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On the relationship between the structure and the barrier performance of plasma deposited silicon dioxide-like films
Surface & Coatings Technology 204 (2010) 4012–4017
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
On the relationship between the structure and the barrier performance of plasma deposited silicon dioxide-like films Anna Maria Coclite a,⁎, Antonella Milella a, Riccardo d'Agostino a, Fabio Palumbo b a b
Dipartimento di Chimica, Università degli Studi di Bari, Italy Istituto di Metodologie Inorganiche e dei Plasmi – CNR, Bari, Italy
a r t i c l e
i n f o
Article history: Received 14 September 2009 Accepted in revised form 14 May 2010 Available online 16 June 2010 Keywords: Barrier coatings FT-IR Plasma deposition Silicon dioxide
a b s t r a c t Silicon dioxide-like barrier films were deposited by plasma enhanced chemical vapor deposition from different siloxane and silane precursors. The variation of the precursor was investigated as a route to obtain silicon dioxide-like films with different structures, densities and hence barrier performances. Although the films were characterized by the same elemental composition, some differences in film density and porosity were evidenced from optical properties measurements and from the shift of the SiOSi infrared absorption band position. These differences were correlated with film microstructure and in turn with barrier performances. The results confirmed that films with high density and low porosity performed better as single inorganic barrier layers for food-packaging. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Over the past few decades, interest in transparent, thin SiO2-like coatings have increasingly grown for several applications like optics [1], interlayer dielectrics [2], and barrier films for the food-packaging and medical devices industries [3]. As gas and vapor diffusion barriers, such coatings have desirable properties, such as transparency, recyclability, and microwave use and are superior in this regard to the thin metal coatings (generally aluminium-based). The use of thin SiO2-like barrier films on polymer substrates is an interesting and challenging research field because the matching of thin inorganic barrier film to plastic substrates creates a new material which is impermeable and at the same time flexible, lightweight and environmental compatible. Although extensively studied [3–8], the SiO2-like film deposition for barrier applications still offers some interesting aspects, concerning the correlation between film properties (e.g. density and porosity) and performances. Previous reports on plasma deposited barrier coatings from organosilicon precursors mainly correlated the film chemical composition with the barrier properties demonstrating that the permeability can be significantly improved minimizing the content of organic moieties and of silanol groups in the films. For example, the Oxygen Transmission Rate (OTR) through a PET substrate can be lowered from 2–5 cm3 m− 2 day− 1 atm− 1 to 0.2–0.5 cm3 m− 2 day− 1 atm− 1 [6,7] and the water vapor transmission rate (WVTR) to 0.5 g m− 2 day− 1 ⁎ Corresponding author. Dipartimento di Chimica, Università degli Studi di Bari, Via Orabona, 4, 70126 Bari, Italy. Tel.: +39 0805443433; fax: + 39 0805443405. E-mail address: [email protected] (A.M. Coclite). 0257-8972/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2010.05.024
[4,5] at 23 °C and 50% of relative humidity (RH). It was demonstrated that the presence of organic and silanol groups deteriorates the barrier properties both because the SiOSi network is interrupted and because they increase the solubility of the permeant molecules in the film favouring their transmission through the barrier film. Beside film chemistry, other factors such as film defects, film adhesion to polymeric substrate and substrate surface defects were demonstrated to influence the film permeability to gas and vapor. da Silva Sobrinho et al. [7] showed that the presence of unwanted but inevitable nano-, microscopic defects and pinholes in the coating, which may result from dust particles on the substrate surface, from geometric shadowing and stress during film growth at sites of high surface roughness, limits the minimum value achievable of OTR. The smaller the defect density, the lower the permeation rate is. Roberts et al. [8], by means of a model of the permeation process in evaporated silicon oxide gas barrier films, identified such defects as distinguishable in three types: macro-defects (cracks N 1 nm), nano-defects (pinholes b 1 nm) and amorphous oxide lattice defects (interstice at grain boundary b 0.3 nm). The barrier properties have never been directly correlated with the film microstructure in terms of density and porosity although extensive studies in literature show that silica coatings can be significantly different in the atomic structure depending on the particular deposition process used: both silicon and oxygen atoms can be found in different environments, allowing the presence of homoatomic bonds in disordered chains. Dense nearly stoichiometric silicon dioxide films can be achieved by Ion Beam Assisted Deposition (IBAD) [9], increasing the substrate temperature [10] or biasing the substrate during deposition [11]. Lefevre et al. [9] studied and modeled the microstructural densification process of silicon dioxide
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thin films induced by medium-energy particles bombardment. They considered the description of the network topology in terms of rings of nearest-neighbour bonds, where an n-membered ring is defined as the shortest closed path of n Si–O bonds originating on a given silicon atom. According to this description, in amorphous silica, the structure is statistically dominated by five, six and seven-fold rings. In deposited silicon dioxide-like films, the presence of small rings (three- and fourmembered) is associated with stress and defects in the network, while large rings (e.g. nine-membered) corresponds to film porosity. Ion bombardment induces the structural densification of the SiO2-like films because small rings rearrange to produce larger and less strained ones and too large rings turn into five- six- and seven-membered ones, reducing the free volume and causing the density to increase. They concluded that the densification process can be identified as a relaxation mechanism resulting in silicon dioxide-like films with low stress and low defect concentration. The microstructure and in turn the properties of the SiO2-like films can be tuned by changing the plasma parameters [12] (e.g. O2 flow rate, power, and bias) and/or by choosing a suitable monomer [13,14]. Hegemann et al. [12] showed for HMDSO/O2 plasmas that the power input per gas flow (W/F) controls the chemical composition of the film, determining the discharge fragmentation, while the mechanical properties are mostly influenced by the potential drop across the plasma sheath. High W/F values, such as the ones required to deposit SiO2-like films, are indicative of a monomer-deficient regime in which the precursor structure becomes more and more negligible due to strong fragmentation. Different monomers were previously used to deposit SiO2-like films resulting in similar chemical composition but different deposition rate [13,14]. This is believed to affect the microstructure and in turn the performance of the SiO2-like films. At the authors knowledge the literature does not report the correlation between the structure of the monomer and the barrier performances. The aim of the present paper was to understand whether the monomer structure could affect the barrier performance of SiO2-like coatings deposited in highly oxidative regimes and high fragmentation conditions. Hence four different monomers were selected as follows: (i) tetraethoxysilane (TEOS) and hexamethyldisiloxane (HMDSO) as usual source of SiO2-like films [3–6,15] and (ii) divinyltetramethyldisiloxane (DVTMDSO) and allyltrimethylsilane (ATMS) to compare film structure derived from siloxane and silane. Furthermore, DVTMDSO and ATMS with vinyl and allyl functional groups are highly attractive precursors for inorganic/organic multilayer barrier film deposition [16]. 2. Materials and methods Silicon dioxide-like films were deposited in a capacitively coupled asymmetric parallel plate reactor from different precursors [17]. The upper electrode was connected to a 13.56 MHz power supply (RFPP) through an automatic LC matching network. The reactor was evacuated by means of a turbomolecular/rotary pumping system and the pressure was monitored with a MKS capacitive gauge. Discharge conditions were chosen to obtain C-free coatings with stoichiometry close to SiO2, thus high RF power (i.e. 250 W) and high O2 to monomer ratio (i.e. N20) were used. The deposition parameters are summarized in Table 1. Because of the different chemical structures of the four precursors used, the O2/monomer for each deposition was chosen largely higher than the stoichiometric value obtained from the complete oxidation reaction. Table 2 shows the chemical structures of the precursors used for the depositions. Film chemical and optical characterization was performed using polished Si as substrate, while barrier performance tests were accomplished depositing the coatings on 125-μm-thick poly(ethylene terephthalate) PET (Melinex® ST 504 manufactured by DuPont Teijin). The morphological analysis was performed on both kinds of substrates for all the films deposited.
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Table 1 Experimental parameters utilized for film deposition, the corresponding deposition rates, refractive indices (at 633 nm) and surface RMS roughness (on Si). The RF power was fixed at 250 W. Monomer
O2/mon
Ar flow rate (sccm)
Pressure (mTorr)
Dep rate (nm/min)
n
RMS (nm)
TEOS DVTMDSO HMDSO ATMS
22 14 20 21
44 200 200 150
100 150 200 100
48 ± 4 119 ± 11 69 ± 3 7±1
1.457 1.458 1.461 1.444
0.13 ± 0.06 0.49 ± 0.04 0.60 ± 0.16 0.76 ± 0.09
The film thickness was measured by means of a profilometer (TENCOR α-step 500). All the films were grown up to 100 nm in thickness. Film chemical composition and structure were investigated by FTIR spectroscopy (BRUKER, Equinox 55). Spectra were recorded from 400 to 4000 cm− 1 in absorbance mode at 4 cm− 1 resolution. In order to minimize effects of water vapor and carbon dioxide, the spectrometer was purged with nitrogen for 15 min between each measurement. Baseline correction was applied to the spectra and they were then normalized to the film thickness. The SiOSi stretching band was fitted using two Gaussian components using the “Fit multi-peaks” procedure of OriginLab software. Atomic percentage composition of the films was determined by Xray photoelectron spectroscopy (XPS) (Thermo VG Scientific). Low and high-resolution spectra were acquired using a monochromatic Al Kα radiation (1486.6 eV). The take off angle was 53°. In order to remove carbon contamination from the sample surface, a 10 s sputter cleaning with 1 KeV, 500 nA Ar+ ions was carried out before the analysis. During the XPS analysis, the sample charge was compensated by a 1 eV electron beam at high neutralization current by means of a flood gun. Surface film morphology was investigated by Atomic Force Microscopy (AFM, AutoProbe CP, Thermomicroscope). Images were acquired in non-contact mode using conical gold-coated silicon tips. Root-mean-square (RMS) roughness was measured on 5 μm × 5 μm surface areas averaging 15 measurements for each sample. Ex-situ refractive index measurements were carried out through variable angle spectroscopy ellipsometry (VASE, JA Woollam M2000). The measurements were done at three different angles (65°, 70° and 75°) in the wavelength range of 200 nm to 1000 nm. The applied optical model consisted of four components: the silicon substrate, the native SiO2 layer, the bulk film, and the surface roughness layer. The bulk components were modeled by the Cauchy function while the top-layer using the Bruggeman Effective Medium Approximation (BEMA). The BEMA layer consisted of 50% bulk film and 50% voids. The model also incorporated possible thickness inhomogeneity within the sampled area [18]. Table 2 Chemical structures of the monomers used for the depositions.
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Water vapor transmission rate (WVTR) of the barrier films on PET was measured on a 100 cm2 active area, at 25 °C and 98% of relative humidity (RH), using a MOCON Permatran-W. The WVTR measured through the bare Melinex was 4.46 g m− 2 day− 1. 3. Results and discussion Inorganic silicon dioxide films were deposited in high fragmentation regime (i.e. high oxygen flow rate and high RF power). Varying the precursor structure the oxygen flow rate was changed accordingly, in order to completely oxidize the organic moieties. The ratios were deduced from preliminary experiments aimed at obtaining Cfree polymers at low pressure. Table 1 shows the deposition conditions exploited for this study. XPS elemental composition analysis showed atomic percent of carbon less than 1% and O/Si ratio near 2 for all deposited films. Film stoichiometry was further investigated by curve fitting of the Si2p signal. The curve fitting was accomplished by first subtracting the spin-orbit splitting contribution [19] and then the signals were regressed in the SiyO4 − y components, with y from 0 to 4, depending on the number of Si–O or Si–Si bonds per silicon atom [20]. The least squares regression to the signals was found by using only two components of FWHM around 2 eV [19]. For all the films, the main component was positioned at 103.4 eV, and thus associated with SiO4 groups. The second contribution, centered at 101.9 eV was attributed to Si2O3 groups and represented the 0.5–1.4% of the total Si 2p area. Similar fitting profiles were found by Alfonsetti et al. [20] for films with SiO1.97 stoichiometry. No significative differences were found between the five samples confirming that they are very close in stoichiometry. Fig. 1 shows the WVTR measured for the deposited films. Different performances were detected notwithstanding the similarity in chemical composition evidenced. Indeed only the films deposited from TEOS and DVTMDSO showed good barrier properties, reducing of one order of magnitude the transmission rate through the bare substrate. In order to explain these differences, the refractive indices were measured. Due to the absence of carbon, the measured refractive index basically results from the density and the fraction of voids in the films: high n values can be regarded as indicative of high film density. In the present case, being the measurements carried out ex-situ, the n values can be affected by the water vapor adsorbed in the pores: generally this phenomenon results in an increased refractive index with respect to the in-situ measurements due to the increase of the pore refractive index from 1 of void to 1.332 of water [18]. Analyzing the indices reported in Table 1, it should be noted that the silicon dioxide films deposited from TEOS, DVTMDSO and HMDSO have similar refractive indices, very close to the one characteristic of silica glass (1.46). Despite this similarity, the WVTRs for these samples are different. Significative differences in the refractive indices can be
Fig. 1. Histogram of the WVTR measured at 98% relative humidity and 25 °C for the bare PET substrate and through the film obtained from the different precursors.
appreciated for the film deposited from the ATMS. It shows a low n (1.44) probably indicative of low structural density, correspondingly this film is the one with the highest WVTR value. Being the ex-situ refractive index measurements probably affected by water adsorption, they cannot be univocally related with the film density. Thus, to better understand the film microstructure and identify a possible correlation with the abovementioned barrier properties an IR investigation was performed. The FT-IR spectra of the SiO2-like films deposited under the condition reported in Table 1 are reported in Fig. 2. The IR spectra show the absorption bands typical of Si–O–Si unit: asymmetric stretching at 1060–1070 cm− 1, bending at 810 cm− 1 and rocking at 460 cm− 1. Absorption characteristic of Si– OH (stretching O–H at 3500 cm− 1 and bending at 950 cm− 1) and organic groups (CHx stretching at 3000–2800 cm− 1, the Si–(CH3)x stretching at 1260–1270 cm− 1, and rocking at 800–840 cm− 1) are not detectable. The spectra present all the same absorption bands, confirming that the deposited films have the same chemical composition. The only differences can be found in the shape and in the position of the SiOSi stretching band: the maximum absorption wavenumbers and the intensities of the high-energy shoulder change depending on the precursor. The band is centered at around 1060 cm− 1 for films deposited from ATMS, TEOS, and DVTMDSO, while for HMDSO the band is centered at about 1070 cm− 1 and the high-energy shoulder is the most intense with respect to the other monomer. The position and the shape of the SiOSi stretching absorption depend on several factors: the thickness, the structural atomic arrangements, and a non perfect stoichiometry of the SiOSi network with homoatomic bonds (i.e. Si–Si, or O–O) [2]. Considering that the compared films have all the same thickness (about 100 nm) and the same stoichiometry (according to XPS), the Si–O stretching band wavenumber and the intensity of the shoulder seem to be linked to the network structure. In particular, literature shows that porosity and the presence of shorter chains in the network structure are associated with the high-energy shoulder, whose intensity is a rough estimate of the disorder [21], whereas the shift of the Si–O–Si asymmetric stretching absorption band could be due to a variation of the SiOSi bonding angle [22]. In thermal silicon dioxide films, the SiOSi absorption band can be fitted by two components: one taking into account the in phase motion (AS1) and the other for the out of phase asymmetric stretching (called AS2). According to the central force model, the maximum absorption wavenumber of the AS1 mode stretching band (v̄ASI) is related to the bonding angle of the Si–O–Si unit (θ) through the relation: θ νAS1 = P ν0 sin 2
P
ð1Þ
where the factor v̄0 = 1130.32 cm− 1 gives the AS1 wavenumber of thermal silica. Previous studies on silicon dioxide films have shown that a 100-nm-thick-thermally grown oxide has a bonding angle of 144° and the SiOSi stretching band falls at 1075 cm− 1. As it can be deduced from Eq. (1), the lower the bonding angle, the lower the absorption wavenumber is. Extensive studies in literature [22,23,24] associate low bonding angles (near 137° for SiOSi stretching band wavenumber of 1055 cm− 1) with high network density. Lucovsky et al. [22] attributed the increases in density for small bonding angles to better packing in the material, due to compressive stress. In the inset of Fig. 2 the fitting of the IR SiOSi stretch absorbance bands are shown for the films deposited from TEOS, and from HMDSO. The fitting has been done with two Gaussian functions to distinguish the AS1 and AS2 vibration modes. Apparently, the oxide deposited from TEOS seems to be characterized by smaller bonding angles due to the lower wavenumber of the AS1 band. This film shows the less intense high-energy shoulder.
A.M. Coclite et al. / Surface & Coatings Technology 204 (2010) 4012–4017
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Fig. 2. Normalized FT-IR spectra of SiO2-like films obtained from the different monomers as reported in Table 1. Inset shows best fitting of the SiOSi asymmetric stretching band for films deposited from TEOS and HMDSO.
Brunet-Brauneau et al. [23] proposed and verified a model for the evaporated SiO2-like films based on the assumption that the nonthermal silica films are composed of a dense matrix with pores, partially or fully filled with water. With this model they related the shift towards lower frequencies of the AS1 mode with an increasing matrix density, and the shift of the AS2 mode wavenumber towards higher values to an increasing of the pore fraction in the material. The pores could be partially or fully filled with water, explaining why a SiO2 film could have a lower density than fused silica (which is not porous) and at the same time a higher refractive index. According to this model, the absorption position of the AS1 and AS2 modes allow to distinguish between matrix density and porosity, while the refractive indices measured by ellipsometry are due to both the density and the fraction of voids, hence they are related with the average material density. In Fig. 3 the AS1 and AS2 absorption wavenumbers are reported for the films obtained from each precursor. The samples deposited from TEOS and DVTMDSO display the lowest AS1 mode absorption wavenumber, followed by the films deposited from ATMS and HMDSO. Assuming that the abovementioned model might be applied also to the present plasma deposited SiO2-like films, a density decrease might be hypothesized when passing from TEOS to HMDSO. For the AS2 mode absorption position, Brunet-Brauneau et al.[23] evidenced that this mode is much less sensitive than the AS1. In the present case, indeed, all samples showed similar absorption wavenumbers (in the range of 1174–1175 cm− 1) except for the film deposited from ATMS (1181 cm− 1). Following the model, the latter
Fig. 3. Trend of the AS1 (■) and AS2 (○) absorption wavenumbers for the films deposited from each precursor as a function of the WVTR.
film seems to be characterized by the highest fraction of voids, and this would reasonably explain the lower refractive index measured for the oxide from ATMS (1.44). It is worth noticing that the refractive indices measurements repeated after one year aging in air revealed almost no change, except for the film deposited from ATMS whose n value passed from 1.44 to 1.46. This latter increase in n would confirm that the sample deposited from ATMS was more porous and probably after one year the voids were fully filled with water. The adopted distinction between porosity and matrix density allows also to explain the trend of the barrier performances, indeed the WVTR values increase monotonically with the AS1 and AS2 absorption wavenumbers (Fig. 3). The best barrier films are the ones which seem to be characterized by both dense silica matrix and a small fraction of voids: with the AS1 and AS2 absorptions positioned towards low wavenumber (i.e. the ones deposited from TEOS and DVTMDSO). The worst barrier performances are obtained for the coatings with low matrix density and/or high fraction of voids (high AS1 and AS2 wavenumbers). The agreement between the barrier properties and the AS2 wavenumber (i.e. with the fraction of voids) is particularly significative confirming that non-porous oxide is needed for barrier applications. Furthermore, Lefevre's model [9] associated small SiOSi bonding angles with disordered networks characterized by three–four-membered rings. According to this model the present barrier performances could be explained considering that disordered network with small SiOSi bonding angles (i.e. low AS1 absorption) are more flexible, while ordered structures seem to be characterized by a higher rigidity which can be responsible of the failure of the mechanical and the barrier properties, as reported by Kawaguchi et al. [25]. Another important factor which could determine differences in permeation rates is the film surface roughness: Erlat et al. [3] showed that the OTR decreases from 9 to 0.4 cm3 m− 2 atm− 1 day− 1 with decreasing film surface roughness from 14 to 1 nm. The coatings reported in this work are characterized by low RMS roughness (listed in Table 1) ranging from 0.13 to 0.76 nm on Si. In particular the SiO2like film deposited from TEOS displays the lowest surfaces roughness, while the one deposited from ATMS showed the highest RMS roughness value. In Fig. 4 the AFM images of the bare plastic substrate (Fig. 4a) and of the one coated with 100-nm-thick film deposited from ATMS (Fig. 4b) are reported. Comparing the RMS roughness value of the bare substrate (1.5± 0. 4 nm) and of the coating from ATMS (1.7 ± 0.7 nm), one can conclude that the polymer surface roughness remains unaffected by film deposition, since the difference is within roughness standard deviation. If this is valid for the roughest films it is quite reasonable to conclude that it holds even for smoother ones. As a
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Fig. 4. Atomic force micrographs (5 µm × 5 µm) of bare substrate (a) (z scale 14 nm) and with 100-nm-thick SiO2-like film deposited from ATMS (b) (z scale 17 nm).
Hence efforts to relate the latter differences with the microstructure were proposed through ellipsometric and FT-IR investigations. The lowest WVTR values were found for films deposited from TEOS and DVTMDSO, characterized by refractive indices near 1.46 and low SiOSi stretching band absorption wavenumber. Melinex surface roughness was almost unaltered by the film deposition, independently of the precursor, thus any influence of the morphology on the barrier properties was ruled out. Although the exact nature of the film structure requires further investigation to be fully understood, a possible correlation between film microstructure and barrier properties was drawn adopting a structural model verified for the evaporated silicon dioxide films, based on infrared spectroscopy. The model allowed to obtain information on matrix density and the fraction of voids by the decomposition of the SiOSi IR stretching band in the AS1 and AS2 modes. The AS1 wavenumbers were associated with the bonding angles (as from the central force model) and the matrix density, while the AS2 wavenumbers with the film porosity: low AS1 and AS2 wavenumbers were seen as indicative of high density and low fraction of voids, respectively. The films deposited from TEOS and DVTMDSO showed the lowest AS1 and AS2 absorption wavenumber thus they were associated with low SiOSi bonding angles, dense matrix and low fraction of voids. Correspondingly, they showed the best barrier performances. The worst barrier properties were detected for the films deposited from HMDSO and ATMS being these apparently characterized by low matrix density (high AS1 wavenumber) and high porosity (high AS2 wavenumber), respectively. In conclusion, the precursor structure seems to influence the density and porosity of the SiOSi bonding network which in turn affect the barrier performances. Thus, to have good single layer inorganic barrier coating both dense matrix and low film porosity seem to be required. The best WVTR value showed a decrease of one order of magnitude respect to the bare PET substrate permeation. Such barrier improvement is promising for medical devices- or food-packaging applications. Future studies will be aimed to a better understanding of the influences of the monomer on the film microstructure deposited under high fragmentation and oxidative regime. Acknowledgments
consequence, the very low values of RMS roughness of deposited films along with the small variations in the polymer roughness itself upon film deposition, seem to rule out any influences of the surface roughness on the barrier properties. Finally, results seem to indicate that the presence of Si four-fold bonded with O in the precursor (i.e. TEOS) gives a network microstructure which positively influences the barrier performance while opposite effect is obtained when the monomer is lacking in SiO bonds (i.e. ATMS). Based on the results discussed so far, it is very difficult to elaborate a rationale of the influence of the monomer structure on the barrier performances. In fact many factors could affect the structure of the deposited coatings, such as energy input per monomer molecule, the addition of reactive gases (O2 and Ar) and the bias voltage responsible of the plasma-surface interactions. 4. Conclusions Nearly stoichiometric silicon dioxide-like films were deposited from different siloxane and silane organosilicon precursors, highly diluted in oxygen and argon. The films were all characterized by the same chemical composition and stoichiometry, notwithstanding the different structure of the precursors. Despite this similarity, some differences were appreciated regarding the barrier performances.
The authors would like to thank Savino Cosmai for his technical assistance, Anett Berndt (Siemens AG) for WVTR measurements. The financial support of the European Community VI Framework Programme (IP Napolyde NMP2-CT-2005-5158461) and of MIUR (PRIN07) are also acknowledged. References [1] L. Martinu, D. Poitras, J. Vac. Sci. Technol. A 18 (2000) 2619. [2] P.G. Pai, S.S. Chao, Y. Takagi, G. Lucovsky, J. Vac. Sci. Technol. A 4 (1986) 689. [3] A.G. Erlat, R.J. Spontak, R.P. Cloarke, T.C. Robinson, P.D. Haaland, Y. Tropsha, N.G. Harvey, E.A. Vogler, J. Phys. Chem. B 103 (1999) 6047. [4] R. Lamendola, R. d'Agostino, Pure Appl. Chem. 70 (1998) 1203. [5] M. Creatore, F. Palumbo, R. d'Agostino, Pure Appl. Chem. 74 (2002) 407. [6] M. Creatore, F. Palumbo, R. d'Agostino, P. Fayet, Surf. Coat. Technol. 142–144 (2001) 163. [7] A.S. da Silva Sobrinho, G. Czeremuskin, M. Latrèche, M.R. Wertheimer, J. Vac. Sci. Technol. A 18 (2000) 2021. [8] A.P. Roberts, B.M. Henry, A.P. Sutton, C.R.M. Grovenor, G.A.D. Briggs, T. Miyamoto, M. Kano, Y. Tsukahara, M. Yanaka, J. Membr. Sci. 208 (2002) 75. [9] A. Lefevre, L.J. Lewis, L. Martinu, M.R. Wertheimer, Phys. Rev. B 64 (2001) 115429-1-9. [10] M. Creatore, J.-C. Cigal, G.M.W. Kroesen, M.C.M. van de Sanden, Thin Solid films 484 (2005) 104. [11] A. Milella, M. Creatore, M.A. Blauw, M.C.M. van de Sanden, Plasma Process. Polym. 4 (2007) 621. [12] a. D. Hegemann, H. Brunner, C. Oehr, Surf Coat Technol. 142–144 (2001) 849; b. D. Hegemann, H. Brunner, C. Oehr, Surf Coat Technol. 174–175 (2003) 253; c. D. Hegemann, M.M. Hossain, Plasma Proc. Polym. 2 (2005) 554.
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On the relationship between the structure and the barrier performance of plasma deposited silicon dioxide-like films Anna Maria Coclite a,⁎, Antonella Milella a, Riccardo d'Agostino a, Fabio Palumbo b a b
Dipartimento di Chimica, Università degli Studi di Bari, Italy Istituto di Metodologie Inorganiche e dei Plasmi – CNR, Bari, Italy
a r t i c l e
i n f o
Article history: Received 14 September 2009 Accepted in revised form 14 May 2010 Available online 16 June 2010 Keywords: Barrier coatings FT-IR Plasma deposition Silicon dioxide
a b s t r a c t Silicon dioxide-like barrier films were deposited by plasma enhanced chemical vapor deposition from different siloxane and silane precursors. The variation of the precursor was investigated as a route to obtain silicon dioxide-like films with different structures, densities and hence barrier performances. Although the films were characterized by the same elemental composition, some differences in film density and porosity were evidenced from optical properties measurements and from the shift of the SiOSi infrared absorption band position. These differences were correlated with film microstructure and in turn with barrier performances. The results confirmed that films with high density and low porosity performed better as single inorganic barrier layers for food-packaging. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Over the past few decades, interest in transparent, thin SiO2-like coatings have increasingly grown for several applications like optics [1], interlayer dielectrics [2], and barrier films for the food-packaging and medical devices industries [3]. As gas and vapor diffusion barriers, such coatings have desirable properties, such as transparency, recyclability, and microwave use and are superior in this regard to the thin metal coatings (generally aluminium-based). The use of thin SiO2-like barrier films on polymer substrates is an interesting and challenging research field because the matching of thin inorganic barrier film to plastic substrates creates a new material which is impermeable and at the same time flexible, lightweight and environmental compatible. Although extensively studied [3–8], the SiO2-like film deposition for barrier applications still offers some interesting aspects, concerning the correlation between film properties (e.g. density and porosity) and performances. Previous reports on plasma deposited barrier coatings from organosilicon precursors mainly correlated the film chemical composition with the barrier properties demonstrating that the permeability can be significantly improved minimizing the content of organic moieties and of silanol groups in the films. For example, the Oxygen Transmission Rate (OTR) through a PET substrate can be lowered from 2–5 cm3 m− 2 day− 1 atm− 1 to 0.2–0.5 cm3 m− 2 day− 1 atm− 1 [6,7] and the water vapor transmission rate (WVTR) to 0.5 g m− 2 day− 1 ⁎ Corresponding author. Dipartimento di Chimica, Università degli Studi di Bari, Via Orabona, 4, 70126 Bari, Italy. Tel.: +39 0805443433; fax: + 39 0805443405. E-mail address: [email protected] (A.M. Coclite). 0257-8972/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2010.05.024
[4,5] at 23 °C and 50% of relative humidity (RH). It was demonstrated that the presence of organic and silanol groups deteriorates the barrier properties both because the SiOSi network is interrupted and because they increase the solubility of the permeant molecules in the film favouring their transmission through the barrier film. Beside film chemistry, other factors such as film defects, film adhesion to polymeric substrate and substrate surface defects were demonstrated to influence the film permeability to gas and vapor. da Silva Sobrinho et al. [7] showed that the presence of unwanted but inevitable nano-, microscopic defects and pinholes in the coating, which may result from dust particles on the substrate surface, from geometric shadowing and stress during film growth at sites of high surface roughness, limits the minimum value achievable of OTR. The smaller the defect density, the lower the permeation rate is. Roberts et al. [8], by means of a model of the permeation process in evaporated silicon oxide gas barrier films, identified such defects as distinguishable in three types: macro-defects (cracks N 1 nm), nano-defects (pinholes b 1 nm) and amorphous oxide lattice defects (interstice at grain boundary b 0.3 nm). The barrier properties have never been directly correlated with the film microstructure in terms of density and porosity although extensive studies in literature show that silica coatings can be significantly different in the atomic structure depending on the particular deposition process used: both silicon and oxygen atoms can be found in different environments, allowing the presence of homoatomic bonds in disordered chains. Dense nearly stoichiometric silicon dioxide films can be achieved by Ion Beam Assisted Deposition (IBAD) [9], increasing the substrate temperature [10] or biasing the substrate during deposition [11]. Lefevre et al. [9] studied and modeled the microstructural densification process of silicon dioxide
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thin films induced by medium-energy particles bombardment. They considered the description of the network topology in terms of rings of nearest-neighbour bonds, where an n-membered ring is defined as the shortest closed path of n Si–O bonds originating on a given silicon atom. According to this description, in amorphous silica, the structure is statistically dominated by five, six and seven-fold rings. In deposited silicon dioxide-like films, the presence of small rings (three- and fourmembered) is associated with stress and defects in the network, while large rings (e.g. nine-membered) corresponds to film porosity. Ion bombardment induces the structural densification of the SiO2-like films because small rings rearrange to produce larger and less strained ones and too large rings turn into five- six- and seven-membered ones, reducing the free volume and causing the density to increase. They concluded that the densification process can be identified as a relaxation mechanism resulting in silicon dioxide-like films with low stress and low defect concentration. The microstructure and in turn the properties of the SiO2-like films can be tuned by changing the plasma parameters [12] (e.g. O2 flow rate, power, and bias) and/or by choosing a suitable monomer [13,14]. Hegemann et al. [12] showed for HMDSO/O2 plasmas that the power input per gas flow (W/F) controls the chemical composition of the film, determining the discharge fragmentation, while the mechanical properties are mostly influenced by the potential drop across the plasma sheath. High W/F values, such as the ones required to deposit SiO2-like films, are indicative of a monomer-deficient regime in which the precursor structure becomes more and more negligible due to strong fragmentation. Different monomers were previously used to deposit SiO2-like films resulting in similar chemical composition but different deposition rate [13,14]. This is believed to affect the microstructure and in turn the performance of the SiO2-like films. At the authors knowledge the literature does not report the correlation between the structure of the monomer and the barrier performances. The aim of the present paper was to understand whether the monomer structure could affect the barrier performance of SiO2-like coatings deposited in highly oxidative regimes and high fragmentation conditions. Hence four different monomers were selected as follows: (i) tetraethoxysilane (TEOS) and hexamethyldisiloxane (HMDSO) as usual source of SiO2-like films [3–6,15] and (ii) divinyltetramethyldisiloxane (DVTMDSO) and allyltrimethylsilane (ATMS) to compare film structure derived from siloxane and silane. Furthermore, DVTMDSO and ATMS with vinyl and allyl functional groups are highly attractive precursors for inorganic/organic multilayer barrier film deposition [16]. 2. Materials and methods Silicon dioxide-like films were deposited in a capacitively coupled asymmetric parallel plate reactor from different precursors [17]. The upper electrode was connected to a 13.56 MHz power supply (RFPP) through an automatic LC matching network. The reactor was evacuated by means of a turbomolecular/rotary pumping system and the pressure was monitored with a MKS capacitive gauge. Discharge conditions were chosen to obtain C-free coatings with stoichiometry close to SiO2, thus high RF power (i.e. 250 W) and high O2 to monomer ratio (i.e. N20) were used. The deposition parameters are summarized in Table 1. Because of the different chemical structures of the four precursors used, the O2/monomer for each deposition was chosen largely higher than the stoichiometric value obtained from the complete oxidation reaction. Table 2 shows the chemical structures of the precursors used for the depositions. Film chemical and optical characterization was performed using polished Si as substrate, while barrier performance tests were accomplished depositing the coatings on 125-μm-thick poly(ethylene terephthalate) PET (Melinex® ST 504 manufactured by DuPont Teijin). The morphological analysis was performed on both kinds of substrates for all the films deposited.
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Table 1 Experimental parameters utilized for film deposition, the corresponding deposition rates, refractive indices (at 633 nm) and surface RMS roughness (on Si). The RF power was fixed at 250 W. Monomer
O2/mon
Ar flow rate (sccm)
Pressure (mTorr)
Dep rate (nm/min)
n
RMS (nm)
TEOS DVTMDSO HMDSO ATMS
22 14 20 21
44 200 200 150
100 150 200 100
48 ± 4 119 ± 11 69 ± 3 7±1
1.457 1.458 1.461 1.444
0.13 ± 0.06 0.49 ± 0.04 0.60 ± 0.16 0.76 ± 0.09
The film thickness was measured by means of a profilometer (TENCOR α-step 500). All the films were grown up to 100 nm in thickness. Film chemical composition and structure were investigated by FTIR spectroscopy (BRUKER, Equinox 55). Spectra were recorded from 400 to 4000 cm− 1 in absorbance mode at 4 cm− 1 resolution. In order to minimize effects of water vapor and carbon dioxide, the spectrometer was purged with nitrogen for 15 min between each measurement. Baseline correction was applied to the spectra and they were then normalized to the film thickness. The SiOSi stretching band was fitted using two Gaussian components using the “Fit multi-peaks” procedure of OriginLab software. Atomic percentage composition of the films was determined by Xray photoelectron spectroscopy (XPS) (Thermo VG Scientific). Low and high-resolution spectra were acquired using a monochromatic Al Kα radiation (1486.6 eV). The take off angle was 53°. In order to remove carbon contamination from the sample surface, a 10 s sputter cleaning with 1 KeV, 500 nA Ar+ ions was carried out before the analysis. During the XPS analysis, the sample charge was compensated by a 1 eV electron beam at high neutralization current by means of a flood gun. Surface film morphology was investigated by Atomic Force Microscopy (AFM, AutoProbe CP, Thermomicroscope). Images were acquired in non-contact mode using conical gold-coated silicon tips. Root-mean-square (RMS) roughness was measured on 5 μm × 5 μm surface areas averaging 15 measurements for each sample. Ex-situ refractive index measurements were carried out through variable angle spectroscopy ellipsometry (VASE, JA Woollam M2000). The measurements were done at three different angles (65°, 70° and 75°) in the wavelength range of 200 nm to 1000 nm. The applied optical model consisted of four components: the silicon substrate, the native SiO2 layer, the bulk film, and the surface roughness layer. The bulk components were modeled by the Cauchy function while the top-layer using the Bruggeman Effective Medium Approximation (BEMA). The BEMA layer consisted of 50% bulk film and 50% voids. The model also incorporated possible thickness inhomogeneity within the sampled area [18]. Table 2 Chemical structures of the monomers used for the depositions.
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Water vapor transmission rate (WVTR) of the barrier films on PET was measured on a 100 cm2 active area, at 25 °C and 98% of relative humidity (RH), using a MOCON Permatran-W. The WVTR measured through the bare Melinex was 4.46 g m− 2 day− 1. 3. Results and discussion Inorganic silicon dioxide films were deposited in high fragmentation regime (i.e. high oxygen flow rate and high RF power). Varying the precursor structure the oxygen flow rate was changed accordingly, in order to completely oxidize the organic moieties. The ratios were deduced from preliminary experiments aimed at obtaining Cfree polymers at low pressure. Table 1 shows the deposition conditions exploited for this study. XPS elemental composition analysis showed atomic percent of carbon less than 1% and O/Si ratio near 2 for all deposited films. Film stoichiometry was further investigated by curve fitting of the Si2p signal. The curve fitting was accomplished by first subtracting the spin-orbit splitting contribution [19] and then the signals were regressed in the SiyO4 − y components, with y from 0 to 4, depending on the number of Si–O or Si–Si bonds per silicon atom [20]. The least squares regression to the signals was found by using only two components of FWHM around 2 eV [19]. For all the films, the main component was positioned at 103.4 eV, and thus associated with SiO4 groups. The second contribution, centered at 101.9 eV was attributed to Si2O3 groups and represented the 0.5–1.4% of the total Si 2p area. Similar fitting profiles were found by Alfonsetti et al. [20] for films with SiO1.97 stoichiometry. No significative differences were found between the five samples confirming that they are very close in stoichiometry. Fig. 1 shows the WVTR measured for the deposited films. Different performances were detected notwithstanding the similarity in chemical composition evidenced. Indeed only the films deposited from TEOS and DVTMDSO showed good barrier properties, reducing of one order of magnitude the transmission rate through the bare substrate. In order to explain these differences, the refractive indices were measured. Due to the absence of carbon, the measured refractive index basically results from the density and the fraction of voids in the films: high n values can be regarded as indicative of high film density. In the present case, being the measurements carried out ex-situ, the n values can be affected by the water vapor adsorbed in the pores: generally this phenomenon results in an increased refractive index with respect to the in-situ measurements due to the increase of the pore refractive index from 1 of void to 1.332 of water [18]. Analyzing the indices reported in Table 1, it should be noted that the silicon dioxide films deposited from TEOS, DVTMDSO and HMDSO have similar refractive indices, very close to the one characteristic of silica glass (1.46). Despite this similarity, the WVTRs for these samples are different. Significative differences in the refractive indices can be
Fig. 1. Histogram of the WVTR measured at 98% relative humidity and 25 °C for the bare PET substrate and through the film obtained from the different precursors.
appreciated for the film deposited from the ATMS. It shows a low n (1.44) probably indicative of low structural density, correspondingly this film is the one with the highest WVTR value. Being the ex-situ refractive index measurements probably affected by water adsorption, they cannot be univocally related with the film density. Thus, to better understand the film microstructure and identify a possible correlation with the abovementioned barrier properties an IR investigation was performed. The FT-IR spectra of the SiO2-like films deposited under the condition reported in Table 1 are reported in Fig. 2. The IR spectra show the absorption bands typical of Si–O–Si unit: asymmetric stretching at 1060–1070 cm− 1, bending at 810 cm− 1 and rocking at 460 cm− 1. Absorption characteristic of Si– OH (stretching O–H at 3500 cm− 1 and bending at 950 cm− 1) and organic groups (CHx stretching at 3000–2800 cm− 1, the Si–(CH3)x stretching at 1260–1270 cm− 1, and rocking at 800–840 cm− 1) are not detectable. The spectra present all the same absorption bands, confirming that the deposited films have the same chemical composition. The only differences can be found in the shape and in the position of the SiOSi stretching band: the maximum absorption wavenumbers and the intensities of the high-energy shoulder change depending on the precursor. The band is centered at around 1060 cm− 1 for films deposited from ATMS, TEOS, and DVTMDSO, while for HMDSO the band is centered at about 1070 cm− 1 and the high-energy shoulder is the most intense with respect to the other monomer. The position and the shape of the SiOSi stretching absorption depend on several factors: the thickness, the structural atomic arrangements, and a non perfect stoichiometry of the SiOSi network with homoatomic bonds (i.e. Si–Si, or O–O) [2]. Considering that the compared films have all the same thickness (about 100 nm) and the same stoichiometry (according to XPS), the Si–O stretching band wavenumber and the intensity of the shoulder seem to be linked to the network structure. In particular, literature shows that porosity and the presence of shorter chains in the network structure are associated with the high-energy shoulder, whose intensity is a rough estimate of the disorder [21], whereas the shift of the Si–O–Si asymmetric stretching absorption band could be due to a variation of the SiOSi bonding angle [22]. In thermal silicon dioxide films, the SiOSi absorption band can be fitted by two components: one taking into account the in phase motion (AS1) and the other for the out of phase asymmetric stretching (called AS2). According to the central force model, the maximum absorption wavenumber of the AS1 mode stretching band (v̄ASI) is related to the bonding angle of the Si–O–Si unit (θ) through the relation: θ νAS1 = P ν0 sin 2
P
ð1Þ
where the factor v̄0 = 1130.32 cm− 1 gives the AS1 wavenumber of thermal silica. Previous studies on silicon dioxide films have shown that a 100-nm-thick-thermally grown oxide has a bonding angle of 144° and the SiOSi stretching band falls at 1075 cm− 1. As it can be deduced from Eq. (1), the lower the bonding angle, the lower the absorption wavenumber is. Extensive studies in literature [22,23,24] associate low bonding angles (near 137° for SiOSi stretching band wavenumber of 1055 cm− 1) with high network density. Lucovsky et al. [22] attributed the increases in density for small bonding angles to better packing in the material, due to compressive stress. In the inset of Fig. 2 the fitting of the IR SiOSi stretch absorbance bands are shown for the films deposited from TEOS, and from HMDSO. The fitting has been done with two Gaussian functions to distinguish the AS1 and AS2 vibration modes. Apparently, the oxide deposited from TEOS seems to be characterized by smaller bonding angles due to the lower wavenumber of the AS1 band. This film shows the less intense high-energy shoulder.
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Fig. 2. Normalized FT-IR spectra of SiO2-like films obtained from the different monomers as reported in Table 1. Inset shows best fitting of the SiOSi asymmetric stretching band for films deposited from TEOS and HMDSO.
Brunet-Brauneau et al. [23] proposed and verified a model for the evaporated SiO2-like films based on the assumption that the nonthermal silica films are composed of a dense matrix with pores, partially or fully filled with water. With this model they related the shift towards lower frequencies of the AS1 mode with an increasing matrix density, and the shift of the AS2 mode wavenumber towards higher values to an increasing of the pore fraction in the material. The pores could be partially or fully filled with water, explaining why a SiO2 film could have a lower density than fused silica (which is not porous) and at the same time a higher refractive index. According to this model, the absorption position of the AS1 and AS2 modes allow to distinguish between matrix density and porosity, while the refractive indices measured by ellipsometry are due to both the density and the fraction of voids, hence they are related with the average material density. In Fig. 3 the AS1 and AS2 absorption wavenumbers are reported for the films obtained from each precursor. The samples deposited from TEOS and DVTMDSO display the lowest AS1 mode absorption wavenumber, followed by the films deposited from ATMS and HMDSO. Assuming that the abovementioned model might be applied also to the present plasma deposited SiO2-like films, a density decrease might be hypothesized when passing from TEOS to HMDSO. For the AS2 mode absorption position, Brunet-Brauneau et al.[23] evidenced that this mode is much less sensitive than the AS1. In the present case, indeed, all samples showed similar absorption wavenumbers (in the range of 1174–1175 cm− 1) except for the film deposited from ATMS (1181 cm− 1). Following the model, the latter
Fig. 3. Trend of the AS1 (■) and AS2 (○) absorption wavenumbers for the films deposited from each precursor as a function of the WVTR.
film seems to be characterized by the highest fraction of voids, and this would reasonably explain the lower refractive index measured for the oxide from ATMS (1.44). It is worth noticing that the refractive indices measurements repeated after one year aging in air revealed almost no change, except for the film deposited from ATMS whose n value passed from 1.44 to 1.46. This latter increase in n would confirm that the sample deposited from ATMS was more porous and probably after one year the voids were fully filled with water. The adopted distinction between porosity and matrix density allows also to explain the trend of the barrier performances, indeed the WVTR values increase monotonically with the AS1 and AS2 absorption wavenumbers (Fig. 3). The best barrier films are the ones which seem to be characterized by both dense silica matrix and a small fraction of voids: with the AS1 and AS2 absorptions positioned towards low wavenumber (i.e. the ones deposited from TEOS and DVTMDSO). The worst barrier performances are obtained for the coatings with low matrix density and/or high fraction of voids (high AS1 and AS2 wavenumbers). The agreement between the barrier properties and the AS2 wavenumber (i.e. with the fraction of voids) is particularly significative confirming that non-porous oxide is needed for barrier applications. Furthermore, Lefevre's model [9] associated small SiOSi bonding angles with disordered networks characterized by three–four-membered rings. According to this model the present barrier performances could be explained considering that disordered network with small SiOSi bonding angles (i.e. low AS1 absorption) are more flexible, while ordered structures seem to be characterized by a higher rigidity which can be responsible of the failure of the mechanical and the barrier properties, as reported by Kawaguchi et al. [25]. Another important factor which could determine differences in permeation rates is the film surface roughness: Erlat et al. [3] showed that the OTR decreases from 9 to 0.4 cm3 m− 2 atm− 1 day− 1 with decreasing film surface roughness from 14 to 1 nm. The coatings reported in this work are characterized by low RMS roughness (listed in Table 1) ranging from 0.13 to 0.76 nm on Si. In particular the SiO2like film deposited from TEOS displays the lowest surfaces roughness, while the one deposited from ATMS showed the highest RMS roughness value. In Fig. 4 the AFM images of the bare plastic substrate (Fig. 4a) and of the one coated with 100-nm-thick film deposited from ATMS (Fig. 4b) are reported. Comparing the RMS roughness value of the bare substrate (1.5± 0. 4 nm) and of the coating from ATMS (1.7 ± 0.7 nm), one can conclude that the polymer surface roughness remains unaffected by film deposition, since the difference is within roughness standard deviation. If this is valid for the roughest films it is quite reasonable to conclude that it holds even for smoother ones. As a
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Fig. 4. Atomic force micrographs (5 µm × 5 µm) of bare substrate (a) (z scale 14 nm) and with 100-nm-thick SiO2-like film deposited from ATMS (b) (z scale 17 nm).
Hence efforts to relate the latter differences with the microstructure were proposed through ellipsometric and FT-IR investigations. The lowest WVTR values were found for films deposited from TEOS and DVTMDSO, characterized by refractive indices near 1.46 and low SiOSi stretching band absorption wavenumber. Melinex surface roughness was almost unaltered by the film deposition, independently of the precursor, thus any influence of the morphology on the barrier properties was ruled out. Although the exact nature of the film structure requires further investigation to be fully understood, a possible correlation between film microstructure and barrier properties was drawn adopting a structural model verified for the evaporated silicon dioxide films, based on infrared spectroscopy. The model allowed to obtain information on matrix density and the fraction of voids by the decomposition of the SiOSi IR stretching band in the AS1 and AS2 modes. The AS1 wavenumbers were associated with the bonding angles (as from the central force model) and the matrix density, while the AS2 wavenumbers with the film porosity: low AS1 and AS2 wavenumbers were seen as indicative of high density and low fraction of voids, respectively. The films deposited from TEOS and DVTMDSO showed the lowest AS1 and AS2 absorption wavenumber thus they were associated with low SiOSi bonding angles, dense matrix and low fraction of voids. Correspondingly, they showed the best barrier performances. The worst barrier properties were detected for the films deposited from HMDSO and ATMS being these apparently characterized by low matrix density (high AS1 wavenumber) and high porosity (high AS2 wavenumber), respectively. In conclusion, the precursor structure seems to influence the density and porosity of the SiOSi bonding network which in turn affect the barrier performances. Thus, to have good single layer inorganic barrier coating both dense matrix and low film porosity seem to be required. The best WVTR value showed a decrease of one order of magnitude respect to the bare PET substrate permeation. Such barrier improvement is promising for medical devices- or food-packaging applications. Future studies will be aimed to a better understanding of the influences of the monomer on the film microstructure deposited under high fragmentation and oxidative regime. Acknowledgments
consequence, the very low values of RMS roughness of deposited films along with the small variations in the polymer roughness itself upon film deposition, seem to rule out any influences of the surface roughness on the barrier properties. Finally, results seem to indicate that the presence of Si four-fold bonded with O in the precursor (i.e. TEOS) gives a network microstructure which positively influences the barrier performance while opposite effect is obtained when the monomer is lacking in SiO bonds (i.e. ATMS). Based on the results discussed so far, it is very difficult to elaborate a rationale of the influence of the monomer structure on the barrier performances. In fact many factors could affect the structure of the deposited coatings, such as energy input per monomer molecule, the addition of reactive gases (O2 and Ar) and the bias voltage responsible of the plasma-surface interactions. 4. Conclusions Nearly stoichiometric silicon dioxide-like films were deposited from different siloxane and silane organosilicon precursors, highly diluted in oxygen and argon. The films were all characterized by the same chemical composition and stoichiometry, notwithstanding the different structure of the precursors. Despite this similarity, some differences were appreciated regarding the barrier performances.
The authors would like to thank Savino Cosmai for his technical assistance, Anett Berndt (Siemens AG) for WVTR measurements. The financial support of the European Community VI Framework Programme (IP Napolyde NMP2-CT-2005-5158461) and of MIUR (PRIN07) are also acknowledged. References [1] L. Martinu, D. Poitras, J. Vac. Sci. Technol. A 18 (2000) 2619. [2] P.G. Pai, S.S. Chao, Y. Takagi, G. Lucovsky, J. Vac. Sci. Technol. A 4 (1986) 689. [3] A.G. Erlat, R.J. Spontak, R.P. Cloarke, T.C. Robinson, P.D. Haaland, Y. Tropsha, N.G. Harvey, E.A. Vogler, J. Phys. Chem. B 103 (1999) 6047. [4] R. Lamendola, R. d'Agostino, Pure Appl. Chem. 70 (1998) 1203. [5] M. Creatore, F. Palumbo, R. d'Agostino, Pure Appl. Chem. 74 (2002) 407. [6] M. Creatore, F. Palumbo, R. d'Agostino, P. Fayet, Surf. Coat. Technol. 142–144 (2001) 163. [7] A.S. da Silva Sobrinho, G. Czeremuskin, M. Latrèche, M.R. Wertheimer, J. Vac. Sci. Technol. A 18 (2000) 2021. [8] A.P. Roberts, B.M. Henry, A.P. Sutton, C.R.M. Grovenor, G.A.D. Briggs, T. Miyamoto, M. Kano, Y. Tsukahara, M. Yanaka, J. Membr. Sci. 208 (2002) 75. [9] A. Lefevre, L.J. Lewis, L. Martinu, M.R. Wertheimer, Phys. Rev. B 64 (2001) 115429-1-9. [10] M. Creatore, J.-C. Cigal, G.M.W. Kroesen, M.C.M. van de Sanden, Thin Solid films 484 (2005) 104. [11] A. Milella, M. Creatore, M.A. Blauw, M.C.M. van de Sanden, Plasma Process. Polym. 4 (2007) 621. [12] a. D. Hegemann, H. Brunner, C. Oehr, Surf Coat Technol. 142–144 (2001) 849; b. D. Hegemann, H. Brunner, C. Oehr, Surf Coat Technol. 174–175 (2003) 253; c. D. Hegemann, M.M. Hossain, Plasma Proc. Polym. 2 (2005) 554.
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