Optimization of N-doped TiO2 multifunctional thin layers by low frequency PECVD process

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Optimization of N-doped TiO2 multifunctional thin layers by low frequency PECVD process Loraine Youssef a,b , Arnaud Joël Kinfack Leoga a , Stéphanie Roualdes a,∗ , Joëlle Bassil b , Mirvat Zakhour b , Vincent Rouessac a , André Ayral a , Michel Nakhl b a Institut Européen des Membranes (IEM), UMR5635 (ENSCM, UM, CNRS), University of Montpellier, Place Eugène Bataillon, CC047, 34095 Montpellier Cedex 5, France b Laboratoire Chimie-Physique des Matériaux (LCPM)/Plateforme de Recherches en Nanosciences (PR2N), Lebanese University, B.P. 90656 Fanar, Jdeidet el Metn, Beirut, Lebanon

a r t i c l e

i n f o

Article history: Received 30 January 2017 Received in revised form 3 May 2017 Accepted 9 May 2017 Available online xxx Keywords: TiO2 Low-frequency plasma N-doping Photocatalysis

a b s t r a c t A one-step low-frequency Plasma Enhanced Chemical Vapor Deposition (PECVD) process, operating at temperature as low as 350 ◦ C, has been implemented to prepare single-oriented pure and N-doped anatase films. The layers have been synthesized using titanium isopropoxide as a precursor, and NH3 as a doping agent. Optimized PECVD conditions have enabled to obtain homogeneous micro-columnar porous thin films with thicknesses close to 500 nm. Depth profiling XPS analyses have proved the nitrogen incorporation into TiO2 lattice after ammonia introduction in the deposition chamber. As another proof of N-doping, Raman and XRD peaks shifting have been observed. Such thin films have been demonstrated as efficient photocatalytic materials which activity region can be tailored from UV to visible region by adjusting the proportion of doping agent in the plasma phase. Due to their microstructural and photocatalytic properties, the prepared thin layers should have an interest as anode materials in solar water splitting cells. © 2017 Elsevier Ltd. All rights reserved.

1. Introduction Direct coupling of separation and photocatalytic degradation using titania-based membranes is an attractive approach intensively investigated in air or water treatment devices [1–4]. Such approach involving multifunctional titania layers can also be envisaged for other applications and in particular for H2 production from solar light by water splitting, which is precisely the ultimate application aimed in this study. For H2 production by solar water splitting, specific original integrated devices have to be conceived combining high energyconversion efficiency, direct hydrogen separation and possible up-scaling with low manufacturing cost. Photo-electro-catalytic cells were originally proposed by Fujishima and Honda in 1972 [5]. Since 2007, some multilayered systems have been described in the literature [6–10] but neither of them is really integrated (that is to say including ␮-architectured uni-directionally aligned layers) or involves plasma-deposited thin films. The main advantage for such a configuration is the facility of charge transfer to

∗ corresponding author. E-mail address: [email protected] (S. Roualdes).

the material surface due to the short particle freeway (thin films). Another advantage for the sandwich-like architecture is the continuous charge transfer among the layers, especially that the different compartments are stuck together using an advanced pressing process. Now plasma processes are very competitive to insure stability, integrity and compactness of material devices. Among plasma processes, Plasma-Enhanced Chemical Vapor Deposition (PECVD) is a very promising low-temperature one-step method for preparing and stacking layers on any kind of substrate. This versatile method enables optimization of structural and functional properties of layers. As reported in the literature, TiO2 thin films have been recently prepared by radio frequency (R.F.) PECVD with titanium tetraisopropoxide (Ti(OC3 H7 )4 , TTIP) as a precursor. Anatase thin film could be obtained either directly through the deposition step by heating the silicon substrate at 450 ◦ C [11] or by applying a bias voltage on it [12,13], or by performing post-annealing at 400 ◦ C after the deposition step [14]. In our group [15], a R.F. PECVD process operating at 150 ◦ C has been implemented to prepare micro-columnar porous TiO2 anatase thin films, performing postannealing at 300 ◦ C for 5 h. Optimized PECVD conditions have enabled to obtain micrometer-thick homogeneous films. The size

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of nano-crystals in prepared anatase thin films has been estimated to be 20 nm by applying Scherrer equation [16]. Besides, the bandgap energy of synthesized anatase thin films on quartz was found as 3.30 eV. At this stage, PECVD TiO2 films can still be optimized in terms of preparation conditions (the simplest and less energetic process being the best), photocatalytic activity and separation properties. The optimization of the properties of such films so that they could be competitive photo-anode in water splitting cells is precisely the scope of this study. Well-known for its high photoactivity in the UV range, TiO2 does not absorb visible light which makes it inefficient for water splitting by solar irradiation. In order to improve the photoactivity of the TiO2 films under solar light, the intrinsic band gap of TiO2 has to be reduced. For that purpose, several strategies have been investigated including the incorporation of either metallic (e.g. Fe, Ni, Ag, Cu) or non-metallic (e.g. C, F, N, P) elements into the TiO2 host material, or the creation of oxygen defects in the TiO2 structure. Recent articles have been focused on N-doped TiO2 thin films since these materials reveal behaviors between metallic TiN and resistive TiO2 compounds. As early as 2001, Asahi et al. [17] have prepared N-TiO2 by reactive sputtering while Ihara et al. [18] have worked on the hydrolysis of TiCl4 and TiCl3 using NH4 OH. A few years later, physical methods such as ion implantation have been introduced by the groups of Ghicov et al. [19] and Yang et al. [20] while ball milling has been proposed in the works of Kang et al. [21]. Other chemical methods have been performed like hydrolyzation of titanium alkoxides by sol-gel methods [22–24], direct hydrolysis of organic or inorganic salts under hydrothermal or solvothermal conditions [25] and oxidation of titanium nitride [26,27]. Other approaches use a plasma process. In the works of Maeda et al. [28], N-doped TiO2 films were prepared by radiofrequency PECVD using titanium isopropoxide with Ar carrier gas and NH3 mixture in the range from 25 to 150 sccm. Nitrogen content increases slightly with increasing NH3 flow rate. It was concluded that both nitrogen atoms substituted for oxygen atoms and crystalline structure contribute to the visible light photocatalytic activity. Reed et al. have prepared mesoporous TiO2 thin films by a surfactant template sol-gel method, followed by a nitrogen plasma treatment as an N-doping approach; the band gap of TiO2 films was reduced from 3.5 eV to a minimum of 3.0 eV after the plasma treatment [29]. As another N-doping competitive approach using a plasma process, it has been recently proved that PECVD makes possible the deposition of N-doped TiO2 materials in one step (adding N2 or NH3 as N-dopant in the plasma chamber during the TiO2 deposition step) [30]. This work aims at proving the feasibility of getting pure and N-doped anatase homogeneous phase by Low-Frequency PECVD process in one step, that is to say without post-annealing step or post-incorporation of doping agent. The novelty of this study, compared to the previous results obtained in the literature, resides in the fact that titania synthesis by Low Frequency PECVD process has never been mentioned yet. In addition, it is true that PECVD doping effect has been proved by structural characterizations in previous works but considering the functional applications, the results came less than expected. The good photocatalytic ability of the prepared films, in particular in the visible region, has also to be proved so that such materials could be envisaged as photocatalytic anode in plasma multilayered devices for H2 production by solar light. Thin films of titanium dioxide were deposited on single crystalline silicon (100) wafers and on borosilicate strips or discs in a low-frequency PECVD reactor using TTIP as a TiO2 precursor (mixed with Ar and O2 ). Two variable synthesis conditions (substrate temperature, plasma power) were modified in order to obtain in-situ anatase without post-annealing but with higher deposition rate, Ti O abundance and photocatalytic activity. Addition of NH3 in the plasma gaseous phase (during the deposition step) was envisaged as an in-situ N-doping approach. Different

ammonia partial pressures were applied to be correlated with the N content in materials and enhancement of the TiO2 photocatalytic efficiency in the visible region. Morphology, microstructure and structure of TiO2 films were studied by Scanning Electron Microscopy (SEM), Fourier-Transformed Infra-Red spectroscopy (FTIR), Energy Dispersive X-ray spectroscopy (EDX), X-Ray Photoelectron Spectroscopy (XPS), X-Ray Diffraction (XRD) and Raman Spectroscopy. Films topographies were studied by Atomic Force Microscopy (AFM). Films functional properties in relation with photocatalysis application were investigated by UV–vis absorbance analysis and Pilkington test.

2. Experimental details 2.1. TiO2 thin films synthesis procedure Titanium dioxide films were deposited on single crystalline silicon (100) substrates (Monsanto Electronic Materials Company), ® borosilicate glass strips (Corning 0211 from Abrisa Technologies) and borosilicate glass discs (20 mm diameter discs from SAS Rossignol) using a capacitively-coupled low-frequency PECVD reactor depicted in Fig. 1. In such a reactor configuration, two symmetrical circular electrodes (here spaced by 30 cm), connected to the low-frequency generator (Alsatherm F 68500, 40 kHz), were placed outside and around a quartz tube playing the role of deposition chamber. Before deposition, the substrates were cleaned using acetone and ethanol. Then, they were introduced into the deposition chamber and maintained under vacuum (at a limit pressure of 1.5 Pa) by primary pumping overnight. The liquid precursor titanium tetraisopropoxide (TTIP), from Sigma-Aldrich, was put in a container immersed in an oil bath heated at 80 ◦ C. Inert argon was bubbling in the TTIP liquid as a carrier gas. The carrying line transporting both argon and TTIP from the container to the deposition chamber was heated at 100 ◦ C in order to avoid any precursor condensation in the line. Oxygen was introduced as the oxidant gas. The fluxes ® of both argon and oxygen were controlled using Bronkhorst HighTech flow-meters. Ammonia gas (NH3 ) was used as nitrogen source in the doping process; its flux was controlled by a millimetric valve. A cold liquid nitrogen trap, separating the deposition chamber from the pump, retained the non-reacted species during the deposition process. The substrate temperature during the deposition step was controlled using a tubular heating device (Eurotherm) surrounding the deposition chamber. The deposition duration of 20 min was kept constant for all the samples. The films were synthesized under 22.5 TTIP + Ar partial pressure (Ar flux equal to 10 sccm) and 17 Pa O2 partial pressure (flux equal to 35.5 sccm). Two different discharge powers, 64 W and 100 W, and four different values of the substrate temperature (Ts ), 150 ◦ C, 250 ◦ C, 300 ◦ C and 350 ◦ C, were investigated in order to study their effect on the films deposition rate and microstructure. After the deposition process (once the reactor has been cooled down (0.5 ◦ C/min)), each PECVD deposit was divided into two pieces: one was removed from the chamber and kept as the as-deposited TiO2 and the other one was left into the deposition chamber for post-annealing treatment. For such a treatment, the sample was heated up to 400 ◦ C (with an average heating rate of 5 ◦ C/min) and maintained at this temperature for 1 h using the same heating device as those used for the deposition step. In the specific case of the highest substrate temperature (350 ◦ C), the addition of NH3 in the deposition chamber at the pressures of 5, 10 or 15 Pa (PNH3 ) was investigated for doping purpose. Depending on the NH3 partial pressure, the total pressure in the deposition chamber during the process was 39.5 (without NH3 ), 44.5, 49.5 or 54.5 Pa.

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Fig. 1. Schematic representation of the capacitively-coupled Low Frequency PECVD reactor.

2.2. Thin film characterizations The morphology and thickness of the TiO2 films deposited on silicon were investigated by SEM (Scanning Electron Microscopy) using a S4800 Hitachi (10% instrumental error on thickness measurement). The surface topography of films deposited on borosilicate glass strips was observed by Atomic Force Microscopy (AFM) using Agilent Technologies 5500 apparatus. The surface ® average roughness was calculated using Picoview SPM software provided with the measuring device. Film chemical structure was analyzed by FTIR (Fourier Transform Infrared) using a Nicolet Nexus 710 and a JASCO FT/IR-6300 (in transmission mode). Element analysis was determined by EDX (X-Ray Dispersive Energy) using an EVO-HD S4500 apparatus bearing a total instrumental error of 5% atomic. As nitrogen and titanium peaks could not be differentiated by EDX analyses, depth profiling XPS (X-ray Photoelectron Spectroscopy) was used in order to detect nitrogen incorporation into doped-TiO2 layers. XPS analyses of the N-doped films surface and bulk were performed on an Escalab 250 from Thermo Electron. The excitation source was a monochromatic Al K␣ (1486.6 eV). The analyzed surface diameter and etched surface were 400 ␮m and 4 mm2 , respectively. The XPS spectra were calibrated according to the C C binding energy (C1s: 284.8 eV). The power parameters were 2 kV and 1 ␮A. The etching times were 0 s (surface), 40 s, 70 s, 310 s, 550 s, 910 s, 1510 s and 2110 s. The crystalline structure of TiO2 films deposited on silicon and borosilicate was characterized by XRD (X-Ray Diffraction) using a Philips XPert Pan Analytical diffractometer with the copper K␣ rays, Bragg-Brentano geometry. Based on the XRD anatase main peak, the crystallite size could be calculated according to the Scherrer equation in the same conditions as those described in a previous paper from our group [30]. The effect of dopant concentration on Raman peaks shifting of deposited TiO2 was investigated using HORIBA scientific from Xplora with a green laser (25 mW, 532 nm) and Olympus focus camera. The instrumental error on this apparatus was estimated to be less than 1 cm−1 . UV/Vis spectropho-

Fig. 2. Dispersion relations of direct and indirect band gaps in the theory of semiconductors.

tometry using a JASCO V-570 apparatus was implemented on films deposited on borosilicate glass discs to reveal the film absorbance shifting towards visible light range. Band gap energy was calculated using Tauc Plot method [31]. In fact, the energy as a function of electron linear momentum plots shows that there are two types of band gap energies: the direct and the indirect gap [32]. In the case of a direct gap semiconductor, the electronic transfer is due to an external energy supply exactly equivalent to the band gap energy without changing the linear momentum of the electron (k = 0). Whereas, in an indirect band gap semiconductor, the electron transfer from valence to conduction band is possible by an energy supply Eg with modification of this electron linear momentum. This is due to the fact that the energy level of the valence band and the one of the conduction band are not perfectly aligned (Fig. 2). Typically, a Tauc plot shows the quantity h (the energy of the light) on the abscissa and the quantity (␣h)1/r on the ordinate, where ␣ is the absorption coefficient of the material. The value of the exponent r denotes the nature of the transition [33]. Since TiO2 is considered as an indirect gap

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semiconductor, the value of r is ½ and the Tauc plot will represent (␣h)1/2 as a function of energy (eV). Lastly, in order to investigate their photocatalytic activity, films were subjected to the Pilkington test. As described in a previous paper from our group [30], the samples were dip-coated in a stearic acid solution (8.8.10−3 mol L−1 in methanol) for 10 min, withdrawn with a withdrawal rate of 0.5 cm min−1 , dried and back-cleaned. Then they were exposed to visible light with an irradiance value of 75 W/m2 (at 420 nm, using a Van Cliff 1 × 150W IP65 visible light lamp) and the stearic acid degradation (absorption bands in the range of 2800–3000 cm−1 ) was followed by FTIR analysis each 20 min of irradiation.

3. Results and discussion 3.1. Structural and microstructural properties of the deposited films 3.1.1. Deposition rate and film morphology The growth of a PECVD material is controlled by four successive mechanisms: the creation of reactive species in the gaseous phase, the mass transfer of these species to the substrate surface, followed by their adsorption and then recombination on the surface. In addition, undesirable homogeneous reactions, i.e. reactions between reactive species in the gaseous phase, can occur inducing a decrease of deposition rate (while leading to the formation of powder in the plasma reactor). Such homogeneous reactions can be caused by an excess of reactive species in the deposition chamber increasing the probability of species collisions in the plasma phase. To avoid homogeneous reactions while providing a noticeable film growth rate, the partial pressure of precursor has to be adjusted to an optimal value. In this study, the choice of 22.5 Pa as the TTIP + Ar partial pressure has been made taking into account such considerations. The substrate temperature has also an important influence on the film growth rate. In-Sun Lee et al. [34] have mentioned that heterogeneous reactions between reactive species (contributing to the film growth, once the reactants have been transferred to the surface) is dominant at Ts ≤ 200 ◦ C. As the temperature further increases, undesirable homogeneous reactions may occur; in such a case, if the plasma input power is not sufficient to create enough new reactive species, the number of reactive species transferred to the substrate surface decreases leading to the growth rate drop. Fig. 3 shows the deposition rate (thickness (nm)/deposition time (min)) of as-deposited non-doped films on silicon wafer as a function of the substrate temperature for two plasma powers. At low plasma power (64 W), the mass transfer to the substrate surface is faster than the creation of reactive species, in the opposite of what happens at high plasma power (100 W) where the precursor is more fragmented. This explains why, for the substrate temperature equal to 150 ◦ C, the film grows faster at 64 W than at 100 W. At 250 ◦ C, the decrease in the deposition rate observed at 64 W is due to the limitation of the film growth by the homogeneous

Fig. 3. Deposition rates of the thin films deposited at 150 ◦ C, 250 ◦ C, 300 ◦ C and 350 ◦ C for (a) 64 W and (b) 100 W.

reactions as explained by In-Sun Lee et al. This is not the case for the film deposited at 100 W where there are enough reactive species created in the gaseous phase. Some of them could contribute to the film growth in spite of the occurrence of homogeneous reactions. Above 250 ◦ C, the deposition rate increases for films prepared at both plasma input powers. This phenomenon is due to the progressive film crystallization (proved by XRD analyses – see later in subsection 3.1.2) that can be morphologically depicted on SEM pictures of films. Indeed, Figs. 4 and 5 show the SEM cross-sections of as-deposited non-doped films on silicon wafer as a function of the substrate temperature for both 64 W and 100 W. Films thicknesses are in the range 160–540 nm. It can be noticed that the film morphology switches from compact mode to columnar mode with the increase of the substrate temperature (already observed by C. Sarantopoulos [35]), whatever the plasma input power may be. At low frequency plasma, highly energetic ions are generated which leads to local surface bombardment [36]. At low temperature (up to 300 ◦ C), the thermal energy is still insufficient to provoke intense surface etching. At 350 ◦ C, ions are sufficiently energetic to etch the surface; fragmented parts of TTIP grow in the vacant spaces. Thus, successive columns are progressively formed. Moreover, it could be shown by SEM analysis (no illustration provided) that performing the post-annealing step (at 400 ◦ C for 1 h) induces a thickness decrease. In the case of the non-doped film deposited at 350 ◦ C, the thickness decrease was of 120 nm (25% drop) in the case of low power (64 W) and 280 nm (52% drop) in the case of high power (100 W). The layer retraction, which is due to particles coalescence caused by sintering effect, is logically higher in the case of film deposited at the highest plasma power, whose higher deposition rate certainly induces a higher porosity in comparison with the film prepared at low plasma power.

Fig. 4. SEM images of the as-deposited non-doped films deposited at 64 W and (a) 150 ◦ C, (b) 250 ◦ C, (c) 300 ◦ C and (d) 350 ◦ C on silicon substrate.

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Fig. 5. SEM images of the as-deposited non-doped films deposited at 100 W and (a) 150 ◦ C, (b) 250 ◦ C, (c) 300 ◦ C and (d) 350 ◦ C on silicon substrate.

SEM observations could not enable to differentiate films prepared without ammonia or from different ammonia partial pressures in terms of morphology and thickness. Performing a topographic study of films deposited on borosilicate glass by AFM, it can be shown that the substrate surface influences the final film surface roughness, being also affected by the ammonia partial pressure in the plasma reactor. Figs. 6 and 7 show the AFM images and average surface roughness of films deposited at 64 W or 100 W and 350 ◦ C with 5, 10 or 15 Pa NH3 . The average surface roughness on borosilicate strips was calculated according to the formula: n 

Ra = 1/n

|yi|, where n is the number of scanned surface points

i=1

and yi is its corresponding roughness (Fig. 8). The thin layers deposited on borosilicate glass strips show a relatively smooth topography, directly related to the nature of the substrate which has been manufactured using a draw process that produces a final product with a high quality, flat surface, low roughness and excellent optical properties. The observed surface nucleation is directly related to size of fragments contributing to the film growth. For the samples deposited at 64 W, the increase in surface roughness when the NH3 pressure increases from 5 Pa to 10–15 Pa is due to the fact that, at a such low plasma input power, the fragments contributing to the film growth are quite big (low fragmentation), and can partially recombine in gaseous phase to form even bigger fragments. This phenomenon is all the most pronounced than the total pressure (partly due to the NH3 presence) in the reactor increases. But since the deposited clusters are not highly fragmented and the borosilicate surface is not porous, this causes a surface roughness increase. At 100 W, the surface roughness seems to be more stable at different ammonia partial pressures, certainly due to the fact that, at such high input plasma power, fragments contributing to the film growth are much smaller; their size is less influenced by the total pressure in the reactor. 3.1.2. Crystallinity Fig. 9 shows the influence of the preparation conditions on the XRD patterns of non-doped thin films. The peak observed at around 2␪ = 33◦ on some patterns (those related to the thinnest films) is specific to the silicon substrate and is not related to the films. As-deposited films synthesized at 150 ◦ C, 250 ◦ C and 300 ◦ C have been found in amorphous phase for both powers 64 W and 100 W. Post-annealing these samples under air for 1 h at 400 ◦ C allows the phase conversion to anatase as it can be seen from the (101) peak at 2␪ = 25.3◦ ; no rutile phase was detected. It can be noticed that increasing the substrate temperature makes the crystallization easier as the anatase main peak becomes more intense. The XRD patterns of the as-deposited films synthesized at 350 ◦ C (at 64 W or 100 W) and anatase TiO2 match completely, which means that anatase phase could be obtained without post-annealing when heating the substrate at 350 ◦ C dur-

ing the deposition step. This result is in accordance with one previous study from our group dealing with High-temperature XRD (HT-XRD) analysis of TiO2 films prepared in a radio-frequency plasma discharge, showing that the spontaneous crystallization of TiO2 thin films occurs between 300 ◦ C and 400 ◦ C [15]. Comparing XRD peaks of the samples prepared in this study (in a low-frequency plasma discharge) to the ones of the films deposited at radiofrequency (previous study [15]), it can be seen that the (101) anatase peak is more intense in the case of the low-frequency discharge (this study) while the (004) anatase peak (that should take place at 2␪ = 37.9◦ ) is almost absent. This is a proof that with lowfrequency PECVD, preferential crystallite growth along the (101) plane can be obtained. Sharp (101) XRD peaks are obtained for the TiO2 deposited at 350 ◦ C. This confirms the hypothesis that sufficient thermal energy leads to TTIP dissociation and TiO2 crystallization without post-annealing at 350 ◦ C. The crystallite sizes of the post-annealed deposited non-doped films prepared on silicon wafer have been calculated using Scherrer equation. The sizes were calculated at 300 ◦ C and 350 ◦ C for both 64 W and 100 W because these are the only two temperatures where we observe the anatase main peak for both powers after post-annealing. First, it can be noticed that crystallites size is less than 50 nm, whatever the plasma parameters. Crystallite sizes are equal to 35 nm and 40 nm for films deposited at 300 ◦ C for 64 W and 100 W, respectively while the have been found equal to 31 nm and 34 nm for films prepared at 350 ◦ C also for 64 W and 100 W, respectively. The crystallites size decrease between 300 ◦ C and 350 ◦ C is an indication of transition between amorphous and crystalline phase. Deposition at 100 W induces slightly larger crystallites than deposition at 64 W due to the more homogeneous precursor fragmentation and deposition on substrate surface. In order to get information on the nitrogen incorporation into the titanium dioxide lattice when ammonia is added in the plasma reactor, XRD analysis was performed on both the non-doped and doped samples deposited on borosilicate substrate in PECVD conditions enabling crystallization without post-annealing (Ts = 350 ◦ C). The main difference between the non-doped and the doped samples is the anatase peak intensity decrease and broadening after nitrogen incorporation, especially at 64 W (Fig. 10). This indicates that the orderliness of the crystal lattice of the anatase decreases by the doping process. Increasing the NH3 partial pressure, this amorphization phenomenon becomes more pronounced, above all at 64 W. This evolution can be ascribed to the replacement of oxygen atoms by nitrogen atoms in these materials. In addition to these phenomena, a peak shifting towards lower diffraction angles (1.2% angle variation at 64 W and 2% angle variation at 100 W) is observed while increasing the dopant concentration in the deposition chamber. Such a peak shifting is probably caused by a compressive stress inside the lattice and change in the chemical composition of the thin films due to the nitrogen incorporation. The lattice crystalline structure at 100 W seems to stay more stable than that of the one at 64 W after the doping process. In fact, at high plasma power, the film’s intrinsic inorganic properties are

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Fig. 6. N-TiO2 films deposited on borosilicate glass at 350 ◦ C, 64 W and (a) 5 Pa NH3 , (b) 10 Pa NH3 and (c) 15 Pa NH3 .

Fig. 7. N-TiO2 films deposited on borosilicate glass at 350 ◦ C, 100 W and (a) 5 Pa NH3 , (b) 10 Pa NH3 and (c) 15 Pa NH3 .

31–35 nm, is barely modified by the increased dopant pressure for 64 W and 100 W, all the samples being deposited at 350 ◦ C on borosilicate glass. More precisely, an increase of the dopant pressure induces a slight increase of the crystallite size. Deposition at 100 W induces slightly larger crystallites than deposition at 64 W, as previously proved in the case of silicon as substrate.

Fig. 8. Average roughness variation by 5, 10 and 15 Pa NH3 doping pressures at (a) 64 W and (b) 100 W for the films deposited on borosilicate glass at 350 ◦ C.

enhanced as proved by spectroscopic analyses (see later Subsection 3.1.3 and 3.1.4). Thus, the TiO2 network is more cross-linked and so less suitable for facilitating nitrogen incorporation into the lattice. First the crystallite sizes between the non-doped and the doped samples have been compared. The results are presented in Table 1. It can be noticed that the crystallite size, found in the range

3.1.3. Chemical structure by FTIR and raman spectroscopies In order to study the influence of the preparation conditions on the Ti O abundance and crystallization degree in films, FTIR analysis (transmission mode) was performed on the as-deposited and annealed non-doped samples. Fig. 11 represents the normalized FTIR spectra of films deposited at different substrate temperatures and discharge powers. The skeletal Ti O bond vibration is in the band below wavenumber 1000 cm−1 . In this fingerprint region, a clear difference appears between amorphous TiO2 and crystalline anatase; more precisely, a sharp absorbance at 450 cm−1 is characteristic for anatase [15]. As shown, quite independently on the PECVD parameters, the as-deposited films (black colored spectra) present a quite broad Ti O band around 500 cm−1 , more characteristic of amorphous phase. Whereas the post-annealed films (grey colored spectra) show a sharper and more intense Ti O band at 450 cm−1 which can be assigned to anatase. In addition, it can be noticed from FTIR results that increasing both substrate temperature and plasma power makes the Ti O band intensity increasing. In parallel, it can be observed that the O H stretching band between 3000 and 3500 cm−1 (constitutive of the film bulk or characteristic of adsorbed water on the film surface and the presence of Ti OH

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Fig. 9. XRD patterns of the as-deposited and annealed TiO2 films deposited on silicon substrate at 64 W and 100 W for substrate temperatures of 150 ◦ C, 250 ◦ C, 300 ◦ C and 350 ◦ C. Table 1 Crystallite size (nm) for non-doped, 5 Pa NH3 doped, 10 Pa NH3 doped and 15 Pa NH3 doped TiO2 films deposited at 350 ◦ C at both 64 W and 100 W on borosilicate glass. Sample

Non doped

Doped 5 Pa NH3

Doped 10 Pa NH3

Doped 15 Pa NH3

Crystallite size at 64 W Crystallite size at 100 W

31 34

33 34

33 34

34 35

bonds) decreases when both substrate temperature and plasma power increase. So the purest and richest Ti O composition has been found for the deposit synthesized at 350 ◦ C and 100 W. As a complementary observation, some FTIR spectra corresponding to samples prepared at high substrate temperatures (300 ◦ C and

above all 350 ◦ C) show shifted baselines, which could be due to detachment of films from the silicon support in drastic preparation conditions. In order to characterize the incorporation of nitrogen in the anatase film prepared at 64 or 100 W and 350 ◦ C (without post-

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Fig. 10. XRD patterns focusing on the anatase main peak for the samples deposited on borosilicate glass at 350 ◦ C doped by 5 Pa NH3 , 10 Pa NH3 and 15 Pa NH3 at (a) 64 W and (b) 100 W.

annealing), FTIR (transmission mode) analyses were performed on the 5 Pa, 10 Pa and 15 Pa NH3 corresponding doped films deposited on silicon. Results are presented in Fig. 12. The strong and sharp band at 500 cm−1 represents the lattice vibration of the TiO2 prepared in this work. It can be noticed that the absorbance of Ti O vibrations at around 500 cm−1 decreases as NH3 concentration increases. The lattice vibration of both treated and untreated TiO2 remains at the same position as the samples were all prepared under the same conditions. This bond weakening might have been caused by the formation of the metal-amine complex (Ti N), depicted at 612 cm−1 [37,38], proving the nitrogen doping. The band characteristic of Ti O bond at 500 cm−1 even disappears in the 10 Pa or 15 Pa NH3 doped TiO2 at 64 W and in the 10 Pa NH3 doped TiO2 at 100 W. Ti O bonds still exist at 15 Pa for films deposited at 100 W. In fact, according to the analysis made by Saha et al. [39], this remaining band could be attributed to an additional oxidation of the N-TiO2 film or to some TiO2 that could have deposited on top of the nitrided titania grains leading to the simultaneous presence of Ti N and Ti O bands (Fig. 13). Further studies by Raman and XPS analyses will be presented afterwards and will enable to deepen the information about the film chemical structure. Quantitatively, it is logically observed that the increase of the Ti N band intensity is proportional to the decrease of the Ti O band intensity. It is worth noting that the N H bending bands of the doped titanium dioxide, which are represented by absorption bands at 1620–1500 cm−1 , slightly shift to lower frequencies compared to absorption bands referenced in the literature for N H at 1644 and 1599 cm−1 [40]. This shifting is due to the weakening of the N H bond by the formation of Ti N bond. Fig. 14 shows the Raman spectra of the non-doped and doped TiO2 thin films prepared at 350 ◦ C, 64 W or 100 W, and different ammonia partial pressures. According to the data reported in the references [41,42], the anatase phase of TiO2 has six Raman bands at 144 cm−1 (Eg ), 197 cm−1 (Eg ), 399 cm−1 (B1g ), 513 cm−1 (A1g ), 519 cm−1 (B1g ) and 639 cm−1 (Eg ). The Raman spectra of all deposited samples are the same as that of the pure anatase phase, with no bands due to the rutile phase. It is known that if chemical bond length of molecules changes due to any internal or external effect, then it may lead to wavenumber shifting in Raman spectra. A shorter bond length causes shifting to a higher wavenumber and vice versa. In the Raman spectra of the films deposited at both 64 W and 100 W, it can be noticed a slight shifting towards lower wavenumbers when adding NH3 in

the deposition chamber. In fact, Ti O bonds have a length equal to 1.96 Å whereas Ti N bonds are 2.12 Å long [37]. So the observed shifting means that Ti O bonds in TiO2 lattice are being replaced progressively by longer bonds as nitrogen is incorporating into the lattice. This bond length increase could be attributed to the formation of Ti N bonds but this progressive replacement is more probably related to the intermediate formation of the Ti O N bond. In fact, according to theoretical DFT calculations performed by Saha et al. [39], there is an overall increase in Ti O bond distances in doped samples which can be attributed to the dominant effect of nitrogen occupying the interstitial position. By that, we cannot directly attribute this length increase to Ti N formation because Ti O in doped samples is also longer than Ti O in the pure anatase phase. For the doped samples prepared at 64 W and 100 W, the anatase bands shift at lower wavenumbers as the NH3 pressure is increased. If we correlate Raman and FTIR in this case, oxygen is partially getting substituted by nitrogen. Indeed, as the dopant concentration is increased, Ti O band in FTIR disappears progressively. The particular point to discuss is the Ti O bond still remaining in the sample deposited at 350 ◦ C, 100 W and 15 Pa NH3 doped. It can be supposed that the Raman shift towards lower wavenumbers in this case is due to both the formation of Ti N bonds and the extension of Ti O bonds by the interstitial nitrogen (Ti N O). It has still to be investigated whether the FTIR Ti O bond is exclusively due to the Ti N O layer or a mix of TiO2 /Ti O N phases. This question is really important since the electrochemical reactions generally occur on the material surface. It should be noted that the sample was analyzed using a green laser (532 nm) so the wavelength may have probed only some layers if the film is multilayered (to say TiO2 /Ti N O/Ti N). The Ti2p XPS spectra presented in the next part will reveal the material state. 3.1.4. Nitrogen content and local structure by XPS Using the classical surface XPS analysis, the quantity of nitrogen depicted in both non-doped and doped samples is not very different. It slightly increases from 0.55 at.% to 0.93 at.% for films prepared at 350 ◦ C and 64 W, and from 0.79 at.% to 1.36 at.% for those synthesized at 350 ◦ C and 100 W. The fact that some nitrogen could be depicted in non-doped samples can certainly be ascribed to an atmospheric contamination by N2 adsorption on the surface when preparing films at high temperature (350 ◦ C). This observation can be confirmed by the N1s peak at binding energy around 400 eV (Fig. 15). Supposing this contamination also concerns doped samples, it is difficult to know if the small difference of nitrogen quantity between non-doped and doped samples is due to nitrogen

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Fig. 11. FTIR spectra of the as-deposited and annealed TiO2 films deposited on silicon substrate at 150 ◦ C, 250 ◦ C, 300 ◦ C and 350 ◦ C for (a) 64 W and (b) 100 W.

incorporation or to surface contamination enhancement. In order to investigate which type of nitrogen incorporation is happening in the lattice, the elemental composition and chemical nature of

N-doped TiO2 films vs. non-doped ones were probed by XPS analyses performed on the surface and also in the bulk (after etching for different durations). Table 2 presents the N1s binding energies

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Fig. 12. FTIR spectra of the doped films deposited on silicon substrate at 5, 10 and 15 Pa NH3 for (a) 350 ◦ C, 64 W and (b) 350 ◦ C, 100 W.

Fig. 13. Schematic representation of the film microstructure for additional oxidation of Ti N in the sample prepared at 350 ◦ C, 100 W and 15 Pa NH3 .

Fig. 14. Raman spectra for the non-doped and 5, 10 and 15 Pa NH3 doped TiO2 at 350 ◦ C and (a) 64 W and (b) 100 W.

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Fig. 15. Depth profiling XPS N1s binding energy for the N-TiO2 samples deposited at 350 ◦ C on glass for (a) 64 W, 5 Pa NH3 , (b) 64 W, 15 Pa NH3 , (c) 100 W, 5 Pa NH3 , and (d) 100 W, 15 Pa NH3 . Table 2 N1s binding energy positions in TiO2 lattice. N1s binding energy (eV)

Peak bond attribution

396–398 397.7 399–400 >400

Ti N Ti (substitutional N-doping) [44] Ti N Ti (Oxygen simple replacement by nitrogen in the TiO2 crystal lattice) [46] Ti O N/Ti N O (oxidized nitrogen lattice-interstitial positions) [46] Adsorbed N2 or N H from NH3 gas [45]

comprised in the range 396–404 eV, in accordance to those given in the literature [43–45]. The XPS depth profiles are presented in Fig. 15 in the case of 5 Pa and 15 Pa NH3 for both 64 W and 100 W (only the films prepared from both extreme dopant concentrations were analyzed). In the above XPS spectra (remarkably in (a), (c) and (d) ones), a progressive N1s binding energy shifting from 400 eV (corresponding to the oxidized intermediate Ti O N layer) on the material surface towards 396–397 eV while getting deeper in the bulk can be noticed. The remarkable point here is for the sample prepared at 350 ◦ C, 64 W and doped by 15 Pa NH3 (Fig. 15, sample (b)). The 396–397 eV peak, corresponding to the substitutional N-doping (Ti N), appears even on the material surface. This phase is continuous throughout the bulk. This is in perfect accordance with the FTIR spectra (Fig. 12) for this sample (where Ti O bond is completely replaced by Ti N bonds) and indicates that the wavenumber decrease in

the Raman spectrum comes from the substitution of Ti O bond by Ti N. Whereas for the same sample at low dopant concentration (5 Pa NH3 ), a Ti N O or/and Ti O N firstly exists on the material surface (wavenumber decrease in Raman spectroscopy due the Ti O bond extension by interstitial nitrogen) then the Ti N starts to form progressively when passing through the bulk [40]. The titanium Ti2p spectra for both samples deposited at 350 ◦ C and 64 W (5 and 15 Pa NH3 ) confirm the previous observations (Fig. 16(a) and (b)). The black curve corresponds to the material surface (no etching). At this stage, the film is exclusively composed of a crystallized anatase TiO2 phase (doublet sharps peaks at 458.5 eV, 464.1 eV for the 5 Pa NH3 film and 459 eV, 464.3 eV for the 15 Pa NH3 material) [39]. As the etching time increases (probing into the depth), this doublet intensity decreases and shifts to lower binding energies (phase amorphization, compatible with XRD observations) giving place to peaks corresponding to Ti2p bounded to nitrogen at 455 eV and 461 eV [40]. It can be remarked that the layer is thicker for

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Fig. 16. Titanium Ti2p spectra for N-doped TiO2 (350 ◦ C) deposited at 64 W and (a) 5 Pa NH3 or (b) 15 Pa NH3 .

Fig. 17. Titanium Ti2p spectra for N-doped TiO2 films (350 ◦ C) deposited at 100 W and (a) 5 Pa NH3 or (b) 15 Pa NH3 .

Fig. 18. O1s binding energy for the 100 W, 350 ◦ C, doped TiO2 (a) 5 Pa NH3 and (b) 15 Pa NH3 .

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Fig. 19. UV–vis absorption spectra of the non-doped and 5, 10 Pa NH3 N-doped samples deposited at (a) 64 W, 350 ◦ C and (b) 100 W, 350 ◦ C.

the 15 Pa NH3 doped sample (because the nitrided phase can be observed even at 40 s etching time) which explains the absence of Ti O FTIR peak and the Raman shift towards lower wavenumbers at this dopant pressure. The Ti2p spectra for N-doped TiO2 thin films (5 and 15 Pa NH3 ) deposited at 100 W, 350 ◦ C have also been plotted in order to explain the Ti O FTIR peak evolution and the reason behind the Raman shift towards lower wavenumbers (Fig. 17). At 100 W, as at 64 W, the film surface is exclusively composed of a crystallized anatase TiO2 phase (doublet sharps peaks at 458.7 eV, 464.2 eV for the 5 Pa NH3 film and 459 eV, 464.5 eV for the 15 Pa NH3 material). This doublet intensity decreases with the progressive growth of the nitrided phase (at 455 eV, 460.7 eV for 5 Pa NH3 and 455 eV, 460.9 eV for 15 Pa NH3 ). This could be related to the Ti N peak observed in the FTIR spectra for films prepared at 100 W. As the etching time increases, the oxide phase peaks shifting is less pronounced than that at 64 W. This is in accordance with the XRD patterns presented previously where the crystalline structure of the 100 W, 350 ◦ C samples (at 5, 10 and 15 Pa NH3 ) is barely affected by the doping process. It can be remarked one more time that the nitrided layer is thicker for the 15 Pa NH3 doped sample which is a normal phenomenon directly related to the dopant concentration increase. Another observation at 100 W is that, in addition to the oxide layer and the underlying the nitrided phase, there is a feature called “intermediate” where Ti2p spectra present a binding energy of 457 eV and the N1s spectra present, in addition to the peak at 395–396 eV, a broad feature at 400–401 eV even in the deep layers (see the N1s binding energy spectra for 100 W in Fig. 15). These peaks were not observed in the case of the doped films deposited at 64 W. This suggests a probability of Ti O N or/and Ti N O phase that exists simultaneously with the Ti N bonds. The O1s binding energy for the 100 W, 350 ◦ C, doped 5 and 15 Pa NH3 is presented in Fig. 18. The broad peak at 530 eV and 534.5 eV for both samples surface is related to the O1s in the TiO2 crystalline phase formed on the top of the nitrided phase. It can be observed a peak shifting from 530 eV to around 531 eV and the disappearance of the peak at 534 eV when the analysis has been performed deeper into the layer. This oxygen is attributed to the O1s in the oxidized nitride phase (Ti O N/Ti N O) according to the theoretical model of Saha et al. [39]. It can be seen now why the Ti O bond remains at 100 W after nitrogen doping; it is because of that intermediate state that co-exists with the Ti N phase (the presence of this phase is proven by substitutional N1s at 396 eV). So the Raman wavenumber decrease is due to both the extension of Ti O by nitrogen insertion and Ti N longer bonds. Comparing the doping modes

Fig. 20. UV/Vis adsorption spectrum of the commercial TiO2 (Evonik P25) as a reference [47].

between 64 W and 100 W, is can be seen that nitrogen prefers substitutional sites at 64 W while both substitutional/interstitial sites can be occupied at higher plasma input (100 W). 3.2. Functional properties in relation with photocatalysis and water-splitting application 3.2.1. Effect of doping on UV–vis absorption UV–vis absorption was investigated on non-doped and doped samples in order to depict the effect of nitrogen doping on such property. Fig. 19 show the UV–vis absorption spectra of the nondoped anatase films deposited at 64 W and 100 W (350 ◦ C) and two N-doped TiO2 films prepared with 5 and 15 Pa NH3 in the same PECVD conditions. For the non-doped samples, the curve interpolation appears at 387.5 nm which is the wavelength that corresponds to the anatase band gap energy of 3.2 eV [46]. The UV/Vis absorption spectrum of the usual TiO2 reference (powder P25 from Evonik) is presented in Fig. 20 as a reference [47]. Comparing this spectrum with those characteristic of deposited non-doped films, it can be noticed that the non-doped thin films (for both powers) present almost the

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Fig. 21. Tauc plot calculation comparison between the non-doped and 15 Pa NH3 doped samples prepared are (a) 64 W, 350 ◦ C and (b) 100 W, 350 ◦ C.

Fig. 22. Stearic acid degradation evolution under visible light for silicon substrate, non-doped TiO2 , 5 Pa and 15 Pa NH3 doped TiO2 prepared at (a) 64 W, 350 ◦ C and (b) 100 W, 350 ◦ C.

same cut-off as the reference (around 380 nm) which confirms the anatase phase growth by LF-PECVD. As the NH3 partial pressure is increased, the sample absorption decreases in the UV range and shifts towards higher wavelengths up to 550 nm for the 15 Pa NH3 –100 W sample and up to 600 nm for the 15 Pa NH3 –64 W sample. So the cut-off is widely shifted into the visible light region after in-situ nitrogen doping process. In fact, nitrogen doping introduces some new energy levels at the valence band edge acting as shallow donors and thus making the original band gap of TiO2 smaller. That is what is called band gap tailoring. The reason behind the best band gap shifting at 64 W (15 Pa NH3 ) is certainly due to the fact that the nitrogen incorporation into the lattice occurs starting from the material surface and remains continuous into the material bulk (according to the XPS depth profiles shown before) for this particular film. The band gap was calculated using Tauc Plot method. The results are presented in Fig. 21 where ␣ is the ratio of absorbance and thickness and E is the energy deduced from wavelength values. The TiO2 band gap energy has been decreased from 3.2 eV to 2.1 eV for the sample doped with 15 Pa NH3 at 100 W, and to 1.9 eV for the sample doped with 15 Pa NH3 at 64 W. 3.2.2. Effect of doping on photocatalytic activity The non-doped and doped samples synthesized at 64 W and 100 W (350 ◦ C) were submitted to Pilkington test in order to study

their photocatalytic activity. Fig. 22 shows the evolution of the degradation percentages of stearic acid as a function of irradiation time for all analyzed samples. No acid degradation was observed on the non-coated silicon substrate. It is also observed that the non-doped samples at both 64 W and 100 W present really low photocatalytic activity under visible light irradiation. In fact, a maximum degradation yield of 20% is obtained in the case of the 100 W TiO2 after 180 min of visible light irradiation while the acid degradation could barely reach 10% after the same irradiation time. This proves once again the necessity to enhance the visible light efficiency of this semiconductor by the doping approach. Increasing the doping agent concentration (from 5 to 15 Pa) leads to a better photocatalytic activity in the visible light for both powers. A faster degradation is obtained for the films deposited at 100 W, whereas nitrogen incorporation is more homogeneous in films deposited at 64 W than for those prepared at 100 W (as previously depicted by XPS analysis). This apparent contradiction is directly related to the fact that the stearic acid degradation efficiency may be mainly controlled by the feasibility for visible light to penetrate the films more than the nitrogen concentration and distribution in the films structure. According to the SEM results, the films deposited at higher plasma power have a more columnar and so porous microstructure than those deposited at 64 W. So the light can reach the material bulk more easily for the samples prepared at high power. It also seems that the simultaneous presence

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of Ti N/Ti O N phases (two types of doping) at 100 W fastens the photocatalytic activity under visible light irradiation. Still, the reaction efficiency for N-TiO2 (15 Pa NH3 ) remains almost similar at 64 W and 100 W since a 100% stearic acid degradation is reached after 2 h (120 min) illumination for both types of samples. Ming Zhou et al. [48] performed the same test for bare silicon substrate and anatase coated silicon substrate under UV illumination. 100% stearic acid degradation had been reached in 100 min. Uk Lee et al. [49] were able to reach 99% dye degradation on C and S doped nanotubes after 70 min visible light illumination. Kun et al. [50] have studied the photocatalytic activity of phenol (0.5 mmol L−1 ) under visible light with N-doped TiO2 and found that acid degradation was complete within 2 h of irradiation. By that, according to what has been published in the literature, the N-TiO2 thin films synthesized by Low Frequency PECVD process show a good photocatalytic activity in the visible light region.

4. Conclusions This study has proved the feasibility to obtain single-oriented pure anatase phase without post-annealing when depositing TiO2 films from titanium tetra-isopropoxide at substrate temperature as low as 350 ◦ C in a low-frequency plasma discharge (40 kHz). When the substrate temperature is lower than 350 ◦ C, films are amorphous and should by thermally post-treated to be crystallized. With the increase of the temperature up to 350 ◦ C, the film growth switches progressively from compact to columnar mode and the chemical structure refine towards a purer TiO2 material reaching a single-oriented (100) anatase phase at 350 ◦ C. At high plasma power (100 W), the deposition rate and the crystallites size are higher than at low plasma power (64 W), leading to a more porous material. This work has also demonstrated the feasibility to introduce nitrogen species into the titanium dioxide lattice, using ammonia as additional gas in the plasma chamber. Incorporation of nitrogen in films as doping element (formation of Ti N Ti in any plasma conditions, with addition of Ti O N/Ti N O bonds at 100 W) has been proved combining results from FTIR and Raman spectra, XPS Ti2p energy changes and XRD peak shifting. At low plasma power (64 W), the nitrogen doping is more homogeneous from the materials surface to the bulk. The size of nano-crystals in prepared N-doped anatase films (deposited on borosilicate) has been estimated to be in the range 31–35 nm (depending on the synthesis conditions) by applying the Scherrer equation. Increasing the added ammonia quantity, the N-doped TiO2 band gap energy has been tailored to absorb wavelengths in the visible range of solar spectrum (up to 600 nm for the more homogeneous film deposited at 64 W). Consequently, N-doped TiO2 films have been proved as efficient for the photocatalytic degradation of stearic acid under visible light (with a faster rate for the more porous film prepared at 100 W). As prospect, N-doped anatase layers will be deposited on porous glass supports and assembled with proton-exchange membrane and Pt cathode to be tested as photocatalytic anode in plasma multilayered devices for H2 production.

Acknowledgments This work’s financial and administrative support is provided by the Doctorate School of Sciences and Technologies (EDST), Lebanese University (UL) in collaboration with AZM&SAADÉ foundation. We also thank Roland HABCHI (PR2N, Beirut) for RAMAN and AFM analyses, Valérie FLAUD (ICGM, Montpellier) for XPS analysis, Didier COT (IEM, Montpellier) for SEM analysis and Arie VAN DER LEE (IEM, Montpellier) for XRD analysis.

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Please cite this article in press as: L. Youssef, et al., Optimization of N-doped TiO2 multifunctional thin layers by low frequency PECVD process, J Eur Ceram Soc (2017), http://dx.doi.org/10.1016/j.jeurceramsoc.2017.05.010