Local Site Distribution of Oxygen in Silicon-Rich Oxide Thin Films: A

Article pubs.acs.org/JPCC

Local Site Distribution of Oxygen in Silicon-Rich Oxide Thin Films: A Tool to Investigate Phase Separation Davor Ristić,† Mile Ivanda,*,‡ Giorgio Speranza,†,§ Zdravko Siketić,‡ Ivančica Bogdanović-Radović,‡ Marijan Marciuš,‡ Mira Ristić,‡ Ozren Gamulin,∥ Svetozar Musić,‡ Krešimir Furić,‡ Giancarlo C. Righini,⊥,# and Maurizio Ferrari† †

CSMFO Laboratory, CNR-IFN, Via alla Cascata 56/C, Povo, 38123 Trento, Italy Ruđer Bošković Institute, P.O. Box 180, 10002 Zagreb, Croatia § Fondazione Bruno Kessler FBK, Via Sommarive 18, Povo, 38123 Trento, Italy ∥ Medical School, Department of Physics and Biophysics, University of Zagreb, Šalata 3b, 10000, Zagreb, Croatia ⊥ MDF Laboratory, CNR-IFAC, Via Madonna del Piano 10, 50019 Sesto Fiorentino, Firenze, Italy # Centro Fermi, Piazza del Viminale 1, 00184 Roma, Italy ‡

ABSTRACT: Thin films of nonstoichiometric silicon oxide (SiOx with x < 2) have been studied extensively during the past few decades because of their importance in many electronic and optoelectronic applications, and particular attention has been paid to models that can better describe their global structure. Herein, we present a detailed study of SiOx films deposited on silicon(111) and silica substrates using the low-pressure chemical vapor deposition (LPCVD) method by thermal oxidation of silane in an oxygen atmosphere at a temperature of 570 °C. The oxygen and silane flows in the reactor were varied to obtain films with different values of oxygen content x. Ellipsometry and m-line measurements were used to determine the complex refractive index of the deposited films. The oxygen contents in the films were measured by infrared spectroscopy, energy-dispersive X-ray spectroscopy (EDX), and time-of-flight elastic recoil detection analysis (TOF-ERDA). The oxygen contents in the films were also estimated from the measured values of the complex refractive indices using Bruggeman’s effective-medium aproximation (EMA). All of the results were in good agreement, except for those obtained from infrared spectroscopy, which corresponded to systematically higher oxygen contents. This effect was interpreted as being due to an inhomogeneous distribution of oxygen atoms in the films (phase separation). This issue was confirmed by X-ray photoelectron spectroscopy (XPS) analysis of the Si 2p core levels, which showed an almostcomplete phase separation of the silicon-rich oxides into amorphous silicon and silicon dioxide, indicating that the mixture model is the most appropriate for the present films.

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INTRODUCTION Silicon-rich oxides (SiOx, 0 < x < 2) are technologically relevant materials because they are frequently present in a variety of situations such as thin boundary layers between crystalline silicon and thermally grown oxides. The electronic properties of silicon are strongly affected by the presence of these boundary oxides.1 In electronic devices based on complementary metal− oxide−semiconductor (CMOS) technology, the performances are dictated by the quality of the gate oxides considering the continuous reduction of their dimensions to reach larger integrations. SiOx was first used as a replacement for silicon dioxide as an insulating layer in bipolar transistors.2 For this reason, SiOx phases have been the subjects of several studies aimed at establishing their composition and structure.3−5 The electronic properties of SiOx systems also drive their optical properties. After the first detection of visible luminescence in porous silicon,6 silicon-rich oxide started to be used as a material from which nanocrystalline silicon (nc-Si) could be produced by annealing. It was shown that Si nanocrystals © 2012 American Chemical Society

embedded in a transparent dielectric SiO2 matrix are strong light emitters,7 which makes them appealing for optoelectronic applications.8−10 More recently, chemical vapor deposition, sputtering, reactive ion etching, and ion implantation have been utilized to fabricate these systems.11−13 The advantage of producing nanocrystalline silicon from SiOx thin films rather than by other methods stems from the fact that SiOx thin films can be produced by completely CMOS-compatible methods. A few models have been used in the literature to describe the short-range order in SiOx, including the random bonding model (RBM),14 the mixture model (MM),15 and the intermediate model (IM).16 The effective use of one model seems to depend on the specific fabrication technique of the films. Received: February 6, 2012 Revised: March 27, 2012 Published: April 5, 2012 10039

dx.doi.org/10.1021/jp301181y | J. Phys. Chem. C 2012, 116, 10039−10047

The Journal of Physical Chemistry C

Article

We focused our investigation on silicon-rich oxide thin films deposited by the low-pressure chemical vapor deposition (LPCVD) technique and produced several samples with varying oxygen contents. We studied the chemical and atomic structures of the deposited films in great detail using different spectroscopic techniques. The bulk oxygen content was directly determined using time-of-flight elastic recoil detection analysis (TOF-ERDA) and energy-dispersive X-ray (EDX) spectroscopy. Surface characterizations in terms of chemical composition and atomic concentrations were performed using X-ray photoelectron spectroscopy (XPS). The experimental results were then compared with estimations of the oxygen content in the SiOx films calculated from complex refractive index measurements using Bruggeman’s effective-medium theory and from IR absorption spectra. Conclusions were drawn about the siliconrich structure and its relation to the oxygen content in the films.

IV ellipsometer was used for measurements at a wavelength of 405 nm for both the thickness and the complex refractive index of thinner films (samples S3−S5). For films thicker than 100 nm (samples S1, S2, and SR), ellipsometry could not be used for the determination of the film thickness because the absorption at 405 nm was too large; instead, we used a Metricon 2010 m-line apparatus equipped with a rutile prism and laser diodes at 1319 and 1542 nm, where the sample absorption is negligible. A JEOL JSM 700F scanning electron microscope fitted with an Oxford Inca EDX apparatus was used for surface imaging and chemical analysis of the produced films. TOFERDA measurements were made using two different ion types: 26 MeV 197Au ions with an angle of incidence of 20° onto the surface of the samples and 20 MeV 81Br ions with a 10° incidence angle. The detection angle of ions recoiled from the sample was 37.5° with respect to the direction of the incident beam. The beam diameter at entry was 2 × 2 mm2, and the beam current was below 1 nA in all cases. Infrared spectra were recorded using a Perkin-Elmer Spectrum GX Fourier transform infrared (FTIR) spectrometer. XPS measurements were performed using a Scienta ESCA 200 analyzer (Gammadata Instrument AB, Uppsala, Sweden) equipped with a monochromatized Al Kα X-ray source. The core lines were acquired at 150 eV pass energy, corresponding to an energy resolution of 0.35 eV taking into account the effects of charge compensation. All spectra were aligned using the pure Si 2p peak as a reference. After linear background subtraction, core-line peak fitting was performed using Gaussian components, each of which was associated with a specific chemical bond formed by oxygen with silicon in a specific oxidation state.

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EXPERIMENTAL PROCEDURES Thin silicon-rich oxide films were deposited in an LPCVD reactor at a temperature of 570 °C using silane gas diluted with argon (SiH4/Ar) and oxygen as the reactant gases. The LPCVD reactor was of the 5-in.-diameter horizontal type with three zones and a total length of 2 m.17 The flow of SiH4/Ar in the reactor was kept constant at 260 cm3/min, whereas the volume fraction of silane in the SiH4/Ar mixture was varied from 2% to 26%. The oxygen flow in the reactor was varied from 9 to 27 cm3/min. In addition, a reference sample of amorphous silicon was deposited by injecting only the silane gas into the reactor. The depositions were made on both silica and silicon wafers, but all of the results reported herein refer to films deposited on silicon substrates, unless noted otherwise. The deposition parameters of the samples are given in Table 1. An AUTO EL

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RESULTS Table 2 reports the thicknesses of the SiOx thin films measured by the different techniques. The deposition rates were calculated using the values of thickness measured by ellipsometry and m-line spectroscopy. The decrease of the deposition rate with increasing oxygen-to-silane flow ratio in the reactor for samples SR and S1−S3 is evident. This can be explained by considering that, with increasing oxygen flow in the reactor, the molecules of silane and oxygen tended to react before reaching the substrates; because of this effect, silicon and silicon dioxide molecules were preferentially deposited at the entrance of the LPCVD reactor tube, and only a fraction of the molecules were deposited onto the substrates, which were placed halfway down the tube. On the other hand, samples

Table 1. Deposition Parameters of the Pure Silicon (SR) and SiOx Films: v is the flow of oxygen in the reactor, V(SiH4)/ V(Ar) the ratio of the SiH4 and Ar volumes in the gas mixture and t is the duration of the depositions sample

v(O2) (cm3/min)

V(SiH4)/V(Ar) (%)

t (min)

SR S1 S2 S3 S4 S5

0 9 18 27 27 27

26 26 26 26 14 2

300 300 300 300 120 166

Table 2. Thicknesses and Complex Refractive Indices of the Pure Silicon (SR) and SiOx Filmsa SR

S1

d (nm) vdep (nm/min)

515 ± 8 1.72 ± 0.03

m-Lineb 378 ± 4 1.26 ± 0.01

d (nm)

−

−

n k

4.3 ± 0.1 1.92 ± 0.05

3.8 ± 0.1 1.95 ± 0.05

ω (cm−1)

1012 ± 3

S2 127 ± 2 0.42 ± 0.01 TOF-ERDA 150 ± 30 Ellipsometry 3.2 ± 0.1 1.53 ± 0.01 IR Absorption 1026 ± 3

S3

S4

70 ± 10 0.23 ± 0.03

Ellipsometry 60 ± 20 0.5 ± 0.2

S5 70 ± 20 0.4 ± 0.1

57 ± 7

38 ± 7

50 ± 10

2.4 ± 0.1 0.62 ± 0.02

2.4 ± 0.1 0.46 ± 0.05

1.46 ± 0.01 0

1037 ± 3

1037 ± 3

1060 ± 3

a d is the film thickness; vdep is the deposition rate; n and k are the real and imaginary parts, respectively, of the complex index of refraction at 405 nm; and ω is the Si−O stretching peak position in the IR absorption spectrum. bm-line measurements were performed on samples deposited on silica substrates.

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S3−S5 were all produced at the same oxygen flow rate, namely, 27 cm3/min. However, with increasing dilution of silane, the reaction between silane and oxygen had less probability to occur before the gases reached the substrate, so the deposition rates for samples S4 and S5 were higher than that for the sample S3. The values of thickness reported in Table 2 for both the ellipsometry and m-line measurements indicate the actual variations across the 2-in. length of samples; actually, the films were deliberately deposited to have a slight gradient of thickness (approximately 10 nm over 1 cm) when moving from the center toward the edge of the substrate. This was achieved by placing the substrate wafer in the reactor between two dummy wafers so that the gases became diluted when they arrived at the center of the substrate wafer. Therefore, the deposition rate was slightly higher at the edge of the substrates when compared to the center. This allowed us to measure the complex refractive index using ellipsometry, as explained below. The complex refractive indices of the SiOx films at the wavelength of 405 nm were determined using ellipsometry. Samples SR, S1, and S2 were thick enough that the incident light was absorbed in the sample, preventing any reflection at the substrate−film interface. This was confirmed by making measurements on the films deposited on both silicon and silica substrates. The measured quantities Δ and Ψ were the same regardless of the type of substrate. The complex refractive indices of these samples were determined directly from the measured quantities Δ and Ψ using the standard ellipsometric formula ⎡ ⎛ r − 1 ⎞2 ⎤ ⎟ ⎥ n2 = sin 2 θ ⎢1 + tan θ ⎜ ⎝ r + 1⎠ ⎦ ⎣

Figure 1. Δ−Ψ plots of samples S3−S5 deposited on silicon substrates. The circles represent the experimental ellipsometric data measured across the surface of the samples, and the solid line is the theoretical fit. The complex refractive index of the film was used as a fit parameter and is indicated inside each panel.

the theoretical curve. The small discrepancies are due to the fact that the complex refractive indices were not constant across the samples but varied slightly depending on the measurement position on the sample. This small variation is responsible for the uncertainties in the determination of the refractive indices presented in Figure 1. For sample S5, the experimental data exactly follow the silicaon-silicon curve, which identifies sample S5 as silicon dioxide. Refractive index measurements were made by m-line spectroscopy on samples S1 and S2 and reference sample SR at the wavelengths of 1319 and 1542 nm. Samples S3−S5 were too thin to support waveguiding modes at the above-mentioned wavelengths and could not be characterized by m-line spectroscopy. The samples were analyzed using a rutile prism, which can couple light into modes that have an effective index of refraction between 2.4 and 1.55 for the transverse electric (TE) modes and between 2.2 and 1.35 for the transverse magnetic (TM) modes. Because the refractive indices of the samples were higher than the upper limit of effective refractive indices that can be detected by the prism, we were not able to detect all of the modes present in the waveguide and detected only those modes falling in the effective refractive index range covered by this prism. The values of the effective indices reported in Table 3 are mean values measured across the surfaces of the samples. The measurements were made along the profiles of the samples: For every point, the thickness and refractive index were calculated, and the analysis of the results allowed us to determine the uncertainties in the thicknesses and refractive indices. Table 2 also shows the mode assignments of the effective indices, made by taking into consideration all possible combinations of the assignments and choosing only those leading to reasonable values of both refractive indices and thicknesses. The dispersion of semiconductor and dielectric materials was also taken into account; in particular, the refractive index of the film at 1542 nm should always be slightly lower than the refractive index at 1319 nm (actually, this decrease for our

(1)

where θ is the angle of incidence of the light on the sample (which, in our case, was 70) and r = tan ΨeiΔ. Thus, complex refractive indices of n = (4.34 ± 0.1) + i(1.97 ± 0.05), n = (3.8 ± 0.1) + i(1.95 ± 0.05), and n = (3.2 ± 0.1) + i(1.53 ± 0.01) were obtained for samples SR, S1, and S2, respectively. For samples S3−S5, we exploited the gradient of the film thickness across the sample. In these cases, the complex refractive index was determined by measuring the Δ and Ψ profiles from the edges toward the center of the samples deposited on silicon substrates. The model used for the Δ and Ψ calculation was the simple single-film model. In that case, the reflection coefficient of the thin film−substrate system is given by s s iδ r12 + r23 e

r=

s s iδ r23e 1 + r12 p p iδ r12 + r23 e p p iδ r23e 1 + r12

(2)

where rs12 and rp12 are the Fresnel reflection coefficients for the s-polarized (electric field of light polarized perpendicular to the plane of incidence) and p-polarized (electric field of light polarized in the plane of incidence) signals, respectively, on the air−film interface; rs23 and rp23 are the corresponding Fresnel coefficients on the film−substrate interface; and δ is the phase difference between light reflected by the air−film and film− substrate interfaces. For each film, the real and imaginary parts of the complex refractive index were determined by fitting the measured data to the theoretical Δ−Ψ function. The Δ−Ψ plots of samples S3−S5 are shown in Figure 1. It can be seen that, for samples S3−S5, the experimental data can be fitted reasonably well with 10041



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Table 3. Effective Refractive Indices of the Modes Detected Using m-Line Spectroscopy and a Rutile Prism, As Well As Mode Assignments, Thicknesses d, and Refractive Indices n at 1319 and 1542 nm Calculated from the Effective Refractive Indices of the Modesa effective refractive indices 1319 nm TE

TM

TE

SR

2.198 TE2

1.506 TM2 1.700 TM1 1.456 TE0

1.650 TE2 2.145 TE1 1.915 TE0

S1 S2 a

1542 nm

sample

2.045 TE0

TM

1.462 TM1

n(1319)

n(1542)

d (nm)

3.74 ± 0.01

3.71 ± 0.01

515 ± 8

3.55 ± 0.02

3.52 ± 0.02

378 ± 4

2.96 ± 0.01

2.93 ± 0.01

127 ± 2

Measurements were performed on samples deposited on silica substrates.

Table 4. Oxygen Content x of Deposited SiOx Films, According to Different Analytical Techniques S1

S2

S3

S4

S5

0.81 ± 0.08

0.88 ± 0.08

2.00 ± 0.08

TOF-ERDA x

0.10 ± 0.04

0.28 ± 0.04 EDX

x

0.12 ± 0.02

x Im(x)/Re(x)a

0.11 ± 0.06 0.08

x

0.08 ± 0.02

x

1.20 ± 0.06

0.30 ± 0.05 >0.5 Ellipsometry + Bruggeman EMA 0.36 ± 0.05 0.84 ± 0.06 0.08 0.03 m-Line + Bruggeman EMA 0.37 ± 0.01 IR Absorption 1.43 ± 0.06 1.62 ± 0.06

>0.8 0.89 ± 0.07 0.04

2.00 ± 0.02 0

1.62 ± 0.06

2.00 ± 0.06

a

Ratio of the real and imaginary parts of the oxygen content calculated using the Bruggeman EMA from the ellipsometric data, which should be equal to zero if application of the Bruggeman EMA is valid.

samples was found to be around 0.03). Consider, as an example, sample SR: If we assume that, with the used prism, the visible modes at 1319 nm are the TE3 and TM3 modes, then there should be two TM modes (TM2 and TM3) visible at 1542 nm. Because only one mode was detected at 1542 nm, we have to conclude that the initial assignment is not correct. If, on the other hand, we assume that the modes visible at 1319 nm are TE1 and TM1, then the condition to fulfill to have a TE1 effective index equal to 1.65 at 1542 nm would require that the difference of the refractive indices, n(1319) − n(1542), be equal to 0.25, which is clearly too large. Finally, if we assign the two modes at 1319 nm to be TE2 and TM2, the calculated refractive index difference, n(1319) − n(1542), must be on the order of 0.03, which is fully reasonable. Thus, for each sample, there was only one assignment that was self-consistent. Table 4 reports the values of oxygen concentration x determined by various methods. The TOF-ERDA depth profiles of the deposited films presented in Figure 2 clearly show that the oxygen content in the films increased with increasing oxygen flow in the reactor. The thicknesses were determined by assuming the density of the SiOx films to be 2.3 g/cm3, because the densities of both pure silicon and pure silicon dioxide are close to this value, and by estimating the mean molecular masses of the films from the measured oxygen contents in the films. In the case of sample S1, a slow continuous decrease of the oxygen content in the sample profile was observed. For this reason, we were unable to estimate the thickness of the film with sufficient precision, and only an average value of the oxygen concentration is given in Table 4. TOF-ERDA also showed small amounts of carbon ( 0.5 for sample S3 and x > 0.8 for sample S5. The oxygen content was also calculated starting from the measured complex refractive index values using Bruggeman’s effective-medium approximation (EMA),20 which is often used for the analysis of thin SiOx films.21,22 Bruggeman’s EMA is based on the assumption that the dielectric constant of a mixture of two compounds can be calculated from the dielectric constants of the two components as ε −ε ε −ε fA A + fB B =0 εA + 2ε εB + 2ε (4)

hydrogen ( x), the oxygen content shown by the IR spectra would be equal to y rather than x. This effect of clustering on the position of the Si−O stretching peak was recently demonstrated by Cueff et al.26,27 The real microscopic structure of SiOx is still a subject of debate. This structure is very important for the applicability of SiOx films as the basis for the production of light-emitting materials (nanocrystalline silicon) because the light-emitting properties of such materials are correlated with the microscopic structure of the amorphous SiOx matrix.28 Two models are commonly used to describe a SiOx network: the mixture model15 and the random bonding model.14 The mixture model approximates the SiOx structure as a mixture of two distinct phases, one of which is richer in oxygen than the other. In the random bonding model, on the other hand, the structure is described as a single-phase system consisting of a number of homogenously distributed randomly bonded Si−(SinO4‑n) tetrahedra (0 < n < 4). In this case, there should not be any significant variations in the number of oxygen atoms in the neighborhood of any given silicon atom, so the oxygen content determined from the position of the Si−O stretching mode should be equal to the macroscopic oxygen content. Recently, the intermediate model (IM) was introduced by Novikov and Gritsenko.16 The IM model assumes smooth variation of the chemical composition at the boundaries between silicon clusters and SiO2 matrix, leading to a continuous distribution of the oxygen content x. The actual structure seems to be greatly determined by the deposition procedure.29 In some works, the structures of SiOx films obtained by radio-frequency sputtering30 and physical evaporation31 were claimed to correspond to the random bonding network model, whereas SiOx films obtained by magneto-sputtering,32 plasma-enhanced chemical vapor deposition,33 and a number of commercially available bulk silicon monoxides34 have been assigned to the mixture model. The intermediate model was used to describe SiOx layers prepared by LPCVD from SiH4 and a N2O source at 750 °C.16 In our case, the above-mentioned discrepancy in the oxygen contents determined from the IR spectra and from the other methods leads to the conclusion that our structure does not correspond to the random bonding model. This finding can

−1

where ω is the Si−O peak position in cm . The entire infrared spectrum is shown in Figure 4a, with an enlargement of the

Figure 4. IR spectra of samples S1−S5 deposited on silicon substrates: (a) entire IR range from 0 to 10000 cm−1, (b) an expanded view of the spectra from 400 to 1300 cm−1, (c) Si−O stretching mode after subtraction of the background.

range from 400 to 1300 cm−1 in Figure 4b; in Figure 4c, the spectrum of the silicon substrate was subtracted from the recorded spectra. These spectra show that there is a clearly visible shift of the position of the Si−O stretching mode peak with an increase of the oxygen content in the samples. The position of the Si−O stretching mode peak and the oxygen content x obtained using eq 6 are reported in Tables 2 and 4, respectively. It can be seen that the oxygen contents obtained in this manner were much larger than those obtained by TOFERDA, EDX, and EMA. A possible explanation for this discrepancy is given in the following section.

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DISCUSSION Figure 5 summarizes the results of the measurements by plotting the values of the oxygen contents x obtained by TOF-ERDA, EDX, IR, and Bruggeman EMA in the visible and infrared ranges. It can be seen that there is a very good agreement between the TOF-ERDA, Bruggeman EMA, and EDX results everywhere these measurements were made. The oxygen contents obtained from the IR spectra, however, appear to be much higher than all of the other values. This can be explained considering that the position of the Si−O stretching peak depends not on the real oxygen content 10044



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Raman spectroscopy of LPCVD-deposited silicon films revealed that, at temperatures in the range of 500−570 °C, the films are amorphous, whereas at 590 °C, the films are polycrystalline.37 Moreover, the width of the Raman band at around 480 cm−1 was found to decrease with increasing deposition temperature,37 indicating a gradual relaxation and ordering of the structure when moving from 500 to 570 °C.38 Therefore, a temperature of 570 °C, although not sufficient to induce crystallization, is high enough to produce an amorphous homogeneous structure in which the variations in the bond angle or length are very low. In the spectrum of sample SR, there was also a peak at around 102.8 eV corresponding to oxidized silicon that is due to the surface oxidation of the sample, occurring in the time between the deposition and the XPS measurements. (The probe depth of XPS was around 6 nm.) The Si 2p core spectra were deconvoluted into Gaussian components using the following procedure. First, the spectrum of the Si0 peak from sample SR was fitted to two Gaussians corresponding to the two spin−orbit Si0 levels. The width of the two Gaussians was 0.3 eV, and the spin−orbit splitting was found to be 0.6 eV. Sample S7 (pure silica) was fitted using a single 0.7 eV broad Gaussian corresponding to the Si4+ line. All of the other spectra were fitted by supposing a constant energy shift of the Si+, Si2+, and Si3+ peaks from the midpoint between the two spin−orbit split components of the Si0 peak and the Si4+ peak. The intensities of the different peaks were used as fit parameters. Also, the position of the Si4+ component and the midpoint between the Si0 components were varied to optimize the fit. The widths of the Gaussian components and their energy separations were kept constant. It was found that, with an increase of oxygen in the samples, the entire spectrum of the Si 2p core levels shifted by about 0.4 eV, with the position of the Si4+ peak varying between 102.9 and 103.3 eV and the midposition of the Si0 levels varying between 99.05 and 99.45 eV. This shift might be due to changes in the Fermi level with increasing oxygen content in the samples.15 In the random bonding network model, the silicon and oxygen atoms are homogeneously dispersed in the SiOx structure. According to this hypothesis, the intensities of the different components of the Si 2p core level in an SiOx structure should be given by15

be also ascertained using X-ray photoelectron spectroscopy (XPS) measurements. XPS measurements of the Si 2p core levels are presented in Figure 6. In the XPS spectrum of sample SR, the spin−orbit

Figure 6. XPS spectra of samples S1−S5 deposited on silicon substrates. In the left panels, the circles represent the experimental XPS spectra, the dashed lines are the individual components of the Si 2p core level corresponding to different oxidation states of silicon, and the solid line is their sum. In the right panels, the circles are the relative intensities of the individual components to the Si 2p core level, the squares are the components calculated according to the randombonding model, and the triangles are the components calculated according to the mixture model. The parameters used in the calculations are given inside the panels.

In(x) =

4! (x /2)n (1 − x /2)4 − n (4 − n) ! n!

(7)

where n is the oxidation state of silicon. In the mixture model, phase separation of the SiOx structure into two components, one richer in silicon and the other richer in oxygen, is assumed y−x x−z SiOx → SiOy + SiOz y−z y−z (8)

splitting of the Si0 line of the Si 2p level around 99.05 eV is visible. The spin−orbit splitting is clearly visible in the presence of crystalline structures.35 In the case of amorphous Si and SiOx, the variations of the bond angle in the structures, defects, and loss of stoichiometry lead to a broadening of the spin− orbit components that masks the spin−orbit splitting.8,15,29 The Raman spectra of our LPCVD SiOx samples deposited on silica substrates showed no sign of the crystalline silicon peak at around 520 cm−1.36 This apparent contradiction with the presence of spin−orbit splitting in the XPS spectra can be explained considering the different descriptions that these two techniques provide. Raman spectra account for phonon vibration occurring at the molecular or supramolecular scale, whereas photoelectron spectra mirror the atomic level. This means that our samples exhibit an amorphous structure characterized by a low number of voids/defects and by a small variation of the bond angle of the tetrahedral units. This can also be explained by the choice of deposition temperature.

Then, the relative contribution of each component in the XPS spectrum is given by In(x , y , z) =

⎛ y ⎞n x−z 4! 4−n ⎜ ⎟ (1 − y /2) y − z (4 − n) ! n! ⎝ 2 ⎠ y−x ⎛ z ⎞n 4! 4−n ⎜ ⎟ (1 − z /2) + y − z (4 − n) ! n! ⎝ 2 ⎠ (9)

In the right part of Figure 6, the fits of the experimental data using eq 9 are shown. The overall oxygen content x was determined from the experimental intensities of the different Si 10045



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2p core-level components using the equation15 x = 0.5I (Si+) + I(Si 2 +) + 1.5I (Si 3 +) + 2I(Si4 +)

optoelectronics” (2010-2013) of the Provincia Autonoma di Trento and of project 098-0982904-2898 of the Ministry of Science, Education and Sports of the Republic of Croatia.

(10)

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where the overall intensity of the Si 2p core level is normalized to unity. The oxygen content in each sample determined in this manner (presented in Figure 6) was found to be slightly greater than the oxygen contents determined using the other methods (TOF-ERDA, ellipsometry, EDX). This can be explained considering that XPS is a surface probing technique and, after deposition, the surface can undergo further oxidation with respect to the bulk. Actually, the oxygen contents measured by XPS were always found to be greater than those in the bulk (i.e., x) by about 0.4−0.5. The oxygen contents in the two separated phases y and z were used as fit parameters. The results of the fit are presented on the right of Figure 6. From this analysis, it is evident that, whereas the random-bonding network model does not give a satisfactory agreement with the experimental data, the mixture model is in almost perfect agreement. In fact, all of the samples were found to exhibit almost complete separation into amorphous silicon (z = 0) and silicon dioxide (y = 2). Actually, the fits always gave a value of y > 1.7 and a value of z < 0.22.

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CONCLUSIONS Thin SiOx films were deposited using the LPCVD method to obtain values of oxygen content 0 < x < 2. Different analytical techniques were used to measure optical properties and to determine the oxygen content. TOF-ERDA and EDX spectroscopy were used to determine the oxygen content directly. EDX spectroscopy was found to be applicable only for samples thinner than 200 nm and using an electron-beam acceleration voltage of less than 3 kV. The complex refractive indices of the films were measured using ellipsometry at a wavelength of 405 nm and using m-line spectroscopy at 1319 and 1542 nm. Bruggemman’s effective-medium approximation was used to determine the oxygen content from the measured complex refractive indices and was found to be in very good agreement with the TOF-ERDA and EDX results. On the contrary, the oxygen contents obtained from the position of the Si−O stretching peak in the infrared absorption spectra were found to be much larger than expected. This was explained by oxygenrich SiOy and silicon-rich SiOz (y > x > z) phase separation. This assumption of phase separation was also ascertained using XPS. The XPS spectra of the Si 2p core levels clearly show an almost-complete phase separation of the SiOx films into amorphous silicon and silica. These results, which indicate that the structure of our films can be better described by the mixture model, can be of general utility to sustain the development of detailed analysis procedures aimed at improving the quality of silicon-rich oxide films, where the oxygen content and its distribution play a crucial technological role.

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REFERENCES

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Phone: 385 1 456 0928. Fax: 385 1 4680 112. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research was performed within the framework of the NSBMO research project “Novel silicon based materials for 10046



Article

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Local Site Distribution of Oxygen in Silicon-Rich Oxide Thin Films: A

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