the erbium oxide films grown at various temperatures reveals that nature of f–f transition depends on the growth temperature (e.g. films grown at 525 1C shows a sharp absorption band due to the transition from the ground state 4I15/2 to 4F9/2, whereas this particular f–f transition is absent in the films grown above 525 1C). This is because the growth temperature affects the composition and structure of the film directly, thus altering the optical characteristics of the erbium oxide films. To examine band-to-band transitions, the measured absorption coefficient ðaÞ may be used, to estimate the direct and indirect transition energies from the plots of ½aðhnÞ n versus photon energy ðhnÞ: The value of the exponent, n; depends on the nature of the particular transition [23]. For a direct transition n ¼ 12 whereas, for indirect transitions, n ¼ 2: The direct bandgap is calculated from the plot of ½aðhnÞ 0:5 versus the photon energy ðhnÞ following the relation: aðhnÞ ¼ A ½hn  E g 0:5 ;

(3)

where A is a constant and E g is the bandgap of the material. The bandgap of the strongly oriented Er2O3, Gd2O3, and Eu2O3 films, as estimated using Eq. (3), were 5.7, 5.4, and 4.4 eV, respectively. A detailed UV–visible study shows that the optical properties and bandgap of REO films depend on the CVD growth condition. For example, the bandgap of the less-strongly oriented polycrystalline c-Gd2O3 film, i.e., the films grown at lower temperatures, is in the range of 5.1–5.4 eV; whereas the bandgap of the mixed phase Gd2O3 films, i.e., the film containing both the cubic and monoclinic phases, is
5.0 eV, which is smaller than that of the various c-Gd2O3 films. This observed variation in the bandgap of the less oriented, polycrystalline Gd2O3 films may be attributed to the differences in the crystallinity, grain-size, and morphology of the films, which are strongly dependent on the CVD conditions. These results compare with the bandgap of 5.3 eV reported for thin gadolinium oxide films grown

by PVD [24]. The difference may, in part, arise from the microstructural features of the mixedphase film.

4. Conclusions In the present study, we have successfully synthesized polycrystalline thin films of the oxides of several rare earths by low-pressure MOCVD, on fused quartz substrates. Irrespective of the growth temperatures, films comprise the cubic phase of REO. Films grown at elevated temperatures were found to be (1 1 1)-textured, despite the amorphous nature of the substrate. This may be understood on the basis of the minimization of surface energy at the substrate–film interface. The presence (absence) of low-intensity reflections other than (1 1 1) has been explained qualitatively on the basis of degree of misfit between the interatomic distances in the disordered fused quartz substrate and those in the crystalline REO film. Optical measurements on the films reveal that their optical properties are affected by the presence of carbon incorporated into the films from the metalorganic precursor itself.

Acknowledgements The authors thank K. Shalini for providing the metalorganic precursors. MPS thanks the Council of Scientific and Industrial Research (CSIR), New Delhi, India, for a senior research fellowship.

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