Structural and optical characterization of pure Si-rich nitride thin films
© Debieu et al.; licensee Springer. 2013
Received: 26 November 2012
Accepted: 29 December 2012
Published: 16 January 2013
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© Debieu et al.; licensee Springer. 2013
Received: 26 November 2012
Accepted: 29 December 2012
Published: 16 January 2013
The specific dependence of the Si content on the structural and optical properties of O- and H-free Si-rich nitride (SiNx>1.33) thin films deposited by magnetron sputtering is investigated. A semiempirical relation between the composition and the refractive index was found. In the absence of Si-H, N-H, and Si-O vibration modes in the FTIR spectra, the transverse and longitudinal optical (TO-LO) Si-N stretching pair modes could be unambiguously identified using the Berreman effect. With increasing Si content, the LO and the TO bands shifted to lower wavenumbers, and the LO band intensity dropped suggesting that the films became more disordered. Besides, the LO and the TO bands shifted to higher wavenumbers with increasing annealing temperature which may result from the phase separation between Si nanoparticles (Si-np) and the host medium. Indeed, XRD and Raman measurements showed that crystalline Si-np formed upon 1100°C annealing but only for SiNx<0.8. Besides, quantum confinement effects on the Raman peaks of crystalline Si-np, which were observed by HRTEM, were evidenced for Si-np average sizes between 3 and 6 nm. A contrario, visible photoluminescence (PL) was only observed for SiNx>0.9, demonstrating that this PL is not originating from confined states in crystalline Si-np. As an additional proof, the PL was quenched while crystalline Si-np could be formed by laser annealing. Besides, the PL cannot be explained neither by defect states in the bandgap nor by tail to tail recombination. The PL properties of SiNx>0.9 could be then due to a size effect of Si-np but having an amorphous phase.
Since the discovery of efficient visible photoluminescence (PL) of silicon nanoparticles (Si-np) due to quantum confinement effects (QCE) , the possibility of bandgap engineering of Si-based materials through the Si-np size control makes Si-based nanostructured material attracting for future applications in optoelectronics as low-cost, miniaturized, and CMOS-compatible, light-emitting devices (LEDs), laser, as well as photovoltaic devices. In the past, researches were focused on luminescent Si-np embedded in Si oxide media. However, the insulating nature of Si oxide remains a barrier for the production of future electrically pumped LEDs and efficient photovoltaic cells. This detrimental aspect can be overcomed to an extent, using a higher conductive host medium like Si nitride which has a lower bandgap energy than SiO2.
The first results on Si nitride are promising since many researchers have reported on efficient visible PL with tunable light emission via the change of the Si nitride composition. However, it also turns out that N-rich nitride [2–4] and Si-rich nitride thin films containing amorphous [5–8] or crystalline [9–14] Si-np or without Si-np [15–18] can exhibit PL in the same spectral range. As a result, the mechanism of the PL in Si nitride is still a controversial subject in the literature. QCE in amorphous or crystalline Si-np, defect states in the bandgap, and band tail recombination have been proposed to account for the PL. However, since the synthesis methods were mostly based on chemical vapor deposition techniques, most of the films contained a significant amount of hydrogen [2, 5, 8, 10, 11, 13, 14, 16] and, in some cases, of oxygen [19, 20], which can both contribute to the PL. Consequently, it is difficult to experimentally distinguish the mechanisms of the PL.
Then, this article is significant since we report on the structural and optical properties of Si-rich SiNx<1.33 thin films devoid of hydrogen and oxygen. The films were deposited by radio frequency (RF) magnetron sputtering. The excess of Si incorporated during the sputtering process makes possible the formation of Si-np during a suitable annealing. The microstructural properties of the films with regard to the composition and the annealing temperature are investigated. The possible contributions of the Si nitride medium and of Si-np formed during thermal annealing, or laser annealing, on the origin of the PL are discussed notably as a function of the Si-np phase (crystalline or amorphous).
In this work, pure amorphous Si-rich SiN x thin films were deposited on p-type 250-μm-thick (100) Si wafers and on fused silica substrates by two methods of RF magnetron sputtering using argon as the main sputtering gas. The films were deposited either by N2-reactive sputtering of a Si target or by co-sputtering of Si3N4 and Si targets. The Si content was monitored either by the N2/Ar partial pressure ratio (≡Ar/N2) or by the RF target power ratio PSi/(PSi + PSi3N4) ≡ Si/Si3N4. The grown temperatures were 200°C and 500°C, and the plasma pressures were 2 and 3 mTorr. We adjusted the deposition time to ensure that the films thicknesses were of the same order of magnitude (100 to 200 nm) in order to avoid any effect on the optical and structural properties. The films were subsequently annealed in a N2 gas flow in a tubular furnace during 1 h.
The layer compositions were determined by Rutherford backscattering spectrometry (RBS). RBS measurements were carried out at room temperature using a 1.9 MeV 4He+ ion beam with an incident direction normal to the sample surface. The backscattered ions were collected at a scattering angle of 165°. The analysis of the RBS spectra, which were performed using the simulation code SIMNRA , enables us to quantify (a) the atomic fraction of the various elements with an accuracy of 0.8 at.% for Si and N and 0.2 at.% for Ar and (b) to determine the atomic areal densities of the films. The infrared absorption properties were investigated by means of a Thermo Nicolet (Nexus model 670) Fourier transform infrared (FTIR) spectrometer. The band positions were obtained by fitting the data with Gaussians. The film microstructure was investigated by Raman spectroscopy with a 532-nm continuous-wave laser illumination with a spot diameter of 0.8 μm. Several neutral density filters were employed to tune the excitation power density from 0.14 to 1.4 MW/cm2. A dispersive Horiba Jobin-Yvon Raman spectrometer with a resolution of 1.57 cm−1, equipped with a confocal microprobe and a CCD camera, was used to acquire the Stokes scattering spectra of the thin layers that were exclusively deposited on fused silica substrates. We also studied the film microstructure by X-ray diffraction (XRD) using a Phillips X’PERT HPD Pro device with Cu K λ radiation (λ = 0.1514 nm) at a fixed grazing incidence angle of 0.5°. Asymmetric grazing geometry was chosen to increase the material volume interacting with the X-ray beam and to eliminate the contribution of the Si substrate. Moreover, the structure was investigated by high-resolution transmission electron microscopy (HRTEM) on cross-sectional samples using a JEOL 2010F (200 kV) microscope.
The optical properties of the films were investigated by spectroscopic ellipsometry using a Jobin-Yvon ellipsometer (UVISEL) with an incident angle of 66.2°. The experimental data were fitted by a dispersion law based on the Forouhi-Bloomer model for amorphous semiconducting and insulating materials  using the DeltaPsi2 software  which determines the refractive index n, the absorption coefficient α versus the photon energy, and the layer thicknesses. The PL spectra were measured using the 457-nm lines of an Ar+ ion laser (12.7 W/cm2) and a fast Hamamatsu photomultiplier after dispersion of the light in a Jobin-Yvon TRIAX-180 monochromator. The PL measurements were corrected from the spectral response of the PL setup.
We first report on the combined analysis of the SiN x film composition by RBS and ellipsometry. Then, the microstructure and the optical properties of the films are investigated as a function of the composition, as well as the annealing temperature.
Nevertheless, one can notice in Figure 3 that our experimental results progressively diverge from the models obtained by this group and also by Hasegawa et al.  while x is decreased. However, the two groups both studied hydrogenated SiN x films (SiN x :H) in contrast to our results. Besides, these latter authors have shown that the Si-H density increased while x was experimentally decreased. Consequently, the drop of n is explained by the H incorporation in their material as suggested elsewhere . However, we could use this model to fit the experimental data but using the refractive index of a-Si (na-Si = 4.37, see Figure 2) instead of hydrogenated a-Si (na-Si:H = 3.3) used by Bustarret et al. . This shows again the influence of H on the optical properties of the films. We obtained = 1.85, which is similar to many previous results [25–27], but is lower than 2.03 that is commonly used for a-Si3N4. This difference could be explained by the incorporation of voids in the microstructure  as attested by the presence of residual Ar atoms detected by RBS in the as-deposited films. Besides, this explanation is confirmed by the density ρ v of our SiN x films which was calculated using the atomic areal density ρ s , and the film thickness d, obtained by RBS and ellipsometry analyses, respectively, with the following relation: ρ v = ρ s / d. We found that the density varied from 2.4 to 2.8 g/cm3, which is again sensibly lower than that of a-Si3N4 of 3.1 g/cm3 reported in the literature .
Considering the RBS and the ellipsometry spectra, we have produced thin SiN x films with various compositions that do not depend on the synthesis method, but only on the Si content. As a consequence, n is a precise indicator of the composition that will be used in the following sections.
Similar blueshifts of the TO band [5, 25, 27, 32–34] and of the LO band [24, 27, 33] were also observed in SiN x :H films. Lucovsky et al.  explained the TO band blueshift by the incorporation of H. They suggested that one of the near silicon atoms of the N(−Si)3 bonding configuration where Si has only one N neighbor is replaced by a H atom. The H incorporation was also evoked to be responsible for the LO band blueshift in SiN x :H [24, 27, 33, 39]. However, our spectra in Figure 5 demonstrate that these two blueshifts are not necessarily linked to H. Besides, similar blueshifts of the TO band [15, 35] and of the LO band  have also been reported in O- and H-free SiN x thin films while the Si content was decreased. As a consequence, these two blueshifts are partly or completely due to some change of the [N]/[Si] ratio in the case of SiN x :H or pure SiN x , respectively. The change in the positions of the TO and the LO modes of Si-N absorption bands are due to some modifications intrinsic to the Si-N binding configuration. In their calculation, Hasegawa et al.  have predicted that the blueshift of the TO mode is linked to the decrease of the Si-N bond length which is caused by a compositional change of SiN x [25, 41]. In addition to this, some stress in the films induced by the Si incorporation may also contribute to such shifts . Moreover, one can assume that the TO-LO coupling of the Si-N asymmetric stretching modes is induced by the disorder in the material in the same manner as that established in Si oxide [42, 43]. Consequently, the increase of the LO band intensity is a signature of the ordering of the films while the Si content is decreased.
where νTO(x) and νLO(x) are the TO and the LO band positions, respectively, and νTO(4/3) and νLO(4/3) are the TO and the LO band positions calculated for x = 4/3, which correspond to the stoichiometric condition, respectively. We found νTO(4/3) = 840 cm−1 which is interestingly the value attributed to the Si-N stretching vibration of an isolated nitrogen in a N-Si3 network [33, 44] and νLO(4/3) = 1197 cm−1. These relations can be used to estimate the composition of as-deposited Si-rich SiN x films in the same way as the empirical one concerning Si-rich silicon oxide .
In order to further investigate the microstructure of the 1100°C-annealed films, HRTEM observations have been performed on several thin films with various n > 2.5. Figure 9b shows the diffraction pattern of one film with n = 2.89. One can observe three quasi-continuous rings corresponding to various orientations of c-Si because of the presence of randomly oriented crystalline Si-np. These numerous crystalline Si-np can be easily distinguished from the host matrix (Figure 9a) because of the lattice fringes of c-Si. They are rather small with an average size of about 6.0 ± 0.5 nm (Figure 9c).
The extensive investigation of the microstructure of SiN x films versus the composition and the annealing treatments enables us to discuss on the PL origin considering that the films do not contain any oxygen and hydrogen. We show that neither defect states within the bandgap nor band tail states could account for all the aspects of the PL. Although we could form crystalline Si-np, we show that the radiative emission is not originating from confined states in crystalline Si-np but could be related to small amorphous Si-np.
Optically active defect states within the bandgap of amorphous SiN x could play a role in the radiative recombination of SiN x as reported by several authors [18, 53]. This interpretation is based on the wide PL spectra that contained distinct PL peaks with several energy levels that corresponded to the calculated values of various defect states found by Robertson [54, 55]. Similar spectra were observed in the 1.75 to 3.1 eV spectral range by Ko et al.  who noticed a redshift of the PL with decreasing Si content. This evolution is in contrast to that of our PL spectra which, moreover, do not contain any distinct PL peaks attributable to distinct defect state levels. As a consequence, we believe that the origin of the PL of our SiN x samples cannot be ascribed to defect states localized within the bandgap.
Let us consider the optical transition between photogenerated carriers localized in the band tail of the material in accordance with the static disorder model . In this model, the carrier distribution in the exponential band tail density of states accounts for the PL band position and the PL shape of SiN x :H . An increase of the width of the localized states results in a blueshift and an increase of the width of the PL band. On one hand, many groups [13, 16] explained that the increase of the structural disorder caused by the nitrogen alloying in Si-rich SiN x :H with a very high Si content (SiNx<0.6) accounts for the widening of the band tail states and then for the PL behavior. On the other hand, many groups [2–4] explained that the increase of the structural disorder induced by the incorporation of more nitrogen in N-rich SiNx>1.33:H films accounts for the widening of the band tails and the PL properties. The increase of disorder in N-rich SiNx>1.33 films with increasing N content is consistent with the FTIR spectra of Huang et al.  which show a drop of the LO band intensity.
Figure 5b showed that the situation is inversed in our Si-rich SiN x films with a low Si excess content since the disorder manifestly increases with the Si incorporation. Therefore, the redshift of the PL band (Figure 12) cannot be explained by the tail-to-tail radiative recombination which would anyway be in contradiction with the widening of the PL band (inset of Figure 12). As a consequence, unlike Si-rich SiN x :H films with a very high Si content (SiNx<0.6) [13, 16], we believe that the static disorder model cannot account for the PL properties of H-free Si-rich SiN x films containing a low Si content (SiNx>0.85). Besides, it has been shown that the hydrogen concentration plays an important role in the PL properties (intensity and peak position) of hydrogenated films .
Crystalline Si-np were detected by Raman, XRD, and HRTEM in numerous SiN x films annealed at 1100°C that had a high n > 2.5 (SiNx<0.8). Furthermore, we have demonstrated in Figure 8 that the progressive redshift of the crystalline Raman peak while n decreased is due to the decrease of the crystalline Si-np average size. The average sizes in the films with n ranging from 2.53 to 2.89 are between 2.5 and 6 nm, respectively. Theses sizes are theoretically small enough to show PL from excitons confined in crystalline Si-np according to the QCE model . This model was proposed to explain the size dependence of the PL peak position that was noticed in oxide systems [1, 59]. This size effect was evidenced in free crystalline Si-np surrounded by a thin Si oxide shell , which however slightly differ from that generally observed while the crystalline Si-np are embedded in a Si oxide host medium [59, 61]. In the case of Si nitride as embedding matrix, several authors suggested that the PL could emanate from confined states in crystalline Si-np, which were present in the materials as attested by HRTEM observations, mainly because of a perceivable size effect on the PL [10–14]. Although our measurements (Figure 12) also show that the PL peak shifted to lower energies with increasing Si content, which is consistent with the QCE model, crystalline Si-np cannot be responsible for the radiative emission for two reasons: (1) Although small (2.5 to 6 nm) Si nanocrystals could be formed in films with n > 2.5 during annealing at 1100°C, we could not detect any PL. PL was detected only for smaller refractive indexes (n < 2.4). Besides, we demonstrated in Figure 7b and Figure 10 that this temperature is necessary to crystallize the excess of Si. Furthermore, (2) the PL of luminescent SiN x films (i.e., with n < 2.4) was quenched while we could form crystalline Si-np by another annealing method using an intense laser irradiation (Figure 14).
Although we have demonstrated that crystalline Si-np are not valid to explain the PL, let us consider Si-np with an amorphous phase as proposed by several authors [5–8]. The annealing temperature dependence of the FTIR spectra of one luminescent SiN x film (n = 2.22) shown in Figure 6 suggests that a phase separation between Si-np and the Si nitride host media occurred during the annealing. The two Raman bands of a-Si at 150 and 485 cm−1 shown in Figure 7 indicate that luminescent films (i.e., with n < 2.4) could contain amorphous Si-np. Besides, the Raman spectra would then show that the density of amorphous Si-np increased with increasing annealing temperature. This explains the absence of PL in the as-deposited samples and why the highest integrated PL intensity (Figure 13) was found at 900°C and not at 1100°C when crystalline Si-np could form. The redshift of the PL bands with increasing Si content (Figure 12) would then be due to a size effect. Also, the increase of the PL band width would then result from the widening of the size distribution as experimentally observed in Si oxide matrices [59, 61]. Then, we have imaged a 1,000°C-annealed SiO x /SiN x multilayer by energy-filtered transmission electron microscopy enabling to distinguish small amorphous Si-np from the host media because of the high contrast of this technique. Because of PL interest, the refractive index of the SiN x sublayer was set between 2.1 and 2.3. We could distinctly observe amorphous Si-np in the 3.5-nm-thick SiO x sublayers, but no particles were perceivable in the 5-nm-thick SiN x sublayers . Si-np could be however very small, below the EFTEM detection threshold of about 1 to 2 nm, and then constituted less than 1000 of Si atoms. Besides, such an amorphous Si-np size seems possible compared to the average size of 2.5 nm of crystalline Si-np detected by Raman spectroscopy in SiN x with n = 2.53. Consequently, the origin of the PL would be related to small amorphous Si-np, and the recombination would originate either from confined states in the Si-np and/or from defect states at the interface between the Si-np and the Si nitride medium .
We have produced pure amorphous Si-rich SiNx < 1.33 thin films by magnetron sputtering with various Si contents using two deposition methods, namely the N2-reactive sputtering of a Si target and the co-sputtering of Si and Si3N4 targets. The dependence of the only Si content on the microstructure and on the optical properties was studied. The two synthesis methods are equivalent since no systematic change could be discerned in the structural and the optical analyses. Besides, no trace of O atoms was detected by RBS and by FTIR, and no H bonded to Si or N could be detected by FTIR. We could then establish an empirical relation between the [N]/[Si] ratio and n based on the random bonding model on pure SiN x which manifestly differs from previous relations that concerned SiN x :H because of the H incorporation induced by the chemical deposition techniques.
Because of the absence of Si-H, N-H, and Si-O absorption bands, we could highlight the Berreman effect on the FTIR spectra of SiN x by the normal incidence and an oblique illumination. The TO-LO pair modes of the two Si-N stretching absorption bands could be unambiguously assigned. A redshift of the two modes and a drop of the LO band intensity were observed while the Si content increased, which indicates that incorporation of more Si generates more disorder in the films. Moreover, a significant blueshift of the two modes with increasing annealing temperature was noticed which may be explained by a phase separation between Si-np and the Si nitride medium. At the same time, the LO band intensity increased indicating a rearrangement of the Si nitride network towards less disorder.
The effect of the annealing temperature on the Raman spectra has been investigated on films with n < 2.5 (SiNx>0.9). The Raman spectra indicate that small amorphous Si-np could be formed during the annealing and that their density increased with the annealing temperature. For higher n (n > 2.5, SiNx<0.8), Raman spectra, as well as XRD patterns, demonstrated that crystalline Si-np are formed upon annealing at 1100°C. Moreover, QCE on the optical phonon in crystalline Si-np embedded in Si nitride was observed. It matches with previous theoretical models concerning Si nanocrystals in Si oxide systems. The average size measured by HRTEM increased from 2.5 to 6 nm with increasing n.
Only SiN x films with n ranging from 2.01 to 2.34 (SiNx>0.9) exhibit visible PL. The PL bands redshifted and widened while n was increased. The tail to tail recombination cannot account for these PL properties since the FTIR spectra showed that the disorder increased with increasing n which would result in a blueshift and a widening of the PL bands. The PL could be then due to a QCE. The annealing temperature dependence of the PL intensity is consistent with the formation of Si-np. Nevertheless, the PL is not related to crystalline Si-np since they have not been detected in luminescent films by XRD and Raman measurements. As an additional proof, the PL quenched while Si crystalline Si-np could be formed by an intense laser irradiation. As a consequence, we believe that the PL is actually related to small amorphous Si-np and/or defect states that could be located at the interface between Si-np and the Si nitride host medium.
The authors acknowledge the French Agence Nationale de la Recherche, which supported this work through the Nanoscience and Nanotechnology Program (DAPHNÉS project ANR-08-NANO-005).
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