Defect Emission and Optical Gain in SiCxOy:H Films - ACS Applied

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Defect emission and optical gain in SiCxOy:H films Zhenxu Lin, Hongfei Li, Rui Huang, Yi Zhang, Jie Song, Hongliang Li, Yanqing Guo, Chao Song, and John Robertson ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b06118 • Publication Date (Web): 26 Jun 2017 Downloaded from http://pubs.acs.org on June 29, 2017

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ACS Applied Materials & Interfaces

Defect emission and optical gain in SiCxOy:H films Zhenxu Lina, †, Hongfei Lib, †, Rui Huanga, *, Yi Zhanga, Jie Songa, Hongliang Lia, Yanqing Guoa, Chao Songa, and John Robertsonb a.School of Materials Science and Engineering, Hanshan Normal University, Chaozhou, Guangdong 521041, China b.Department of Engineering, University of Cambridge, Cambridge, UK Keyword: Silicon oxycarbide, Plasma enhanced chemical vapor deposition, Defect, Photoluminescence, Optical gain

ABSTRACT : Luminescent SiCxOy:H films，which are fabricated at different CH4 flow rates using plasma-enhanced chemical vapor deposition (PECVD) technique, exhibit strong photoluminescence (PL) with tuning from near-infrared to orange regions. The PL features an excitation wavelength-independent recombination dynamics. The silicon dangling bond (DB) defects identified by electron paramagnetic resonance spectra are found to play a key role in the PL behavior. The first-principles calculation shows that the Si DB defects introduce a mid-gap state in the band gap, which is in good agreement with the PL energy. Moreover, the band gap of a-SiCxOy:H is found to be mainly determined by Si and C atoms. Thus, the strong light emission is believed to result from the recombination of excited electrons and holes in Si DB defects, while the tunable light emission of the films is attributed to the substitution of stronger Si–C

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bonds for weak Si–Si bonds. It is also found that the light emission intensity shows a superlinear dependence on the pump intensity. Interestingly, the film exhibits a net optical gain under ultraviolet excitation. The gain coefficient is 53.5cm-1 under a pumping power density of 553mW cm-2. The present results demonstrate that the SiCxOy system can be a very competitive candidate in the applications of photonics and optoelectronics.

1. Introduction The development of Si-based luminescent materials is motivated by the demand for cheap and complementary metal-oxide semiconductor (CMOS) light sources that are compatible with onchip optical interconnect. Progress has been recently made in understanding and optimizing light emission from Si-based materials, such as SiOx, SiNx, and SiNxOy. 1-13 A breakthrough in this field is the realization of optical gain and light amplification from SiOx-based material.14-16 Owing to the difficulty of electrical injection in SiOx,2 silicon dioxide is commonly replaced with silicon nitride as a promising luminescent material for efficient electrical excitation.17-21 Compared with silicon nitride, however, silicon oxynitride more effectively balances carrier injections in light-emitting devices, thus remarkably improving the probability of carrier recombination.22,23 Extensive studies have recently focused on dielectric silicon oxycarbide (SiCxOy) due to its efficient light emission and applicability in Si chips. SiCxOy, which is widely used as an interlayer dielectric to reduce parasitic capacitance and resistance capacitance (RC) delay in Si chips, generally features strong white light emission.24-26 Moreover, a SiCxOy host matrix provides a remarkably higher solid solubility for rare earths, such as Eu and Er, compared with SiO2.27,28 In addition, the SiCxOy matrix can also be used as donor that contributes to the light emission of Eu2+ ions via energy transfer.29 So far, defect states, such as C-related oxygen


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vacancies (C-NOVs), Si-NOVs, and Si-related oxygen deficiency centers, have been proposed to demonstrate the PL mechanism in SiCxOy.30-32 Nikas et al recently reported that white light emission results from the band tail states related to the Si–O–C and /or Si–C bonds.33 Although considerable effort has been devoted to understanding PL characters in SiCxOy, the origin of PL remains contradictory because of complicated Si phase structures, such as Si, SiC, and SiOx, which coexist in SiCxOy. In addition, orange to near-infrared emitters that are based on silicon oxycarbide have not yet been achieved. In particular, the optical gain in a-SiCxOy films has never been studied. In this work, we report on strong orange to near-infrared switching PL and net optical gain in SiCxOy:H film. The silicon dangling bond (DB) defects identified by electron paramagnetic resonance spectra are found to play a key role in the PL behavior. The first-principles calculation shows that the Si DB defects introduce a mid-gap state. Moreover, the band gap of a-SiCxOy:H is found to be mainly determined by Si and C atoms. The origin of PL is discussed and a threelevel-luminescence model is put forward to illustrate the PL and optical gain in a-SiCxOy films.

2. Samples and experimental details Amorphous silicon oxycarbide films were fabricated on quartz and Si wafers substrates via very high frequency-PECVD technique. SiH4, O2, and CH4 were used as the reactant gas sources. The flow rates of SiH4 and O2 were maintained at 3.5 and 1.2 sccm, respectively, whereas the flow rates of CH4 (Fr(CH4)) are changed from 5 to 20 sccm. The as-deposited films were designated as Sx(x = 1, 2, 3, 4) with the Fr(CH4) of 5, 10, 15, and 20 sccm. During deposition, the radio frequency power, substrate temperature, and pressure were 30 W, 250 °C, and 20 Pa,


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respectively. The as-deposited films were also dehydrogenated at 400 °C for 1 h, and then annealed at 500 °C for 1 h under a nitrogen atmosphere. PL spectra were measured by Jobin Yvon fluorolog-3 spectrophotometer. The band gaps of the films were evaluated using Shimadzu UV-3600 spectrophotometer. The chemical compositions of Si, C, and O were determined by XPS. The microstructures of the films were characterized by high-resolution transmission electron microscopy (HRTEM). The bonding configurations of the films were recorded using Fourier Transform Infrared Spectroscopy (FTIR). The paramagnetic defects in the films were analyzed with EPR measurements. The SiNx/a-SiCxOy/SiO2 strip-loading waveguide (30 µm×1 cm) was fabricated on quartz. Optical gain was measured by the standard VSL technique. The sX hybrid functional was also used to calculate the band properties of SiCO systems, as well as the involved point defect. The DFT simulations were performed using the CASTEP plane wave pseudopotential code. The norm conserving pseudopotential was used with the cutoff energy of 750 eV in all calculations. The amorphous SiCxOy structures were calculated by the GGA exchange correlational function with the Gamma point scheme due to a large supercell. The screened exchange hybrid functional was applied in all electronic property calculations to correct the well-known bad gap error and provide the correct electron localization distribution. The amorphous SiCO models were generated by the first principles molecular dynamics (MD) method according the following steps: first, the system with a given ratio of Si, C, and O atoms underwent a high-temperature melting process at 2500 K for 10ps. Second, the system was then gradually cooled down to 300 K at the quenching rate of 100K/ps. Third, a rough geometry optimization was applied and H atoms were included to passivate the remaining dangling bonds (DB) on Si and C atoms. Finally, an accurate geometry optimization was performed until the residual force on each atom was lower than 0.01eV/Å.


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3. Results and discussion Film microstructure was analyzed with HRTEM. Figure 1(a) presents the cross-section HRTEM image of S2. S2 does not show any sharply contrasting structures, which clearly indicate that no Si or SiC nanostructures exist in the film and that the film is amorphous. The Si 2p core level spectra of the films (Figure 1(b)), which indicate the ionic states of Si atoms, were used to analyze the phase structures in the films. The binding energy of the Si 2p peak is 101.5 eV for S1, which was grown at a CH4 flow rate of 5 sccm. The binding energy gradually increases to 102.1 eV as the Fr(CH4) increases to 20 sccm. The binding energy of the Si 2p peak for all the films is higher than that of SiC (100.8 eV) and lower than that of SiO2 (103.2 eV).34 These results indicate that the SiCxOy phase is the dominant phase in the films. Figure 2 shows the dependence of the optical band gap (Eopt) of the films on the Fr(CH4). Eopt is calculated from the Tauc plot equation (αhν)1/2=B1/2(hν- Eopt), where hν is the photon energy and α is the absorption coefficient.22 Eopt increases from 2.6 to 2.9 eV as the Fr(CH4) increases from 5 to 20 sccm. This behavior is a result of the increased C content in the films as revealed in Figure 2, similar to that observed in a-SiCx.35 Figure 3(a) shows the room-temperature PL spectra of the SiCxOy film grown at various Fr(CH4). The PL emission band gradually blueshifts from 735 nm to 630 nm as the Fr(CH4) increases from 5 to 20 sccm. Moreover, PL intensity is enhanced. Increasing the Fr(CH4) to 20 sccm significantly increases the integrated PL intensity by more than two orders of magnitude compared with the PL intensity of the film grown at the Fr(CH4) of 5 sccm. Interestingly, the tunable light emissions are visible to the naked eye even at an excitation wavelength of 325 nm from a Xe lamp (Figure 3 (b)). The tunable light emission from a–SiCxOy can be controlled by modulating the Fr(CH4). Figure 3(c) shows that that the PL peak positions for all the films are independent of excitation wavelengths. Meanwhile, the


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FWHM of the PL spectra for all the films almost do not widen as the excitation wavelength decreases from 375 nm to 275 nm (Figure 3(d)). These phenomena are completely different from that observed PL that is governed by the band tail state recombination mechanism.36 In that mechanism, the FWHM of PL gradually widens with the decrease of the excitation wavelength. These results indicate that the PL in our case may result from defect-related luminescence centers. The EPR spectra (Figure 4) of the films were examined to gain more insight on the origin of PL from a–SiCxOy films. The g value of the EPR signal from the films can be calculated according to gµBH=hv, where µB, H, h and v is the Bohr magnetron, the spin Hamiltonian, Planck’s constant and the X band frequency, respectively.37 The spectra for all the films exhibit an EPR signal with the same g value of 2.0047 and the same line width △Hpp of approximately 10 G, which is assigned to unpaired silicon DB defects.38 These characteristics confirm the existence of Si DB defects in the films. As the Fr(CH4) increases from 5 to 20 sccm, the integrated EPR signal intensity, which is proportional to the calculated concentration Ns of the Si DB defects, gradually increases from 1.17×1019 cm-3 to 2.19×1019 cm-3. Obviously, the increased concentration of Si DB defects plays an important role in the increase of PL shown in Figure 3(a). To further understand the origin of PL, S2 was annealed at 500 ºC for 1 h in N2 atmosphere. From Figure 5, one can see that the PL spectrum of the S2 annealed at 500 ºC in N2 is nearly doubled with that of S2. However, the PL positions of the two samples do not noticeably change. The Figure 5 inset shows that thermal annealing at 500 ºC in N2 also remarkably enhanced the EPR signal with the same g value of 2.0047 and the same line width △Hpp of approximately 10 G as that of S2. Based on the EPR results, the concentration of unpaired Si DB defects of the annealed S2 in N2 is calculated to be increased by two-fold compared with that of S2. This


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change is similar to that of PL intensity. Thus, the enhanced PL of the annealed film results from the increased concentration of Si DB defects. The independence of PL peak positions with the measured temperatures shown in Figure S1 (Shown in the supporting information) further confirms the origin of PL from the Si DB defects. It is well known that the annealing process in N2 ambient induces the diffusion of nitrogen atoms inside the film, and then the nitrogen atom will replace the Si in the network to form Si-N bonds and leave a new dangling bond, which may explain the increased concentration of Si DB defects after annealing treatment. However, the nitridation of SiCxOy film is unlikely to occur at 500 ºC because this annealing temperature is not enough to decompose the N2 gaseous molecules. It is considered that Si DB defects maybe come from Si-H bonds since Si-H bonds would be broken at the annealing temperatures larger than 350 ºC.39 From Figure 5, one can see that hydrogen annealing treatment can also improve the PL intensity. To demonstrate the evolution of bonding configuration within the films, FTIR absorption spectra of the S2 and the annealed S2 were examined, as shown in Figure 6. The S2 FTIR spectrum shows the following absorption bands: two strong absorption bands at 800 cm-1 and 1020 cm-1, which are attributed to the C–Si and Si–O–Si stretching vibration modes, respectively.40,41 In addition to the above two bands, a weak band at 1260 cm-1 is ascribed to the Si–CH3 stretching vibration mode.42 The 2950 cm-1 bands result from the C–H stretching vibration band.42 The 2150 cm-1 band is due to the stretching mode of Si–H bonds.43 The evident feature in the FTIR spectra (Figure 6) is that the signal of Si–H stretching vibration band decreases after annealing at 500 ºC in N2 and/or H2 atmosphere. This result implies that the Si–H bonds are broken and that numerous silicon dangling bonds form during the annealing process. Therefore, the increased concentration of Si DB defects can be attributed to broken Si–H bonds,


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which significantly increased after annealing S2 at 500 ºC. It is worth noticing that the signal of C–H bond stretching mode also decreases after annealing at 500 ºC in N2 and/or H2 atmosphere. Furthermore, the ratio of the C-H bond intensity to the Si-H bond intensity of the film annealed under the H2 atmosphere is higher than that of the film annealed under the N2 atmosphere. This indicates that hydrogen prefer to passivate the C DBs rather than the Si DBs under hydrogen annealing treatments. This is because the C-H bond (414KJ/mol) is stronger than Si-H bond (377KJ/mol). Thus, combining the PL with the FTIR spectra, C related defects are suggested mainly as nonradiative recombination centers in the films. After the hydrogen annealing treatments, C related defects are partially passivated by hydrogen. Consequently, the PL intensity of the films annealed under the H2 atmosphere is higher than that of the films annealed under the N2 atmosphere. It is also found that the intensity of Si-C stretching vibration band is almost unchanged before and after annealing at 500 ºC, which is completely different from that observed in the Ref. [35]. It indicates that the annealing at 500 ºC cannot cause the remaining Si and C atoms to form Si-C bonds. To gain more insight on the optical properties of a–SiCxOy films, the band characters of SiCO systems and involved point defect were calculated with the sX hybrid function. As shown in Figure 7(a), the defect-free atomic structure of the amorphous SiCO model was constructed according to the procedures described in the experimental section. The component ratio in the model is close to the component ratio of S4 as measured by the XPS (Table 1). In the model, each atom binds in the correct coordination and no 4-membered ring appears. The density of states (DOS) of the model, as calculated by the sX exchange correlation functional, is shown in Figure 7(b). The calculated band gap is 2.8 eV, which is close to the Eopt (2.9 eV) estimated by the Tauc plot equation (Figure 2). The DOS figure shows a clean band gap for this defect-free


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model. Moreover, the valence band edge and conduction band edge are contributed by Si and C atoms. Therefore, the band gap of a–SiCxOy is mainly determined by the interaction among Si and C atoms. Thus, the widening of the band gap with increasing Fr(CH4) can be ascribed from the replacement of weak Si–Si bonds by stronger Si–C bonds, which had been demonstrated in the a-SiCx.44 Figure 8(a) shows the electronic orbital of the Si dangling bond (DB) defect in the SiCxOy system, as given by the sX functional. The defect orbital is highly localized on two 3-fold Si atoms. The partial DOS (Figure 8(b)) indicates that Si DB defects introduce a mid-gap state at 1.85 eV above the valence band minimum. This result is in good agreement with the PL peak energy observed in S4. Accordingly, the strong light emission of the films can be ascribed to the recombination of excited electrons and holes in Si DB defects. The tunable light emission with the increasing Fr(CH4) can be attributed to the substitution of stronger Si–C bonds for weak Si– Si bonds. It is well known that the C sp2 clusters can produce light emission. However, no any D peak or G peak can be found in the Raman spectra (not shown) of the films. This indicates no any the C sp2 clusters existing in the films, thus excluding the possibility of the light emission from the the C sp2 clusters. To measure the optical gain coefficient, the geometry structure of SiNx/a-SiCxOy (S4)/SiO2 strip-loaded waveguide were fabricated, and the light emission is detected by VSL technique, as shown in the inset of Figure 9. Figure 9 shows the integrated PL intensity of S4 as a function of pump power. The integrated PL intensity is found to rise up with the pumping laser power in a superliner increase manner. This strongly indicates the existence of stimulated emission in the waveguide.45 Figure 10 shows the amplified spontaneous emission (ASE) intensity IASE versus the pumping length from the strip-loaded SiNx/a-SiCxOy (S4)/SiO2 waveguide measured at a


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wavelength of 325nm. One can see that the IASE tends to increase with increasing the pumping length, as shown in the inset of Figure 10. For the shorter pumping length (

Defect Emission and Optical Gain in SiCxOy:H Films - ACS Applied

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