- Email: [email protected]
Structure, microstructure and optical properties of cerium oxide thin films prepared by electron beam evaporation assisted with ion beams
Solid State Sciences 11 (2009) 1456–1464
Contents lists available at ScienceDirect
Solid State Sciences journal homepage: www.elsevier.com/locate/ssscie
Structure, microstructure and optical properties of cerium oxide thin films prepared by electron beam evaporation assisted with ion beams Catalina Mansilla* Instituto de Ciencia de Materiales de Sevilla, CSIC-US, Avda. Ame´rico Vespucio 49, 41092-Sevilla, Spain
a r t i c l e i n f o
a b s t r a c t
Article history: Received 15 December 2008 Received in revised form 9 March 2009 Accepted 4 May 2009 Available online 12 May 2009
Cerium oxide thin films were prepared by combined electron beam evaporation and ion beam assisted deposition techniques (EBE–IBAD). Their crystallographic structures, microstructures, and optical properties were studied as a function of the substrate temperature (200 C and 500 C) and the dose of Arþ or Oþ 2 ion assistance during growth. X-ray diffraction was used to estimate the crystallographic texture, grain size, microstrain and lattice constant. Sample microstructure was studied by scanning electron microscopy. Transmission UV–vis spectroscopy was employed to obtain optical information (band gap, density, and refractive index). All films showed a cubic CeO2 structure with different preferential growth depending on the preparation conditions. The bombardment with Arþ ions during film deposition proved to be very effective for changing the film structure, hindering columnar growth and producing smaller grain sizes and higher values of microstrain and lattice constant. Films grown at 200 C and Arþ ion assistance showed the highest density, the smallest grain size (w10 nm) and a high expansion of the lattice constant (up to w1%). This expansion is related to the presence of Ce3þ at the grain boundaries. Ion assistance during the growth leads to films with higher values of refractive index and lower values of band gap. Ó 2009 Elsevier Masson SAS. All rights reserved.
Keywords: CeO2 Band gap Refractive index Grain size XRD XPS SEM
1. Introduction Cerium oxide-based compounds have attracted much attention in the context of solid oxide fuel cells that operate at intermediate temperatures (IT-SOFC) [1] because they may be incorporated as thin films in different parts of these devices. Ceria films have been proposed as the solid electrolyte in IT-SOFC due to their high ionic conductivity at relatively low temperatures [2], or as the buffer layer between the solid electrolyte and the anode, because these films may act as a diffusion barrier [3]. In general, the films involved in IT-SOFC have to be as thin as possible to minimize ohmic losses in the final devices [4], but, depending on the particular application, they have to present adequate compactness and electrical performance [1–3,5,6]. Cerium oxide films have also been proposed for optical, electro-optical, microelectronic and optoelectronic devices [7–13]. In these applications, these films are interesting due to their high
* Present address: Centre de Reserche Public Gabriel Lippmann, Rue du Brill, 41, L4422 Belvaux, Luxembourg. Tel./fax: þ352470261. E-mail address: [email protected] 1293-2558/$ – see front matter Ó 2009 Elsevier Masson SAS. All rights reserved. doi:10.1016/j.solidstatesciences.2009.05.001
refractive index, dc permittivity and transparency in the visible and near- and mid-IR. Various reports have addressed the optical properties of ceria films [7,9,14–17]; different values for the refractive index can be found in the literature, between 1.6 and 2.4 [12,18–20]. The indirect band gap varies between 2.9 and 3.3 eV, while direct band gap is found to be between 3.2 and 3.6 eV. However, higher values of the direct band gap up to 3.95 and 4.17 eV can be found for nanorods and nanoparticles, respectively [21,22]. Cerium oxide films can be prepared by several methods, such as MOCVD [23], magnetron sputtering [15,19], spray pyrolysis [24], sol-gel [25], PECVD [26], spin coating [18] and electron beam evaporation (EBE) [20,27]. The last has the advantages of minimum contamination and precise control of stoichiometry in the case of mixed oxide growth. Assistance with energetic ions is a well-known procedure to increase compactness of a growing thin film. Thus, thin film growth by EBE combined with ion beam assisted deposition (EBE–IBAD) is a promising procedure to tailor the growth and characteristics of thin films. The aim of this work is to study the influence of several experimental parameters (substrate temperature and assistance with Oþ 2 or Arþ ions) on the characteristics (grain size, morphology, density,
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
oxidation state, etc.) and optical properties of cerium oxide thin films prepared by EBE–IBAD. 2. Experimental details Cerium oxide thin films were deposited by EBE–IBAD in a high vacuum chamber (base pressure < 106 mbar) on quartz and Si polished wafers for optical and microstructural characterizations, respectively. CeO2 pellets sintered at 1500 C were used as the evaporation material. The distance between the evaporator and the samples was w50 cm to obtain a uniform thickness of the deposits, which was fixed between 300 nm and 1 mm. The experimental system allows for control of the substrate temperature, deposition rate, and dose and type of ion assistance during deposition. Tailored control of the film microstructure and optical response can be achieved varying these experimental parameters. Two substrate temperatures (Ts ¼ 200 C and Ts ¼ 500 C) and three atmospheres (vacuum, O2 and Ar) were used. The dose of Arþ or Oþ 2 ion bombardment was controlled by means of three different transport ratios (I/A), with I being bombardment current and A the deposition rate: I/A ¼ 0 (no bombardment, I ¼ 0); I/A ¼ 0.2 (moderate bombardment, one bombardment ion per five deposited atoms); and I/A ¼ 1 (high dose of bombardment, one bombardment ion per deposited atom). The energy of the ions was fixed at 600 eV, which produced a discharge current in the sample holder (I) of w2 mA/cm2 (i.e., I w 1015 ions/s/cm2). The angle of incidence of the ion beam with respect to the substrate surfaces was 45 . Two deposition rates (A) of 0.2 and 1.0 Å/s (i.e., A w 1015 and 5 1015 atoms/s/cm2) controlled by a quartz microbalance were used. Non-assisted and ion-assisted (either Arþ or Oþ 2 ) films were prepared. Non-assisted films were prepared in an O2, Ar or vacuum atmosphere. A summary of the deposition conditions of the films used in this study is displayed in Table 1. The samples are labeled in order to summarize their preparation conditions as Fn1Atn2; F indicates ‘film’, n1 indicates the substrate temperature (2 for 200 C and 5 for 500 C), At refers to the growing atmosphere (Ar for Argon, Ox for Oxygen and V for vacuum), and n2 is a second number that indicates the transport ratio (0, 0.2 or 1). The texture, mean grain size, microstrain, and mean lattice constant of the films were determined by X-ray diffraction (XRD) in
1457
a Siemens D-5000 diffractometer operating in a Bragg–Brentano configuration with Cu Ka radiation. The texture of the films was characterized by the textural parameters (Thkl) [28,29]:
Thkl ¼
1 m
Pm
0 Ihkl =Ihkl
(1)
0 i ¼ 1 Ihi ;ki ;li =Ihi ;ki ;li
where m is the number of peaks considered, h, k and l are the Miller 0 are indexes of the diffraction reflection considered, and Ihkl and Ihkl the peak intensities of the reflection of the (hkl) plane observed on the sample and a misoriented reference, respectively. A textural parameter Thkl > 1 indicates the preferential growth of the (hkl) plane parallel to the sample surface, while Thkl < 1 implies the restricted growth of the (hkl) plane. A totally misoriented sample would show Thkl ¼ 1 for each (hkl) plane. Grain sizes (D) and microstrain (3) were evaluated using the Williamson–Hall (W–H) method [30] considering linear regressions according to the following expression:
Bhkl
cos qhkl
l
¼
0:9 sin qhkl þ 43 l D
(2)
where l is the wavelength of the radiation, qhkl is the diffraction angle, and Bhkl is the full width at half maximum of the (hkl) diffraction peak, excluding the instrumental broadening. If the contribution of the microstrain is negligible, the previous expression is reduced to the well-known Scherrer equation [31]:
Dhkl ¼
0:9l Bhkl cos qhkl
(3)
This expression can be used for all the peaks of the diffractogram. Since the presence of crystallographic texture is often present in thin films, the textural parameters were used as weight factors in order to average the crystal size ðDÞ values obtained from different peaks (Dhkl):
D ¼
m 1X Tðhi ; ki ; li ÞDðhi ; ki ; li Þ m i¼1
(4)
Table 1 Schematic summary of the preparation conditions and main characteristics of the ceria thin films. Preparation conditions T ( C)
200
500
Atmosphere
I/A
Properties Label
XRD Mean grain size (nm)
Microstrain (%)
Lattice constant (Å)
Textural parameters
SEM
UV–vis
Thickness (nm)
Thickness (nm)
Refractive index
Density (g/cm3)
Band gap (eV)
T111
T220
T311
IND
DIR
Vacuum
–
F2V
29
0.03
5.405
2.87
0.08
0.05
600
560
1.82
5.1
3.25
3.50
O2
0 0.2 1
F2Ox0 F2Ox02 F2Ox1
29 28 23
0.23 0.21 0.02
5.426 5.418 5.421
2.18 2.54 0.10
0.20 0.35 0.93
0.62 0.11 1.93
1400 – 775
1152 464 700
1.87 1.85 1.93
5.3 5.2 5.5
2.93 3.18 3.13
3.22 3.47 3.36
Ar
0 0.2 1
F2Ar0 F2Ar02 F2Ar1
28 – 8
0.13 – 1.18
5.410 – 5.465
2.91 – 0.66
0.06 – 0.70
0.03 – 1.64
1131 – 226
963 664 228
1.84 1.87 2.41
5.2 5.3 7.2
3.03 3.18 3.16
3.29 3.43 3.45
Vacuum
–
F5V
67
0.03
5.409
0.17
2.41
0.42
530
673
1.99
5.8
3.25
3.45
O2
0 0.2 1
F5Ox0 F5Ox02 F5Ox1
55 73 49
0.18 0.11 0.07
5.409 5.415 5.414
1.51 1.12 0.67
0.95 1.36 1.82
0.54 0.52 0.51
506 – 510
757 458 410
1.77 1.90 2.02
4.8 5.4 5.9
2.93 3.22 3.05
3.25 3.46 3.31
Ar
0 0.2 1
F5Ar0 F5Ar02 F5Ar1
76 29 13
0.04 0.18 0.88
5.414 5.422 5.440
0.24 1.27 1.53
2.47 1.05 0.48
0.29 0.67 0.99
450 – 287
544 512 491
1.86 2.03 2.20
5.2 5.9 6.5
3.18 3.15 3.09
3.42 3.40 3.35
1458
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
The lattice constant (a) of the fluorite cubic structure of the cerium oxide was calculated from the peak position in the diffraction patterns according to the expression:
a ¼
l 2sin qhkl
pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi h2 þ k2 þ l2
(5)
In a manner similar to the calculations for the grain size, the textural parameters were used as weight factors when averaging values from different peaks in order to obtain a mean lattice constant ðaÞ. The microstructure and thickness of the films were characterized by planar-view and cross-sectional images taken by a field emission scanning electron microscope (Hitachi FESEM S-5200) working at an extraction voltage of 6.5 keV and an acceleration potential of 5 keV. The oxidation state of the cerium atoms at the surface of the films was determined by X-ray photoelectron spectroscopy (XPS) using a VG-ESCALAB 210 electron spectrometer. Unmonochromatized Mg Ka radiation was used as the excitation source. XPS measurements were performed using normal and grazing angle (60 with respect to the surface normal) detection. The refractive index (n) and optical band gap (Eg) were determined by transmission UV–vis spectroscopy on thin films deposited on fused quartz [18,32–34]. The acquired transmittance (T) spectra were simulated according to a procedure developed by Swanepoel [32], where n is obtained as the best fitting parameter to the transmittance data. Eg was evaluated by assuming that the absorption coefficient a due to interband transitions near the band gap (a approximately ln(T)) is well described by the following expression [35]:
aEf E Eg
h
(6)
where E is the photon energy (in eV) and h is a constant with the value 0.5 or 2 for direct or indirect allowed transitions, respectively. Eg was obtained using Tauc plots where (aE)1/h was plotted against E [20,36]. The density (d) of the films was estimated according to the Lorentz expression [18,33]:
n2 1 n2 1 d ¼ c2 dc 2 n þ2 nc þ 2
(7)
where nc and dc are the refractive index at a wavelength far from the absorption region and the density of the compact material (in this case, a CeO2 single crystal), respectively. The value nc ¼ 2.4 was taken as a mean value among the values reported by several authors [13,15,37–39], and dc ¼ 7.12 g/cm3 was taken from the JCPDS cards. 3. Results 3.1. Structural analysis Fig. 1 shows the diffraction patterns of the cerium oxide thin films obtained under the different experimental conditions as described in Table 1. The diffraction pattern of a CeO2 powder sintered at 1500 C was included as a reference at the top of these figures. All the diffraction patterns of the samples correspond to the fluorite cubic structure of CeO2 (JCPDS ref. 34-0394). The film texture depends on the growth conditions (see Table 1). The (200) peak could be detected only in few films, so only (111), (220) and (311) diffractions were used to do the texture calculations (i.e., m ¼ 3 in Eq. (1)). The samples prepared at 200 C show the strong preferential growth of the (111) peak (T111 > 2), except for those
prepared with a high dose of ion bombardment (I/A ¼ 1), which present a preferred (311) orientation, although less intense than the other films. The samples prepared at 500 C are less oriented than the ones prepared at 200 C, with the main peak being (111) or (220). In this case, the bombardment has different effects depending on the impinging ion; in the case of Arþ, the preferential growth evolves from (220) to (111) when increasing the transport ratio, while the opposite is seen for Oþ 2 . The (111) plane of the fcc CeO2 unit cell has the highest atom density and, hence, the lowest surface energy [40]. Therefore, in general, fcc CeO2 tends to develop a (111) orientation under normal conditions. The growth of the (220) and (311) orientations, with higher surface energies, is more pronounced under ion bombardment and a reducing atmosphere. The Arþ bombardment during film growth leads to more amorphous films, as can be deduced from the broadening of the peaks corresponding to the samples F2Ar1, F5Ar02 and F5Ar1. In order to study the broadening of the XRD peaks, a Williamson–Hall (W–H) analysis was performed. Fig. 2 displays this representation for some selected samples. A horizontal line indicates a lack of microstrain, and the y-intercept of the linear regression according to Eq. (2) is related to the mean grain size. A typical y-axis representing Bhklcos q/l has been placed at the right side in order to show in the left y-axis the grain size values. Samples prepared at 200 C (F2Ar1 and F2Ox0) show smaller crystal sizes than the corresponding samples prepared at 500 C (F5Ar1 and F5Ox0). Microstrain is clearly present in the samples grown with high doses of Arþ assistance (F2Ar1 and F5Ar1), in contrast with sample F2Ox1, which shows an almost horizontal line. A summary of the microstrain values obtained from the previous W–H analysis (cf. Fig. 2) is presented in Fig. 3a. All the films, except those grown with Arþ at I/A ¼ 1 (with 3 ¼ 0.8–1.2%), show low microstrain values, typically below w0.2%. In addition, samples grown at 200 C show slightly larger 3 than those grown at 500 C. Mean grain sizes obtained according to the Scherrer equation are displayed in Fig. 3b. For these samples, the Scherrer equation is preferred to the W–H method for the grain size estimation; the W–H method assumes the same weight for all the peaks considered, which is not adequate here due to the presence of preferred orientation. In addition, there is a large uncertainty in the determination of the ordinate at the origin, since only three peaks are detected. Nevertheless, a strict application of the W–H method for crystallites of nanometric size with strong microstrain is not straightforward [41]. To take into account a possible underestimation of the mean grain size due to the presence of microstrain, asymmetric error bars are shown for some samples. The Scherrer and W–H expressions can be considered as a lowest and highest limit for the determination of the mean grain size. Samples prepared at 500 C present larger mean grain sizes (60–80 nm) in comparison with those prepared at 200 C (25–35 nm). In both þ cases, the bombardment with Oþ 2 or Ar ions decreases D. This effect is more pronounced when Arþ ions are used, and grain sizes as small as 10–20 nm are obtained. The mean lattice constant of the ceria films obtained for the different deposition conditions is shown in Fig. 3c. The Arþ bombardment causes an increase in the lattice constant. This effect is not so evident for Oþ 2 bombardment. The largest lattice expansions are obtained when the samples are grown with a high dose of Arþ assistance, as in the case of microstrain. The values are higher for films grown at 200 C. Thus, the lattice constant of a relaxed CeO2 (a ¼ 5.410 Å) changes up to 5.440 Å for samples deposited at 500 C (increase of 0.6 %) and up to 5.465 Å (increase of 1.0%) for samples deposited at 200 C. This is in agreement with the work of Patsalas et al. [20], who observed a high increment of the lattice constant when the Arþ ion dose increased, while the effect of the temperature was minor.
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
1459
Fig. 1. Diffraction patterns of the CeOx thin films prepared at substrate temperatures of (a) 200 C and (b) 500 C under different atmospheres and ion assistance.
The variation of the lattice constant may be attributed to macrostress effects. However, several works in the literature have explained this expansion of the lattice constant by the increased influence of the Ce2O3 shell due to the reduction of grain size [42–46]. Fig. 4a shows the dependence of the lattice constant on the mean grain size of the films prepared in this work. Other results reported in the literature [42,43,47] are included for comparison. It can be observed that there is an increase in the lattice constant when the grain size decreases. This is attributed to the presence of Ce3þ and oxygen vacancies at the grain boundaries [42–48], which diminishes the average lattice constant when the mean grain size is reduced. The concentration of Ce2O3 in the films [Ce2O3] can be estimated by considering a linear variation of the lattice constant between the value of the bulk CeO2 (a0 ¼ 5.410 Å) and half of that of the C-type Ce2O3 (a1 ¼ 5.610 Å) [47]. Under these conditions:
½Ce2 O3 ¼
a a0 a1 a0
(8)
The Oþ 2 bombardment causes less change in the lattice constant than the Arþ bombardment (cf. Fig. 3c). This effect can be partially explained by the lower reduction of the crystal size caused by the Oþ 2 ions. However, the different reactivities of the impinging ions may also play an important role; the Oþ 2 ions can oxidize the Ce2O3 shell easily, and thus the increase in the lattice constant is smaller. 3.2. Chemical analysis In order to confirm the presence of the abovementioned surfacial Ce3þ species, XPS measurements were performed. Fig. 5 shows the Ce3d XPS spectra of three selected samples (F5Ox0, F2Ox1 and F2Ar1) with a progressive increase in the lattice constant and a decrease in the crystal size, with normal (left) and grazing (right) angle detection. The Ce3d peak has a complex structure due to the strength of the final state effects caused by the strong hybridation of the Ce4f and O2p states in cerium oxides. Characteristic features
The result of this calculation is shown on the right y-axis of Fig. 4a. The maximum amount of Ce3þ (about w30%) is shown for the sample grown at 200 C under intense Arþ bombardment (F2Ar1). Other samples in the literature, with even smaller grain sizes, show amounts of Ce3þ reaching up to 100% [43]. Assuming spherical particles composed of a Ce2O3 shell (Ce3þ tends to segregate to the grain boundaries [42–48]) and a CeO2 core (see inset Fig. 4b), the thickness of the Ce2O3 shell (t) can be estimated according to the expression:
t ¼ Rð1 ð1 ½Ce2 O3 ÞÞ1=3
(9)
where R is the radius of the whole particle, taken as half of the mean grain size. The results of this calculation are shown in Fig. 4b. It can be seen that the size of the Ce2O3 layer is not constant and that it grows roughly linearly while the crystal size decreases. Thus, the variation of the lattice constant is related not only to a decrease in the size of the CeO2 core but also to an increase of the Ce2O3 shell of the ceria particles [42–46,48].
Fig. 2. Williamson–Hall plot for some selected samples. The crystal sizes are displayed on the left y-axis. See text for details.
1460
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
Fig. 4. (a) Relation between the lattice constant and crystal size for our samples and some works in the literature. The dashed line is a guide for the eye. The estimated concentration of Ce2O3 is plotted on the right y-axis. (b) Thickness of the Ce2O3 shell. The dashed line is a linear fit. See text for details. Inset: the spherical model used for the calculation of the thickness of the Ce2O3 shell. Fig. 3. Parameters obtained from the XRD analysis: (a) microstrain, (b) crystal size estimated by the Scherrer equation, and (c) lattice constant. See text for details.
of Ce4þ and Ce3þ are well established [49] and can be used to identify the presence of these species. The bar diagram at the bottom of the figures indicates the energy position of the corresponding Ce3þ (circles) and Ce4þ (triangles) species. The F5Ox0 sample shows only the peaks related to the Ce4þ oxidation state at both detection angles. The F2Ox1 sample shows a similar profile when measured at the normal condition, but a small contribution of Ce3þ can be seen in the grazing measurement (the valley at w885 eV is partially filled). In the last sample (F2Ar1), the presence of Ce3þ can be detected not only in the spectrum registered at the grazing angle, but also in the normal angle measurement. These measurements address the increment of the Ce3þ species for samples showing smaller crystal sizes and larger lattice constants. In addition, grazing angle measurements further demonstrate Ce3þ major accumulation in the surface. Thus, XPS analysis confirms the tendency shown in Fig. 4. 3.3. Microstructural analysis Fig. 6 displays the planar (left) and cross-sectional (right) SEM images of the films prepared without ion assistance during growth at 200 C and 500 C. Triangular features can be seen in all surface images, in agreement with the textural parameters obtained by XRD, which showed preferential (111) growth for samples prepared at 200 C and preferential (111) and (220) growth for those prepared at 500 C. These features are more diffuse in samples
prepared at 200 C with enhanced tips of the pyramids (with three fold symmetry) in the micrographs. This is especially evident for the sample F2V. The size of these grains is small, about 10–30 nm. Below each tip, a pile up of grains with triangular shapes of increasing size (the length of the sides of the triangles was approximately 30–200 nm) can be seen. The distribution of grains is more homogeneous in the case of F2V than in the cases of F2Ar0 and F2Ox0. Cross-sectional images of these samples confirm the structure formed by piling slim triangular-shaped crystallites with different sizes. Films prepared at 500 C show clearer triangular features, although more randomly oriented. The size of the features seems more uniform in the case of F5Ar0 (the length of the sides of the triangles w100 nm) than in F5Ox0, where a distribution of sizes can be identified (30–100 nm). Cross-section micrographs show a clear columnar growth independent of the atmosphere, in opposition to the grain piling observed for the samples grown at 200 C. This observation is in agreement with the bigger crystal sizes measured by XRD. The width of these columns is about 60–80 nm. These observations are in good agreement with the growth modes presented by Movchan and Demchish [50] and later expanded by Thornton and others [51–54]. Considering the melting temperature of ceria (Tm w 2600 C), films grown at 200 C (Ts/Tm w 0.2) belong to the type I zone, formed by the piling of small grains. In contrast, samples prepared at 500 C (Ts/Tm w 0.3) belong to the type II zone, formed by columns, due to the better diffusion along the surface. Ion assistance during deposition also affects the film morphology. The SEM images of samples prepared under ion assistance (I/A ¼ 1) at 200 C and 500 C are displayed in Fig. 7a and b, respectively. In the
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
1461
Fig. 5. Ce3d XPS spectra for samples F5Ox0, F2Ox1 and F2Ar1 at normal incidence (a) and grazing incidence (b). The energy positions of Ce4þ and Ce3þ contributions are indicated by red and black vertical lines, respectively.
sample F5Ox1, some triangular features can still be identified, but with smaller randomly orientated pieces in between. No clear columnar growth can be seen in the cross-sectional image of this sample. The triangular shapes observed previously are completely lost in the case of the sample grown with Arþ bombardment (F5Ar1), which is composed of small rounded grains (20–40 nm) and
a compact structure. Regarding the samples grown at 200 C, ion assistance completely changed the grain shapes observed previously at the surface. In the case of the F2Ox1 film, the triangular features were replaced by rounded ones (100 nm) formed by long (150–200 nm) and thin (w10 nm) needles, while the cross-section seems more compact. In case of F2Ar1 film, the grains detected are
Fig. 6. SEM images in planar and cross-sectional views for films prepared without ion assistance at substrate temperatures of 200 C (left) and 500 C (right).
1462
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
Fig. 7. SEM images in planar and cross-sectional views for films prepared with high dose of ion assistance (I/A ¼ 1) at substrate temperatures of 200 C (left) and 500 C (right).
round and small, 10–20 nm, and form a smooth surface. Examining cross-sectional image confirms the changes in the surface, showing a very dense morphology and no signs of any structure (grains or columns). Thus, it is clear that the ion bombardment makes columnar growth more difficult, leading to more dense structures and smaller crystal sizes, as reported in the analysis of the XRD patterns. This is particularly clear for the sample prepared at 200 C with Arþ bombardment. The effect of bombardment is to reduce the grain size and increase the film density, due to an increase of the nucleation processes. The ion bombardment is not present explicitly in the Thornton [51,52] or the Messier models [53,54]. However, bombardment may be considered an obstacle for grain growth, comparable to pressure in the Thornton model and substrate voltage in the Messier model. 3.4. Optical characterization The optical characterization of the cerium oxide thin films was performed by UV–vis transmittance measurements. The refractive index and band gap energy were obtained by the fitting methods described in Section 2. As an example, Fig. 8 presents the transmittance spectra of the F5V (left) and F5Ar1 (right) samples. Circles represent experimental data, while solid lines represent the simulated transmission spectra. Similar fits were obtained for other samples. The spectra are characterized by the typical oscillations due to the interference effects found when transparent thin films are deposited on transparent substrates [32]. In general, the larger the amplitude of these oscillations is, the higher the refractive index of the films. On the other hand, the number of oscillations increases with the film thickness. It is worth mentioning that the thicknesses obtained according to this fitting procedure showed in general a good agreement with those obtained from the crosssectional SEM images (cf. Table 1). The insets of Fig. 8 show the corresponding absorption coefficients at the threshold of absorption used in the band gap determination. It can be seen that the direct band gap for film F5Ar1 is w3.45 eV, while for sample F5V, it is w3.35 eV. A shift towards the visible was observed for the bombarded sample (F5Ar1). Fig. 9a summarizes the effect of ion assistance and substrate temperature during growth on the refractive index of the films. As a general trend, higher substrate temperatures and greater ion assistance increase the refractive index of the films. On the other hand, bombardment with Arþ ions during the film growth is more
effective at increasing n than with Oþ 2 ions. The values of n vary between w1.75 for samples prepared without ion assistance (F5Ox0) and w2.40 for samples prepared with Arþ assistance at I/A ¼ 1 (F2Ar1), which are close to the values reported for CeO2 crystals and to the values obtained by Gorman et al. [18] (between 1.75 and 2.3), and they are in the same range as those reported by Patsalas et al. [12] (between 1.6 and 2.15) and Bueno et al. [19] (between 2.28 and 2.36). The film densities, obtained according to Eq. (7), can be checked on the right y-axis of Fig. 9a. They follow the same trends as the refractive index described above. The films with the highest refractive index and density also showed the smallest mean grain size. Ion bombardment during film growth increases the density of the films, with Arþ ions again more effective than Oþ 2 ions. This can be justified by the smaller grain size and lack of columnar growth, which allows for closer packing, in good agreement with SEM observations. In fact, the densest film (with a density similar to that of the CeO2 reference) is the sample prepared at 200 C under a high dose of Arþ bombardment. Thus, changing experimental parameters, such as substrate temperature and introducing ion beam assistance during film growth, allow tailoring of the density and refractive index of ceria films. Fig. 9b displays the band gap of the ceria films for different preparation conditions. For the direct band gap, the values are between 3.2 and 3.5 eV, in agreement with the literature [12,15,19,20,55–57]. These values are in the range of those reported by Sundaram et al. [58] (3.34–3.38 eV) and are slightly lower than those reported by Guo et al. [15] and Murali et al. [25], who observed values between 3.53 and 3.6 eV. In contrast, much higher values were observed by Zhang et al. [21] and Tsunekawa et al. [22], up to 3.95 and 4.17, respectively. These authors explain such high values by quantum confinement effects for nanorods and nanoparticles, respectively. Evaluation of the indirect band gap with the same spectra gives values systematically w0.3 eV lower than those for the direct band gap for all the films. The values of the indirect band gap are also in good agreement with the literature; Guo et al. [15] observed a variation between 2.96 and 3.3 eV, while Patsalas et al. [20] reported data between 2.9 and 3.3 eV. In addition, other authors reported data in a range near ours, between 3.0 and 3.3 eV [57,59,60]. In general, a decrease in the optical band gap can be observed in our samples from w3.5 to w3.3 eV when the bombardment dose is increased. Three samples (F2Ox0, F5Ox0 and F2Ar0) do not follow this general tendency, since they show low values of Eg (down to w2.9 eV). These samples are the ones with
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
1463
Fig. 8. Example of the fitting of UV–vis spectra by the Swanepoel method [32]: (a) sample F5V and (b) sample F5Ar1. Insets: representations used for band gap determinations. Circles and red lines indicate the experimental data and fittings, respectively.
higher thicknesses estimated by UV–vis (cf Table 1), and this is possibly the reason for such low values of the band gap [61]. This observed effect of the bombardment is the opposite of that predicted by quantum confinement, but it can be explained by grain size reduction [55]. Furthermore, it can be justified by the
higher concentration of grain boundaries [56] or the increased amount of Ce3þ [12] in the shell of the grains. However, more parameters are affected when bombardment is increased (density, strain, etc.), so claiming a unique explanation is difficult in our case. 4. Conclusions The effects of substrate temperature (200 C and 500 C) and ion assistance (no bombardment, low and high dose) with two types of þ ions (Oþ 2 and Ar ) were tested on the growth of cerium oxide thin films prepared by EBE–IBAD. In general, the cerium films show an fcc unit cell, but each film shows a particular texture. All of them show a (111) orientation when prepared without ion bombardment and in the presence of O2, but the orientation changes to the (220) and (311) textures with ion bombardment and under a reducing atmosphere. The substrate temperature affects mainly the crystallite size; in the absence of bombardment, the films grown at 500 C show well-formed columns, with grain sizes larger than 60 nm and low values of microstrain. In contrast, the films grown at 200 C are formed by piling of the grains, whose sizes are smaller than 40 nm. These effects are explained by a better diffusion along the surface, which leads to grain growth and structure relaxation. Ion assistance during growth hinders columnar growth, so the crystallographic texture observed in the non-bombarded samples decreases, and denser ceria films with smaller grain size and higher values of lattice constant are obtained. The assistance of the film growth with Arþ ions is more effective than with Oþ 2 , leading to higher values of microstrain (up to 3 w 1.0), smaller grain sizes (w10 nm) and higher expansions of the lattice constant (up to w1%). This expansion is related to the presence of Ce3þ at the grain boundaries, as demonstrated by XPS. Ion assistance during the growth enables films with high values of refractive index (n ¼ 2.2– 2.4 at wavelengths of 500 nm) compared with non-assisted films (n ¼ 1.8–2.0), due to the density improvement. On the other hand, ion assistance produces ceria films with lower band gap values. Acknowledgments
Fig. 9. Some parameters obtained from the UV–vis analysis: (a) refractive index and density, (b) direct band gap (patterned bars) and indirect band gap (solid bars).
The author would like to thank the Spanish Ministry of Science and Education for financial support (project MAT-2004-01558).
1464
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30]
M. Mogensen, N.M. Sammes, G.A. Tompsett, Solid State Ionics 129 (2000) 63. H. Inaba, H. Tagawa, Solid State Ionics 83 (1996) 1. D. Lee, J.H. Han, E.G. Kim, R.H. Song, D.R. Shin, J. Power Sources 185 (2008) 207. B.C.H. Steele, J. Mater. Sci. 36 (2001) 1053. B.C.H. Steele, Solid State Ionics 134 (2000) 3. D. Beckel, A. Bieberle-Hu¨tter, A. Harvey, A. Infortuna, U.P. Muecke, M. Prestat, J.L.M. Rupp, L.J. Gauckler, J. Power Sources 173 (2007) 325. G. Atanassov, R. Thielsch, D. Popov, Thin Solid Films 223 (1993) 288. T. Ami, M. Suzuki, Mater. Sci. Eng. B 54 (1998) 84. M.G. Krishna, A. Hartridge, A.K. Bhattacharya, Mater. Sci. Eng B 55 (1998) 14. S. Logothetidis, P. Patsalas, E.K. Evangelou, N. Konofaos, I. Tsiaoussis, N. Frangis, Mater. Sci. Eng. B 109 (2004) 69. F. Marabelli, P. Wachter, Phys. Rev. B 36 (1987) 1238. P. Patsalas, S. Logothetidis, C. Metaxa, Appl. Phys. Lett. 81 (2002) 466. T. Inoue, Y. Yamamoto, S. Koyama, S. Suzuki, Y. Ueda, Appl. Phys. Lett. 56 (1990) 1332. D.P. Norton, J.D. Budai, M.F. Chisholm, Appl. Phys. Lett. 76 (2000) 1677. S. Guo, H. Arwin, S.N. Jacobsen, K. Jarrendahl, U. Helmersson, J. Appl. Phys. 77 (1995) 5369. L. Mechin, A. Chabli, F. Bertin, M. Burdin, G. Rolland, C. Vannuffel, J.C. Villegier, J. Appl. Phys. 84 (1998) 4935. A.H. Morshed, M.E. Moussa, S.M. Bedair, R. Leonard, S.X. Liu, N. ElMasry, Appl. Phys. Lett. 70 (1997) 1647. B.P. Gorman, V. Petrovsky, H.U. Anderson, T. Petrovsky, J. Mater. Res. 19 (2004) 573. R.M. Bueno, J.M. Martinez-Duart, M. Herna´ndez-Ve´lez, L. Va´zquez, J. Mater. Sci. 32 (1997) 1861. P. Patsalas, S. Logothetidis, L. Sygellou, S. Kennou, Phys. Rev. B 68 (2003). D.E. Zhang, X.M. Ni, H.G. Zheng, X.J. Zhang, J.M. Song, Solid State Sci. 8 (2006) 1290. S. Tsunekawa, J.T. Wang, Y. Kawazoe, A. Kasuya, J. Appl. Phys. 94 (2003) 3654. K. Frohlich, J. Souc, D. Machajdik, M. Jergel, J. Snauwaert, L. Hellemans, Chem. Vap. Deposition 4 (1998) 216. J.L.M. Rupp, A. Infortuna, L.J. Gauckler, Acta Mater. 54 (2006) 1721. K.R. Murali, J. Mater. Sci.: Mater. Electron. 19 (2008) 369. D. Barreca, G. Bruno, A. Gasparotto, M. Losurdo, E. Tondello, Mater. Sci. Eng. C23 (2003) 1013. C. Mansilla, J.P. Holgado, J.P. Espino´s, A.R. Gonza´lez-Elipe, F. Yubero, Surf. Coat. Technol. 202 (2007) 1256. G.B. Harries, Philos. Mag. 43 (1952) 113. L.M. Zhang, Y.S. Gong, C.B. Wang, Q. Shen, M.X. Xia, Thin Solid Films 496 (2006) 371. G.K. Williamson, W.H. Hall, Acta Metall. 1 (1953) 22.
[31] P. Scherrer, Nachr. Ges. Wiss. Gott. 26 (1918) 98. [32] R. Swanepoel, J. Phys, E: Sci. Instrum. 16 (1983) 1214. [33] V. Petrovsky, B.P. Gorman, H.U. Anderson, T. Petrovsky, J. Appl. Phys. 90 (2001) 2517. [34] S. Tsunekawa, T. Fukuda, A. Kasuya, J. Appl. Phys. 87 (2000) 1318. [35] O. Stenzel, The Physics of Thin Film Optical Spectra: an Introduction, Springer, 2005. [36] J. Tauc, R. Grigrovici, A. Vancu, Phys. Status Solidi B 15 (1966) 627. [37] M.S. Alrobaee, M.G. Krishna, K.N. Rao, S. Mohan, J. Vac. Sci. Technol., A 9 (1991) 3048. [38] M.S. Alrobaee, K.N. Rao, S. Mohan, J. Appl. Phys. 71 (1992) 2380. [39] M.S. Alrobaee, L. Shivalingappa, K.N. Rao, S. Mohan, Thin Solid Films 221 (1992) 214. [40] Y.J. Mao, X.H. Liu, F. Zhang, C.X. Ren, S.C. Zou, Surf. Coat. Technol. 104 (1998) 78. [41] S. Brandstetter, P.M. Derlet, S. Van Petegem, H. Van Swygenhoven, Acta Mater. 56 (2008) 165. [42] F. Zhang, S.W. Chan, J.E. Spanier, E. Apak, Q. Jin, R.D. Robinson, I.P. Herman, Appl. Phys. Lett. 80 (2002) 127. [43] L. Wu, H.J. Wiesmann, A.R. Moodenbaugh, R.F. Klie, Y. Zhu, D.O. Welch, M. Suenaga, Phys. Rev. B 69 (2004) 1254151. [44] S. Tsunekawa, K. Ishikawa, Z.Q. Li, Y. Kawazoe, A. Kasuya, Phys. Rev. Lett. 85 (2000) 3440. [45] A. Kossoy, Y. Feldman, E. Wachtel, K. Gartsman, I. Lubomirsky, J. Fleig, J. Maier, Phys. Chem. Chem. Phys. 8 (2006) 1111. [46] S. Tsunekawa, R. Sahara, Y. Kawazoe, K. Ishikawa, Appl. Surf. Sci. 152 (1999) 53. [47] S. Tsunekawa, S. Ito, Y. Kawazoe, Appl. Phys. Lett. 85 (2004) 3845. [48] S. Deshpande, S. Patil, S.V. Kuchibhatla, S. Seal, Appl. Phys. Lett. 87 (2005) 133113. [49] P. Burroughs, A. Hamnett, A.F. Orchard, G. Thornton, J. Chem. Soc., Dalton Trans. (1976) 1686. [50] B.A. Movchan, A.V. Demchish, Phys. Met. Metall. 28 (1969) 83. [51] J.A. Thornton, J. Vac. Sci. Technol. 11 (1974) 666. [52] J.A. Thornton, Annu. Rev. Mater. Sci. 7 (1977) 239. [53] R. Messier, A.P. Giri, R.A. Roy, J. Vac. Sci. Technol. A2 (1984) 500. [54] R. Messier, R.C. Ross, J. Appl. Phys. 53 (1982) 6220. [55] T. Suzuki, I. Kosacki, V. Petrovsky, H.U. Anderson, J. Appl. Phys. 91 (2002) 2308. [56] R. Gallage, A. Matsuo, T. Watanabe, N. Matsushita, M. Yoshimura, J Electroceram. (2008) 1. [57] N. Ozer, Sol. Energy Mater. Sol. Cells 68 (2001) 391. [58] K.B. Sundaram, P. Wahind, Phys. Status Solidi B 161 (1990) K63. [59] S.Y. Zheng, A.M. Andersson-Faldt, B. Stjerna, C.G. Granqvist, Appl. Opt. 32 (1993) 6303. [60] C.A. Hogarth, Z.T. Al-Dhhan, Phys. Status Solidi B 137 (1986). [61] S. Debnath, M.R. Islam, M.S.R. Khan, Bull. Mater. Sci. 30 (2007) 315.
Contents lists available at ScienceDirect
Solid State Sciences journal homepage: www.elsevier.com/locate/ssscie
Structure, microstructure and optical properties of cerium oxide thin films prepared by electron beam evaporation assisted with ion beams Catalina Mansilla* Instituto de Ciencia de Materiales de Sevilla, CSIC-US, Avda. Ame´rico Vespucio 49, 41092-Sevilla, Spain
a r t i c l e i n f o
a b s t r a c t
Article history: Received 15 December 2008 Received in revised form 9 March 2009 Accepted 4 May 2009 Available online 12 May 2009
Cerium oxide thin films were prepared by combined electron beam evaporation and ion beam assisted deposition techniques (EBE–IBAD). Their crystallographic structures, microstructures, and optical properties were studied as a function of the substrate temperature (200 C and 500 C) and the dose of Arþ or Oþ 2 ion assistance during growth. X-ray diffraction was used to estimate the crystallographic texture, grain size, microstrain and lattice constant. Sample microstructure was studied by scanning electron microscopy. Transmission UV–vis spectroscopy was employed to obtain optical information (band gap, density, and refractive index). All films showed a cubic CeO2 structure with different preferential growth depending on the preparation conditions. The bombardment with Arþ ions during film deposition proved to be very effective for changing the film structure, hindering columnar growth and producing smaller grain sizes and higher values of microstrain and lattice constant. Films grown at 200 C and Arþ ion assistance showed the highest density, the smallest grain size (w10 nm) and a high expansion of the lattice constant (up to w1%). This expansion is related to the presence of Ce3þ at the grain boundaries. Ion assistance during the growth leads to films with higher values of refractive index and lower values of band gap. Ó 2009 Elsevier Masson SAS. All rights reserved.
Keywords: CeO2 Band gap Refractive index Grain size XRD XPS SEM
1. Introduction Cerium oxide-based compounds have attracted much attention in the context of solid oxide fuel cells that operate at intermediate temperatures (IT-SOFC) [1] because they may be incorporated as thin films in different parts of these devices. Ceria films have been proposed as the solid electrolyte in IT-SOFC due to their high ionic conductivity at relatively low temperatures [2], or as the buffer layer between the solid electrolyte and the anode, because these films may act as a diffusion barrier [3]. In general, the films involved in IT-SOFC have to be as thin as possible to minimize ohmic losses in the final devices [4], but, depending on the particular application, they have to present adequate compactness and electrical performance [1–3,5,6]. Cerium oxide films have also been proposed for optical, electro-optical, microelectronic and optoelectronic devices [7–13]. In these applications, these films are interesting due to their high
* Present address: Centre de Reserche Public Gabriel Lippmann, Rue du Brill, 41, L4422 Belvaux, Luxembourg. Tel./fax: þ352470261. E-mail address: [email protected] 1293-2558/$ – see front matter Ó 2009 Elsevier Masson SAS. All rights reserved. doi:10.1016/j.solidstatesciences.2009.05.001
refractive index, dc permittivity and transparency in the visible and near- and mid-IR. Various reports have addressed the optical properties of ceria films [7,9,14–17]; different values for the refractive index can be found in the literature, between 1.6 and 2.4 [12,18–20]. The indirect band gap varies between 2.9 and 3.3 eV, while direct band gap is found to be between 3.2 and 3.6 eV. However, higher values of the direct band gap up to 3.95 and 4.17 eV can be found for nanorods and nanoparticles, respectively [21,22]. Cerium oxide films can be prepared by several methods, such as MOCVD [23], magnetron sputtering [15,19], spray pyrolysis [24], sol-gel [25], PECVD [26], spin coating [18] and electron beam evaporation (EBE) [20,27]. The last has the advantages of minimum contamination and precise control of stoichiometry in the case of mixed oxide growth. Assistance with energetic ions is a well-known procedure to increase compactness of a growing thin film. Thus, thin film growth by EBE combined with ion beam assisted deposition (EBE–IBAD) is a promising procedure to tailor the growth and characteristics of thin films. The aim of this work is to study the influence of several experimental parameters (substrate temperature and assistance with Oþ 2 or Arþ ions) on the characteristics (grain size, morphology, density,
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
oxidation state, etc.) and optical properties of cerium oxide thin films prepared by EBE–IBAD. 2. Experimental details Cerium oxide thin films were deposited by EBE–IBAD in a high vacuum chamber (base pressure < 106 mbar) on quartz and Si polished wafers for optical and microstructural characterizations, respectively. CeO2 pellets sintered at 1500 C were used as the evaporation material. The distance between the evaporator and the samples was w50 cm to obtain a uniform thickness of the deposits, which was fixed between 300 nm and 1 mm. The experimental system allows for control of the substrate temperature, deposition rate, and dose and type of ion assistance during deposition. Tailored control of the film microstructure and optical response can be achieved varying these experimental parameters. Two substrate temperatures (Ts ¼ 200 C and Ts ¼ 500 C) and three atmospheres (vacuum, O2 and Ar) were used. The dose of Arþ or Oþ 2 ion bombardment was controlled by means of three different transport ratios (I/A), with I being bombardment current and A the deposition rate: I/A ¼ 0 (no bombardment, I ¼ 0); I/A ¼ 0.2 (moderate bombardment, one bombardment ion per five deposited atoms); and I/A ¼ 1 (high dose of bombardment, one bombardment ion per deposited atom). The energy of the ions was fixed at 600 eV, which produced a discharge current in the sample holder (I) of w2 mA/cm2 (i.e., I w 1015 ions/s/cm2). The angle of incidence of the ion beam with respect to the substrate surfaces was 45 . Two deposition rates (A) of 0.2 and 1.0 Å/s (i.e., A w 1015 and 5 1015 atoms/s/cm2) controlled by a quartz microbalance were used. Non-assisted and ion-assisted (either Arþ or Oþ 2 ) films were prepared. Non-assisted films were prepared in an O2, Ar or vacuum atmosphere. A summary of the deposition conditions of the films used in this study is displayed in Table 1. The samples are labeled in order to summarize their preparation conditions as Fn1Atn2; F indicates ‘film’, n1 indicates the substrate temperature (2 for 200 C and 5 for 500 C), At refers to the growing atmosphere (Ar for Argon, Ox for Oxygen and V for vacuum), and n2 is a second number that indicates the transport ratio (0, 0.2 or 1). The texture, mean grain size, microstrain, and mean lattice constant of the films were determined by X-ray diffraction (XRD) in
1457
a Siemens D-5000 diffractometer operating in a Bragg–Brentano configuration with Cu Ka radiation. The texture of the films was characterized by the textural parameters (Thkl) [28,29]:
Thkl ¼
1 m
Pm
0 Ihkl =Ihkl
(1)
0 i ¼ 1 Ihi ;ki ;li =Ihi ;ki ;li
where m is the number of peaks considered, h, k and l are the Miller 0 are indexes of the diffraction reflection considered, and Ihkl and Ihkl the peak intensities of the reflection of the (hkl) plane observed on the sample and a misoriented reference, respectively. A textural parameter Thkl > 1 indicates the preferential growth of the (hkl) plane parallel to the sample surface, while Thkl < 1 implies the restricted growth of the (hkl) plane. A totally misoriented sample would show Thkl ¼ 1 for each (hkl) plane. Grain sizes (D) and microstrain (3) were evaluated using the Williamson–Hall (W–H) method [30] considering linear regressions according to the following expression:
Bhkl
cos qhkl
l
¼
0:9 sin qhkl þ 43 l D
(2)
where l is the wavelength of the radiation, qhkl is the diffraction angle, and Bhkl is the full width at half maximum of the (hkl) diffraction peak, excluding the instrumental broadening. If the contribution of the microstrain is negligible, the previous expression is reduced to the well-known Scherrer equation [31]:
Dhkl ¼
0:9l Bhkl cos qhkl
(3)
This expression can be used for all the peaks of the diffractogram. Since the presence of crystallographic texture is often present in thin films, the textural parameters were used as weight factors in order to average the crystal size ðDÞ values obtained from different peaks (Dhkl):
D ¼
m 1X Tðhi ; ki ; li ÞDðhi ; ki ; li Þ m i¼1
(4)
Table 1 Schematic summary of the preparation conditions and main characteristics of the ceria thin films. Preparation conditions T ( C)
200
500
Atmosphere
I/A
Properties Label
XRD Mean grain size (nm)
Microstrain (%)
Lattice constant (Å)
Textural parameters
SEM
UV–vis
Thickness (nm)
Thickness (nm)
Refractive index
Density (g/cm3)
Band gap (eV)
T111
T220
T311
IND
DIR
Vacuum
–
F2V
29
0.03
5.405
2.87
0.08
0.05
600
560
1.82
5.1
3.25
3.50
O2
0 0.2 1
F2Ox0 F2Ox02 F2Ox1
29 28 23
0.23 0.21 0.02
5.426 5.418 5.421
2.18 2.54 0.10
0.20 0.35 0.93
0.62 0.11 1.93
1400 – 775
1152 464 700
1.87 1.85 1.93
5.3 5.2 5.5
2.93 3.18 3.13
3.22 3.47 3.36
Ar
0 0.2 1
F2Ar0 F2Ar02 F2Ar1
28 – 8
0.13 – 1.18
5.410 – 5.465
2.91 – 0.66
0.06 – 0.70
0.03 – 1.64
1131 – 226
963 664 228
1.84 1.87 2.41
5.2 5.3 7.2
3.03 3.18 3.16
3.29 3.43 3.45
Vacuum
–
F5V
67
0.03
5.409
0.17
2.41
0.42
530
673
1.99
5.8
3.25
3.45
O2
0 0.2 1
F5Ox0 F5Ox02 F5Ox1
55 73 49
0.18 0.11 0.07
5.409 5.415 5.414
1.51 1.12 0.67
0.95 1.36 1.82
0.54 0.52 0.51
506 – 510
757 458 410
1.77 1.90 2.02
4.8 5.4 5.9
2.93 3.22 3.05
3.25 3.46 3.31
Ar
0 0.2 1
F5Ar0 F5Ar02 F5Ar1
76 29 13
0.04 0.18 0.88
5.414 5.422 5.440
0.24 1.27 1.53
2.47 1.05 0.48
0.29 0.67 0.99
450 – 287
544 512 491
1.86 2.03 2.20
5.2 5.9 6.5
3.18 3.15 3.09
3.42 3.40 3.35
1458
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
The lattice constant (a) of the fluorite cubic structure of the cerium oxide was calculated from the peak position in the diffraction patterns according to the expression:
a ¼
l 2sin qhkl
pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi h2 þ k2 þ l2
(5)
In a manner similar to the calculations for the grain size, the textural parameters were used as weight factors when averaging values from different peaks in order to obtain a mean lattice constant ðaÞ. The microstructure and thickness of the films were characterized by planar-view and cross-sectional images taken by a field emission scanning electron microscope (Hitachi FESEM S-5200) working at an extraction voltage of 6.5 keV and an acceleration potential of 5 keV. The oxidation state of the cerium atoms at the surface of the films was determined by X-ray photoelectron spectroscopy (XPS) using a VG-ESCALAB 210 electron spectrometer. Unmonochromatized Mg Ka radiation was used as the excitation source. XPS measurements were performed using normal and grazing angle (60 with respect to the surface normal) detection. The refractive index (n) and optical band gap (Eg) were determined by transmission UV–vis spectroscopy on thin films deposited on fused quartz [18,32–34]. The acquired transmittance (T) spectra were simulated according to a procedure developed by Swanepoel [32], where n is obtained as the best fitting parameter to the transmittance data. Eg was evaluated by assuming that the absorption coefficient a due to interband transitions near the band gap (a approximately ln(T)) is well described by the following expression [35]:
aEf E Eg
h
(6)
where E is the photon energy (in eV) and h is a constant with the value 0.5 or 2 for direct or indirect allowed transitions, respectively. Eg was obtained using Tauc plots where (aE)1/h was plotted against E [20,36]. The density (d) of the films was estimated according to the Lorentz expression [18,33]:
n2 1 n2 1 d ¼ c2 dc 2 n þ2 nc þ 2
(7)
where nc and dc are the refractive index at a wavelength far from the absorption region and the density of the compact material (in this case, a CeO2 single crystal), respectively. The value nc ¼ 2.4 was taken as a mean value among the values reported by several authors [13,15,37–39], and dc ¼ 7.12 g/cm3 was taken from the JCPDS cards. 3. Results 3.1. Structural analysis Fig. 1 shows the diffraction patterns of the cerium oxide thin films obtained under the different experimental conditions as described in Table 1. The diffraction pattern of a CeO2 powder sintered at 1500 C was included as a reference at the top of these figures. All the diffraction patterns of the samples correspond to the fluorite cubic structure of CeO2 (JCPDS ref. 34-0394). The film texture depends on the growth conditions (see Table 1). The (200) peak could be detected only in few films, so only (111), (220) and (311) diffractions were used to do the texture calculations (i.e., m ¼ 3 in Eq. (1)). The samples prepared at 200 C show the strong preferential growth of the (111) peak (T111 > 2), except for those
prepared with a high dose of ion bombardment (I/A ¼ 1), which present a preferred (311) orientation, although less intense than the other films. The samples prepared at 500 C are less oriented than the ones prepared at 200 C, with the main peak being (111) or (220). In this case, the bombardment has different effects depending on the impinging ion; in the case of Arþ, the preferential growth evolves from (220) to (111) when increasing the transport ratio, while the opposite is seen for Oþ 2 . The (111) plane of the fcc CeO2 unit cell has the highest atom density and, hence, the lowest surface energy [40]. Therefore, in general, fcc CeO2 tends to develop a (111) orientation under normal conditions. The growth of the (220) and (311) orientations, with higher surface energies, is more pronounced under ion bombardment and a reducing atmosphere. The Arþ bombardment during film growth leads to more amorphous films, as can be deduced from the broadening of the peaks corresponding to the samples F2Ar1, F5Ar02 and F5Ar1. In order to study the broadening of the XRD peaks, a Williamson–Hall (W–H) analysis was performed. Fig. 2 displays this representation for some selected samples. A horizontal line indicates a lack of microstrain, and the y-intercept of the linear regression according to Eq. (2) is related to the mean grain size. A typical y-axis representing Bhklcos q/l has been placed at the right side in order to show in the left y-axis the grain size values. Samples prepared at 200 C (F2Ar1 and F2Ox0) show smaller crystal sizes than the corresponding samples prepared at 500 C (F5Ar1 and F5Ox0). Microstrain is clearly present in the samples grown with high doses of Arþ assistance (F2Ar1 and F5Ar1), in contrast with sample F2Ox1, which shows an almost horizontal line. A summary of the microstrain values obtained from the previous W–H analysis (cf. Fig. 2) is presented in Fig. 3a. All the films, except those grown with Arþ at I/A ¼ 1 (with 3 ¼ 0.8–1.2%), show low microstrain values, typically below w0.2%. In addition, samples grown at 200 C show slightly larger 3 than those grown at 500 C. Mean grain sizes obtained according to the Scherrer equation are displayed in Fig. 3b. For these samples, the Scherrer equation is preferred to the W–H method for the grain size estimation; the W–H method assumes the same weight for all the peaks considered, which is not adequate here due to the presence of preferred orientation. In addition, there is a large uncertainty in the determination of the ordinate at the origin, since only three peaks are detected. Nevertheless, a strict application of the W–H method for crystallites of nanometric size with strong microstrain is not straightforward [41]. To take into account a possible underestimation of the mean grain size due to the presence of microstrain, asymmetric error bars are shown for some samples. The Scherrer and W–H expressions can be considered as a lowest and highest limit for the determination of the mean grain size. Samples prepared at 500 C present larger mean grain sizes (60–80 nm) in comparison with those prepared at 200 C (25–35 nm). In both þ cases, the bombardment with Oþ 2 or Ar ions decreases D. This effect is more pronounced when Arþ ions are used, and grain sizes as small as 10–20 nm are obtained. The mean lattice constant of the ceria films obtained for the different deposition conditions is shown in Fig. 3c. The Arþ bombardment causes an increase in the lattice constant. This effect is not so evident for Oþ 2 bombardment. The largest lattice expansions are obtained when the samples are grown with a high dose of Arþ assistance, as in the case of microstrain. The values are higher for films grown at 200 C. Thus, the lattice constant of a relaxed CeO2 (a ¼ 5.410 Å) changes up to 5.440 Å for samples deposited at 500 C (increase of 0.6 %) and up to 5.465 Å (increase of 1.0%) for samples deposited at 200 C. This is in agreement with the work of Patsalas et al. [20], who observed a high increment of the lattice constant when the Arþ ion dose increased, while the effect of the temperature was minor.
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
1459
Fig. 1. Diffraction patterns of the CeOx thin films prepared at substrate temperatures of (a) 200 C and (b) 500 C under different atmospheres and ion assistance.
The variation of the lattice constant may be attributed to macrostress effects. However, several works in the literature have explained this expansion of the lattice constant by the increased influence of the Ce2O3 shell due to the reduction of grain size [42–46]. Fig. 4a shows the dependence of the lattice constant on the mean grain size of the films prepared in this work. Other results reported in the literature [42,43,47] are included for comparison. It can be observed that there is an increase in the lattice constant when the grain size decreases. This is attributed to the presence of Ce3þ and oxygen vacancies at the grain boundaries [42–48], which diminishes the average lattice constant when the mean grain size is reduced. The concentration of Ce2O3 in the films [Ce2O3] can be estimated by considering a linear variation of the lattice constant between the value of the bulk CeO2 (a0 ¼ 5.410 Å) and half of that of the C-type Ce2O3 (a1 ¼ 5.610 Å) [47]. Under these conditions:
½Ce2 O3 ¼
a a0 a1 a0
(8)
The Oþ 2 bombardment causes less change in the lattice constant than the Arþ bombardment (cf. Fig. 3c). This effect can be partially explained by the lower reduction of the crystal size caused by the Oþ 2 ions. However, the different reactivities of the impinging ions may also play an important role; the Oþ 2 ions can oxidize the Ce2O3 shell easily, and thus the increase in the lattice constant is smaller. 3.2. Chemical analysis In order to confirm the presence of the abovementioned surfacial Ce3þ species, XPS measurements were performed. Fig. 5 shows the Ce3d XPS spectra of three selected samples (F5Ox0, F2Ox1 and F2Ar1) with a progressive increase in the lattice constant and a decrease in the crystal size, with normal (left) and grazing (right) angle detection. The Ce3d peak has a complex structure due to the strength of the final state effects caused by the strong hybridation of the Ce4f and O2p states in cerium oxides. Characteristic features
The result of this calculation is shown on the right y-axis of Fig. 4a. The maximum amount of Ce3þ (about w30%) is shown for the sample grown at 200 C under intense Arþ bombardment (F2Ar1). Other samples in the literature, with even smaller grain sizes, show amounts of Ce3þ reaching up to 100% [43]. Assuming spherical particles composed of a Ce2O3 shell (Ce3þ tends to segregate to the grain boundaries [42–48]) and a CeO2 core (see inset Fig. 4b), the thickness of the Ce2O3 shell (t) can be estimated according to the expression:
t ¼ Rð1 ð1 ½Ce2 O3 ÞÞ1=3
(9)
where R is the radius of the whole particle, taken as half of the mean grain size. The results of this calculation are shown in Fig. 4b. It can be seen that the size of the Ce2O3 layer is not constant and that it grows roughly linearly while the crystal size decreases. Thus, the variation of the lattice constant is related not only to a decrease in the size of the CeO2 core but also to an increase of the Ce2O3 shell of the ceria particles [42–46,48].
Fig. 2. Williamson–Hall plot for some selected samples. The crystal sizes are displayed on the left y-axis. See text for details.
1460
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
Fig. 4. (a) Relation between the lattice constant and crystal size for our samples and some works in the literature. The dashed line is a guide for the eye. The estimated concentration of Ce2O3 is plotted on the right y-axis. (b) Thickness of the Ce2O3 shell. The dashed line is a linear fit. See text for details. Inset: the spherical model used for the calculation of the thickness of the Ce2O3 shell. Fig. 3. Parameters obtained from the XRD analysis: (a) microstrain, (b) crystal size estimated by the Scherrer equation, and (c) lattice constant. See text for details.
of Ce4þ and Ce3þ are well established [49] and can be used to identify the presence of these species. The bar diagram at the bottom of the figures indicates the energy position of the corresponding Ce3þ (circles) and Ce4þ (triangles) species. The F5Ox0 sample shows only the peaks related to the Ce4þ oxidation state at both detection angles. The F2Ox1 sample shows a similar profile when measured at the normal condition, but a small contribution of Ce3þ can be seen in the grazing measurement (the valley at w885 eV is partially filled). In the last sample (F2Ar1), the presence of Ce3þ can be detected not only in the spectrum registered at the grazing angle, but also in the normal angle measurement. These measurements address the increment of the Ce3þ species for samples showing smaller crystal sizes and larger lattice constants. In addition, grazing angle measurements further demonstrate Ce3þ major accumulation in the surface. Thus, XPS analysis confirms the tendency shown in Fig. 4. 3.3. Microstructural analysis Fig. 6 displays the planar (left) and cross-sectional (right) SEM images of the films prepared without ion assistance during growth at 200 C and 500 C. Triangular features can be seen in all surface images, in agreement with the textural parameters obtained by XRD, which showed preferential (111) growth for samples prepared at 200 C and preferential (111) and (220) growth for those prepared at 500 C. These features are more diffuse in samples
prepared at 200 C with enhanced tips of the pyramids (with three fold symmetry) in the micrographs. This is especially evident for the sample F2V. The size of these grains is small, about 10–30 nm. Below each tip, a pile up of grains with triangular shapes of increasing size (the length of the sides of the triangles was approximately 30–200 nm) can be seen. The distribution of grains is more homogeneous in the case of F2V than in the cases of F2Ar0 and F2Ox0. Cross-sectional images of these samples confirm the structure formed by piling slim triangular-shaped crystallites with different sizes. Films prepared at 500 C show clearer triangular features, although more randomly oriented. The size of the features seems more uniform in the case of F5Ar0 (the length of the sides of the triangles w100 nm) than in F5Ox0, where a distribution of sizes can be identified (30–100 nm). Cross-section micrographs show a clear columnar growth independent of the atmosphere, in opposition to the grain piling observed for the samples grown at 200 C. This observation is in agreement with the bigger crystal sizes measured by XRD. The width of these columns is about 60–80 nm. These observations are in good agreement with the growth modes presented by Movchan and Demchish [50] and later expanded by Thornton and others [51–54]. Considering the melting temperature of ceria (Tm w 2600 C), films grown at 200 C (Ts/Tm w 0.2) belong to the type I zone, formed by the piling of small grains. In contrast, samples prepared at 500 C (Ts/Tm w 0.3) belong to the type II zone, formed by columns, due to the better diffusion along the surface. Ion assistance during deposition also affects the film morphology. The SEM images of samples prepared under ion assistance (I/A ¼ 1) at 200 C and 500 C are displayed in Fig. 7a and b, respectively. In the
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
1461
Fig. 5. Ce3d XPS spectra for samples F5Ox0, F2Ox1 and F2Ar1 at normal incidence (a) and grazing incidence (b). The energy positions of Ce4þ and Ce3þ contributions are indicated by red and black vertical lines, respectively.
sample F5Ox1, some triangular features can still be identified, but with smaller randomly orientated pieces in between. No clear columnar growth can be seen in the cross-sectional image of this sample. The triangular shapes observed previously are completely lost in the case of the sample grown with Arþ bombardment (F5Ar1), which is composed of small rounded grains (20–40 nm) and
a compact structure. Regarding the samples grown at 200 C, ion assistance completely changed the grain shapes observed previously at the surface. In the case of the F2Ox1 film, the triangular features were replaced by rounded ones (100 nm) formed by long (150–200 nm) and thin (w10 nm) needles, while the cross-section seems more compact. In case of F2Ar1 film, the grains detected are
Fig. 6. SEM images in planar and cross-sectional views for films prepared without ion assistance at substrate temperatures of 200 C (left) and 500 C (right).
1462
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
Fig. 7. SEM images in planar and cross-sectional views for films prepared with high dose of ion assistance (I/A ¼ 1) at substrate temperatures of 200 C (left) and 500 C (right).
round and small, 10–20 nm, and form a smooth surface. Examining cross-sectional image confirms the changes in the surface, showing a very dense morphology and no signs of any structure (grains or columns). Thus, it is clear that the ion bombardment makes columnar growth more difficult, leading to more dense structures and smaller crystal sizes, as reported in the analysis of the XRD patterns. This is particularly clear for the sample prepared at 200 C with Arþ bombardment. The effect of bombardment is to reduce the grain size and increase the film density, due to an increase of the nucleation processes. The ion bombardment is not present explicitly in the Thornton [51,52] or the Messier models [53,54]. However, bombardment may be considered an obstacle for grain growth, comparable to pressure in the Thornton model and substrate voltage in the Messier model. 3.4. Optical characterization The optical characterization of the cerium oxide thin films was performed by UV–vis transmittance measurements. The refractive index and band gap energy were obtained by the fitting methods described in Section 2. As an example, Fig. 8 presents the transmittance spectra of the F5V (left) and F5Ar1 (right) samples. Circles represent experimental data, while solid lines represent the simulated transmission spectra. Similar fits were obtained for other samples. The spectra are characterized by the typical oscillations due to the interference effects found when transparent thin films are deposited on transparent substrates [32]. In general, the larger the amplitude of these oscillations is, the higher the refractive index of the films. On the other hand, the number of oscillations increases with the film thickness. It is worth mentioning that the thicknesses obtained according to this fitting procedure showed in general a good agreement with those obtained from the crosssectional SEM images (cf. Table 1). The insets of Fig. 8 show the corresponding absorption coefficients at the threshold of absorption used in the band gap determination. It can be seen that the direct band gap for film F5Ar1 is w3.45 eV, while for sample F5V, it is w3.35 eV. A shift towards the visible was observed for the bombarded sample (F5Ar1). Fig. 9a summarizes the effect of ion assistance and substrate temperature during growth on the refractive index of the films. As a general trend, higher substrate temperatures and greater ion assistance increase the refractive index of the films. On the other hand, bombardment with Arþ ions during the film growth is more
effective at increasing n than with Oþ 2 ions. The values of n vary between w1.75 for samples prepared without ion assistance (F5Ox0) and w2.40 for samples prepared with Arþ assistance at I/A ¼ 1 (F2Ar1), which are close to the values reported for CeO2 crystals and to the values obtained by Gorman et al. [18] (between 1.75 and 2.3), and they are in the same range as those reported by Patsalas et al. [12] (between 1.6 and 2.15) and Bueno et al. [19] (between 2.28 and 2.36). The film densities, obtained according to Eq. (7), can be checked on the right y-axis of Fig. 9a. They follow the same trends as the refractive index described above. The films with the highest refractive index and density also showed the smallest mean grain size. Ion bombardment during film growth increases the density of the films, with Arþ ions again more effective than Oþ 2 ions. This can be justified by the smaller grain size and lack of columnar growth, which allows for closer packing, in good agreement with SEM observations. In fact, the densest film (with a density similar to that of the CeO2 reference) is the sample prepared at 200 C under a high dose of Arþ bombardment. Thus, changing experimental parameters, such as substrate temperature and introducing ion beam assistance during film growth, allow tailoring of the density and refractive index of ceria films. Fig. 9b displays the band gap of the ceria films for different preparation conditions. For the direct band gap, the values are between 3.2 and 3.5 eV, in agreement with the literature [12,15,19,20,55–57]. These values are in the range of those reported by Sundaram et al. [58] (3.34–3.38 eV) and are slightly lower than those reported by Guo et al. [15] and Murali et al. [25], who observed values between 3.53 and 3.6 eV. In contrast, much higher values were observed by Zhang et al. [21] and Tsunekawa et al. [22], up to 3.95 and 4.17, respectively. These authors explain such high values by quantum confinement effects for nanorods and nanoparticles, respectively. Evaluation of the indirect band gap with the same spectra gives values systematically w0.3 eV lower than those for the direct band gap for all the films. The values of the indirect band gap are also in good agreement with the literature; Guo et al. [15] observed a variation between 2.96 and 3.3 eV, while Patsalas et al. [20] reported data between 2.9 and 3.3 eV. In addition, other authors reported data in a range near ours, between 3.0 and 3.3 eV [57,59,60]. In general, a decrease in the optical band gap can be observed in our samples from w3.5 to w3.3 eV when the bombardment dose is increased. Three samples (F2Ox0, F5Ox0 and F2Ar0) do not follow this general tendency, since they show low values of Eg (down to w2.9 eV). These samples are the ones with
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
1463
Fig. 8. Example of the fitting of UV–vis spectra by the Swanepoel method [32]: (a) sample F5V and (b) sample F5Ar1. Insets: representations used for band gap determinations. Circles and red lines indicate the experimental data and fittings, respectively.
higher thicknesses estimated by UV–vis (cf Table 1), and this is possibly the reason for such low values of the band gap [61]. This observed effect of the bombardment is the opposite of that predicted by quantum confinement, but it can be explained by grain size reduction [55]. Furthermore, it can be justified by the
higher concentration of grain boundaries [56] or the increased amount of Ce3þ [12] in the shell of the grains. However, more parameters are affected when bombardment is increased (density, strain, etc.), so claiming a unique explanation is difficult in our case. 4. Conclusions The effects of substrate temperature (200 C and 500 C) and ion assistance (no bombardment, low and high dose) with two types of þ ions (Oþ 2 and Ar ) were tested on the growth of cerium oxide thin films prepared by EBE–IBAD. In general, the cerium films show an fcc unit cell, but each film shows a particular texture. All of them show a (111) orientation when prepared without ion bombardment and in the presence of O2, but the orientation changes to the (220) and (311) textures with ion bombardment and under a reducing atmosphere. The substrate temperature affects mainly the crystallite size; in the absence of bombardment, the films grown at 500 C show well-formed columns, with grain sizes larger than 60 nm and low values of microstrain. In contrast, the films grown at 200 C are formed by piling of the grains, whose sizes are smaller than 40 nm. These effects are explained by a better diffusion along the surface, which leads to grain growth and structure relaxation. Ion assistance during growth hinders columnar growth, so the crystallographic texture observed in the non-bombarded samples decreases, and denser ceria films with smaller grain size and higher values of lattice constant are obtained. The assistance of the film growth with Arþ ions is more effective than with Oþ 2 , leading to higher values of microstrain (up to 3 w 1.0), smaller grain sizes (w10 nm) and higher expansions of the lattice constant (up to w1%). This expansion is related to the presence of Ce3þ at the grain boundaries, as demonstrated by XPS. Ion assistance during the growth enables films with high values of refractive index (n ¼ 2.2– 2.4 at wavelengths of 500 nm) compared with non-assisted films (n ¼ 1.8–2.0), due to the density improvement. On the other hand, ion assistance produces ceria films with lower band gap values. Acknowledgments
Fig. 9. Some parameters obtained from the UV–vis analysis: (a) refractive index and density, (b) direct band gap (patterned bars) and indirect band gap (solid bars).
The author would like to thank the Spanish Ministry of Science and Education for financial support (project MAT-2004-01558).
1464
C. Mansilla / Solid State Sciences 11 (2009) 1456–1464
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30]
M. Mogensen, N.M. Sammes, G.A. Tompsett, Solid State Ionics 129 (2000) 63. H. Inaba, H. Tagawa, Solid State Ionics 83 (1996) 1. D. Lee, J.H. Han, E.G. Kim, R.H. Song, D.R. Shin, J. Power Sources 185 (2008) 207. B.C.H. Steele, J. Mater. Sci. 36 (2001) 1053. B.C.H. Steele, Solid State Ionics 134 (2000) 3. D. Beckel, A. Bieberle-Hu¨tter, A. Harvey, A. Infortuna, U.P. Muecke, M. Prestat, J.L.M. Rupp, L.J. Gauckler, J. Power Sources 173 (2007) 325. G. Atanassov, R. Thielsch, D. Popov, Thin Solid Films 223 (1993) 288. T. Ami, M. Suzuki, Mater. Sci. Eng. B 54 (1998) 84. M.G. Krishna, A. Hartridge, A.K. Bhattacharya, Mater. Sci. Eng B 55 (1998) 14. S. Logothetidis, P. Patsalas, E.K. Evangelou, N. Konofaos, I. Tsiaoussis, N. Frangis, Mater. Sci. Eng. B 109 (2004) 69. F. Marabelli, P. Wachter, Phys. Rev. B 36 (1987) 1238. P. Patsalas, S. Logothetidis, C. Metaxa, Appl. Phys. Lett. 81 (2002) 466. T. Inoue, Y. Yamamoto, S. Koyama, S. Suzuki, Y. Ueda, Appl. Phys. Lett. 56 (1990) 1332. D.P. Norton, J.D. Budai, M.F. Chisholm, Appl. Phys. Lett. 76 (2000) 1677. S. Guo, H. Arwin, S.N. Jacobsen, K. Jarrendahl, U. Helmersson, J. Appl. Phys. 77 (1995) 5369. L. Mechin, A. Chabli, F. Bertin, M. Burdin, G. Rolland, C. Vannuffel, J.C. Villegier, J. Appl. Phys. 84 (1998) 4935. A.H. Morshed, M.E. Moussa, S.M. Bedair, R. Leonard, S.X. Liu, N. ElMasry, Appl. Phys. Lett. 70 (1997) 1647. B.P. Gorman, V. Petrovsky, H.U. Anderson, T. Petrovsky, J. Mater. Res. 19 (2004) 573. R.M. Bueno, J.M. Martinez-Duart, M. Herna´ndez-Ve´lez, L. Va´zquez, J. Mater. Sci. 32 (1997) 1861. P. Patsalas, S. Logothetidis, L. Sygellou, S. Kennou, Phys. Rev. B 68 (2003). D.E. Zhang, X.M. Ni, H.G. Zheng, X.J. Zhang, J.M. Song, Solid State Sci. 8 (2006) 1290. S. Tsunekawa, J.T. Wang, Y. Kawazoe, A. Kasuya, J. Appl. Phys. 94 (2003) 3654. K. Frohlich, J. Souc, D. Machajdik, M. Jergel, J. Snauwaert, L. Hellemans, Chem. Vap. Deposition 4 (1998) 216. J.L.M. Rupp, A. Infortuna, L.J. Gauckler, Acta Mater. 54 (2006) 1721. K.R. Murali, J. Mater. Sci.: Mater. Electron. 19 (2008) 369. D. Barreca, G. Bruno, A. Gasparotto, M. Losurdo, E. Tondello, Mater. Sci. Eng. C23 (2003) 1013. C. Mansilla, J.P. Holgado, J.P. Espino´s, A.R. Gonza´lez-Elipe, F. Yubero, Surf. Coat. Technol. 202 (2007) 1256. G.B. Harries, Philos. Mag. 43 (1952) 113. L.M. Zhang, Y.S. Gong, C.B. Wang, Q. Shen, M.X. Xia, Thin Solid Films 496 (2006) 371. G.K. Williamson, W.H. Hall, Acta Metall. 1 (1953) 22.
[31] P. Scherrer, Nachr. Ges. Wiss. Gott. 26 (1918) 98. [32] R. Swanepoel, J. Phys, E: Sci. Instrum. 16 (1983) 1214. [33] V. Petrovsky, B.P. Gorman, H.U. Anderson, T. Petrovsky, J. Appl. Phys. 90 (2001) 2517. [34] S. Tsunekawa, T. Fukuda, A. Kasuya, J. Appl. Phys. 87 (2000) 1318. [35] O. Stenzel, The Physics of Thin Film Optical Spectra: an Introduction, Springer, 2005. [36] J. Tauc, R. Grigrovici, A. Vancu, Phys. Status Solidi B 15 (1966) 627. [37] M.S. Alrobaee, M.G. Krishna, K.N. Rao, S. Mohan, J. Vac. Sci. Technol., A 9 (1991) 3048. [38] M.S. Alrobaee, K.N. Rao, S. Mohan, J. Appl. Phys. 71 (1992) 2380. [39] M.S. Alrobaee, L. Shivalingappa, K.N. Rao, S. Mohan, Thin Solid Films 221 (1992) 214. [40] Y.J. Mao, X.H. Liu, F. Zhang, C.X. Ren, S.C. Zou, Surf. Coat. Technol. 104 (1998) 78. [41] S. Brandstetter, P.M. Derlet, S. Van Petegem, H. Van Swygenhoven, Acta Mater. 56 (2008) 165. [42] F. Zhang, S.W. Chan, J.E. Spanier, E. Apak, Q. Jin, R.D. Robinson, I.P. Herman, Appl. Phys. Lett. 80 (2002) 127. [43] L. Wu, H.J. Wiesmann, A.R. Moodenbaugh, R.F. Klie, Y. Zhu, D.O. Welch, M. Suenaga, Phys. Rev. B 69 (2004) 1254151. [44] S. Tsunekawa, K. Ishikawa, Z.Q. Li, Y. Kawazoe, A. Kasuya, Phys. Rev. Lett. 85 (2000) 3440. [45] A. Kossoy, Y. Feldman, E. Wachtel, K. Gartsman, I. Lubomirsky, J. Fleig, J. Maier, Phys. Chem. Chem. Phys. 8 (2006) 1111. [46] S. Tsunekawa, R. Sahara, Y. Kawazoe, K. Ishikawa, Appl. Surf. Sci. 152 (1999) 53. [47] S. Tsunekawa, S. Ito, Y. Kawazoe, Appl. Phys. Lett. 85 (2004) 3845. [48] S. Deshpande, S. Patil, S.V. Kuchibhatla, S. Seal, Appl. Phys. Lett. 87 (2005) 133113. [49] P. Burroughs, A. Hamnett, A.F. Orchard, G. Thornton, J. Chem. Soc., Dalton Trans. (1976) 1686. [50] B.A. Movchan, A.V. Demchish, Phys. Met. Metall. 28 (1969) 83. [51] J.A. Thornton, J. Vac. Sci. Technol. 11 (1974) 666. [52] J.A. Thornton, Annu. Rev. Mater. Sci. 7 (1977) 239. [53] R. Messier, A.P. Giri, R.A. Roy, J. Vac. Sci. Technol. A2 (1984) 500. [54] R. Messier, R.C. Ross, J. Appl. Phys. 53 (1982) 6220. [55] T. Suzuki, I. Kosacki, V. Petrovsky, H.U. Anderson, J. Appl. Phys. 91 (2002) 2308. [56] R. Gallage, A. Matsuo, T. Watanabe, N. Matsushita, M. Yoshimura, J Electroceram. (2008) 1. [57] N. Ozer, Sol. Energy Mater. Sol. Cells 68 (2001) 391. [58] K.B. Sundaram, P. Wahind, Phys. Status Solidi B 161 (1990) K63. [59] S.Y. Zheng, A.M. Andersson-Faldt, B. Stjerna, C.G. Granqvist, Appl. Opt. 32 (1993) 6303. [60] C.A. Hogarth, Z.T. Al-Dhhan, Phys. Status Solidi B 137 (1986). [61] S. Debnath, M.R. Islam, M.S.R. Khan, Bull. Mater. Sci. 30 (2007) 315.