Abstract
Silicon oxycarbide (SiCxOy) has been extensively investigated due to its wide use in the Si semiconductor industry in applications that include low-k dielectrics, passivation layers, and etch-stop layers. Furthermore, SiCxOy research has been exploring its prospective use in numerous other technological usages, such as lighting, energy, and biological applications. The latter include white light-emitting materials, hydrogen storage materials, gas sensors, anode materials for lithium batteries, and biomedical devices. SiCxOy materials can intensively luminescence in a broad emission spectral range that spans the ultraviolet, the visible, and even the near-infrared spectrum, when doped with erbium. Herein, we present pertinent results on the material behaviors from chemically synthesized SiCxOy thin films and nanowires. Moreover, their light-emitting properties and underlying mechanisms for light emission are explored in conjunction with data from their thin film counterparts, which are also employed as baseline comparison metric. We further highlight major challenges and promises of such materials.
Keywords
- silicon oxycarbide
- SiCxOy
- Si‐O‐C
- Si‐C‐O
- nanowires
- thin films
- luminescence
- CVD
- e-beam lithography
- structural properties
- optical properties
- band tails
- disorder
1. Introduction
In silicon (Si) complementary metal-oxide-semiconductor (CMOS) chip technology, silicon oxycarbide (SiCxOy) materials have been extensively employed to serve multiple purposes. For example, SiCxOy materials have been the focus of extensive study due to their applicability as low-k dielectrics, passivation layers, and etch-stop layers to name a few [1, 2]. SiCxOy materials have also been the focus of studies due to their potential application in a plethora of other technological applications (e.g. light emission, energy, and bioapplications). In particular, they have been proposed as candidates for white light-emitting materials [3–5], filters, porous adsorbents, and catalytic supports [6, 7], as hydrogen storage materials [8], gas sensors [9], negative electrode materials for lithium batteries [10, 11], and in biomedical devices [12]. Additionally, it has been also shown that SiCxOy can be utilized as a host material to optically active impurities (rare earth ions). To this end, europium (Eu2+)-doped SiCxOy thin films synthetized by RF magnetron sputtering [13] as well as erbium (Er3+)-doped SiCxOy thin films for near-infrared (IR) emission at the commercially useful telecommunication wavelength of 1540 nm have been recently reported [14–16].
The fabrication of luminescent Si-based nanostructured materials for light emission applications is highly desirable, similarly to how Si-based nanophotonics has undergone great advancements in recent years [17, 18]. Due to the seamless integration of Si-based materials with process protocols and technologies developed for semiconductor CMOS technology, manufacturing costs and process complexity could be also reduced. Furthermore, the extreme down-scaling methods achieved by CMOS technology offer the opportunity to study new compelling properties owing to possible confinement effects in one (1D) or two (2D) dimensions (e.g., the reduction of exciton-phonon interactions, the increase of extraction efficiency of spontaneous emission, and suppression of Auger recombination) [19, 20]. Therefore, the functionality of such nanostructured materials and their devices can be employed in a ubiquitous way in light emission applications [21]. To this extent, SiCxOy nanowires (NWs) have been recently shown to exhibit strong room-temperature visible luminescence [22–24].
In parallel, the identification and elimination of potential obstacles that could deteriorate the luminescence efficiency of such materials (e.g., temperature and excitation power density) needs to be taken into consideration. Indeed, the luminescence efficiency is known to be influenced by environmental fluctuations in temperature and pump-power-density changes [25, 26]. These fluctuations can become critical in luminescence applications such as light-emitting diodes (LEDs). In these applications, the operating temperature and power density can reach respectively ~150°C and ~200 W/cm2, thus greatly influencing the light output and color chromaticity.
Silicon oxycarbide films have been grown predominantly through low-temperature plasma-enhanced chemical vapor deposition (PECVD) using an array of silane-based precursors for the needs of semiconductor industry [27–30]. However, the incorporation of source precursor fragments and decomposition byproducts in the resulting films has led to the observation of enhanced stress levels and increased defect density, both of which have detrimental effects on the optical performance of resulting materials and device structures. Recently, Lin et al. prepared amorphous SiCxOy using a very high-frequency plasma-enhanced chemical vapor deposition (VHF-PECVD) technique [31]. The resulting films exhibited intense room-temperature blue luminescence, characteristic of Si-related neutral oxygen vacancy defect centers. Ryan et al. and Vasin et al. showed that a wide variety of SiCxOy with a continuous range of compositions could be produced by reactive RF-magnetron sputtering [32, 33]. Karakuscu et al. and Abbass et al. have reported sol-gel-prepared SiCxOy thin films [34, 35], while Vasin et al. have reported a-SiCxOy:H thin films growth by low-temperature oxidation of carbon-rich a-SixCy:H thin films [36].
Nevertheless, it is desirable to identify alternative deposition methods, which can inhibit or minimize processing induced structural and/or compositional damage to SiCxOy materials due to phase separation owing to the non-stoichiometric composition. Herein, key findings are summarized pertaining to the development of a thermal chemical vapor deposition (TCVD) strategy for the growth of SiCxOy thin films approximating the SiCxO2(1−x) (0 < x < 1) stoichiometry for light emission applications [37, 38]. These studies led to the identification of a deposition process window for the growth of SiCxOy thin films with strong room-temperature light emission [3]. Additionally, we present the findings pertaining to room-temperature visible photoluminescence (PL) from SiCxOy sub-100-nm nanowire materials fabricated by electron beam lithography (EBL) and reactive ion etching (RIE). These metal-free non-toxic Si-based nanostructured materials may offer an alternative and environmentally friendly pathway toward efficient visible light-emitting materials and devices.
2. Synthesis and fabrication of silicon oxycarbide thin films and nanostructures
2.1. SiCxOy materials grown by thermal chemical vapor deposition: composition control
SiCxOy thin films and their nanostructures are deposited on Si(100) or SiO2 substrates in a hot-wall quartz tube reactor by thermal CVD at 800°C. A single source oligomer (2,4,6-trimethyl-2,4,6-trisila-heptane (C7H22Si3)) is utilized as the source precursor for silicon and carbon atoms and ultra-high purity oxygen (O2) and argon (Ar)/nitrogen (N2) are also employed as co-reactant and dilution gases. The composition of the resulting amorphous SiCxOy≤1.73 (0.11 < x < 0.65) thin films is regulated by modifying the oxygen flow rate and, hence, the oxygen partial pressure in the reaction zone. The films are deposited onto three types of substrates: single-crystal silicon (c-Si) (100), for composition, structural and optical analysis; double-polished intrinsic Si and high-quality UV transparent fused silica, for infrared and optical studies. A subset of the as-deposited SiCxOy films was annealed for 1 h in different ambient (e.g., O2, Ar, N2, or forming gas (FG, 5% H2 and 95% N2)) at temperatures in the range from 500 to 1100°C. Detailed description of the deposition and annealing processes can be found elsewhere [37, 38]. The resulting samples were subsequently employed in a comparative analysis of as-deposited and post-annealed films to determine the effects of thermal treatment on film structural and optical properties as well as their photoluminescence performance.
Rutherford backscattering spectroscopy (RBS) and X-ray photoelectron spectroscopy (XPS) were employed to determine the Si, C, and O content in the SiCxOy materials and they were separated in three different classes, defined by their composition: SiC-like ([O] < 5 at.%), Si‐C‐O, and SiO2-like ([C] < 5 at.%). The compositional evolution of the SiCxOy thin films was plotted in the Si‐C‐O ternary diagram of Figure 1a along with the SiCxO2(1−x) (0 < x < 1) line, with the latter representing stoichiometric silicon oxycarbide without any excess of carbon. The upper and lower limits in SiCxO2(1−x) correspond to SiC and SiO2, for, respectively, x = 1 and 0 [39, 40].
As is shown in Figure 1a, the composition of the thermal CVD-grown SiCxOy was properly tailored to closely follow the pure stoichiometric oxycarbide formula over the whole range of the synthesis process parameters. The SiCxO2(1−x) behavior observed in the samples produced herein suggests the substitution of two divalent oxygen atoms by one tetravalent carbon atom as the C concentration in the SiCxOy materials increases. The latter will be further discussed in Section 3.1. Furthermore, the observed small deviation from the stoichiometric SiCxO2(1−x) trend suggests that there is much less excess of C compared to non-stoichiometric compositions reported for SiCxOy with SiC/SiO2 phases and free carbon, as it is shown in Figure 1 of reference [39].
Due to the presence of methyl groups in the CVD reactants, it is expected to have hydrogen atoms in the grown silicon oxycarbide thin films whose concentration cannot be quantified by the RBS or XPS techniques. Instead, the nuclear reaction analysis (NRA) was conducted by using the 15N + 1H → 12C + 4He + γ-ray resonant nuclear reaction at 6.835 MeV [41]. Figure 1b shows the NRA-derived hydrogen concentration measurements of the silicon oxycarbide films as a function of post-deposition annealing temperature
The mass density
where
The density of the silicon oxycarbide films was observed to increase with higher annealing temperature, due at least partly to the reduction in hydrogen concentration, as shown in Figure 1c. More specifically, the densities of the as-deposited films are 2.8, 2.2, and 2.1 for, respectively, SiC-like, Si‐C‐O, and SiO2-like films. After annealing at 1100°C, the densities were measured to be significantly higher, with values of 3.3, 2.6, and 2.4 g/cm3 for, respectively, SiC-like, Si‐C‐O, and SiO2-like films.
2.2. Nanofabrication of silicon oxycarbide nanostructures
SiCxOy nanostructured systems were fabricated by electron beam lithography and reactive ion etching, namely periodically ordered sub-100-nm nanowire arrays. A representation of the baseline nanofabrication scheme of the SiCxOy NW arrays is schematically depicted in Figure 2. Following the synthesis of SiCxOy thin films, negative hydrogen silsesquioxane (HSQ) (6% HSQ in methyl-isobutyl-ketone solvent) resist is spun onto SiCxOy wafer pieces—deposited on Si (100)—at 1000 rpm, followed by a bake procedure for 4 min at 80°C. Line patterns are exposed using electron beam lithography and then the resulting wafer piece is developed in a chemical solution bath (low concentration (2.38%) of tetramethylammonium hydroxide (TMAH) developer), yielding 2D-nanowire HSQ patterns. The HSQ patterns then underwent a hardening annealing process in Ar/N2 ambient at 500°C, followed by a fluorine-based (e.g., combination of CHF3 and CF4 gases) anisotropic RIE to transfer the pattern into the SiCxOy thin films. The HSQ residue is then removed by wet etching in buffered hydrofluoric (BHF) acid, resulting in periodically well-defined NWs [22].
Pertaining to the NW fabrication, certain samples underwent a sidewall image transfer (SIT) process which was performed by conformal deposition of a thin silicon carbide (SiC) hard mask (sidewall layer) on the patterned photoresist followed by anisotropic RIE [23]. This allowed for a significant reduction of the critical dimensions as, during the SIT process, the NW width is defined by the thickness of the SiC layer rather than the resolution of the lithography step. The SiCxOy NWs underwent different post-fabrication thermal treatments for up to 2 h at annealing temperatures in the range of 50–700°C in Ar, O2, or forming gas (5% H2 and 95% N2) atmospheres.
3. Bonding configuration of silicon oxycarbide materials and nanowires
3.1. Fourier transform infrared spectroscopy (FTIR) characterization of silicon oxycarbide thin films
Figure 3a shows the Fourier transform infrared spectroscopy (FTIR) data of SiCxOy thin films in the 400–1700-cm−1 range. The spectra are characterized by three absorption bands and a shoulder assigned to the Si‐O‐C rocking, Si‐C stretching, and Si‐O transverse and longitudinal-stretching vibration modes, centered at, respectively, ~440, ~800, ~1000, and ~1150 cm−1 [22, 37, 38].
More specifically, the Si‐O‐C vibration mode (~440 cm−1) is caused by the ≡Si‐O‐Si≡ rocking mode due to out-of-plane motion of O in O(3−x)Cx≡Si‐O‐Si≡CyO(3-y) and is unaffected by the Si‐O‐Si-bridging bond-angle variation [42–45]. For samples with C content higher than 20 at.%, the density of ≡Si‐O‐Si≡-bonding groups decreased significantly, as dictated by the SiCxO2(1−x) stoichiometry, suggesting that the backbone-bonding network related to the SiO4 tetrahedral in SiCxOy changed toward SiC-like structures with significant presence of SiC4 tetrahedral. The replacement of O atoms by C atoms with increasing C content in films is reflected in the monotonic increase of the bond area ratio of [Si‐C] and [Si‐O] as clearly depicted in Figure 3b.
Additionally, incorporation of the less electronegative C atoms leads to a reduced Si‐O‐Si bond angle between tetrahedral (see Figure 3c) [30, 44]. As a result, the density of the as-grown films is expected to increase with increasing C content, which is in agreement with the density values shown in Section 2.
Figure 4 displays the FTIR spectra collected in the range from 400 to 2300 cm−1 for the three classes of as-deposited silicon oxycarbide materials and for their annealed counterparts at 900 and 1100°C [37].
Deconvolution of the FTIR spectra reveals several bonding components in the as-deposited SiCxOy material systems [30, 37, 45, 46]. In particular, the deconvolution of the absorption bands in the range of 400–1400 cm−1 for the SiC-like film indicates the presence of a weak C‐H mode at ~530 cm−1, a major Si‐C-stretching absorption mode at ~764 cm−1, and a shoulder assigned to the Si‐O-stretching mode at ~960 cm−1. The hydrogen-related absorption bands (Si‐H and C‐H) are located at ~2100 and 2900 cm−1, respectively. Finally, the absorption observed at ~1846 cm−1 is attributed to a C‐O-stretching mode.
As seen in Figure 4b, the IR spectrum of Si‐C‐O film has three characteristic absorption band regions, originating from the Si‐C and Si‐O functional groups [37]. In comparison to the SiC-like matrix films, a new absorption peak is seen at ~440 cm−1, attributed to the Si‐O‐C vibration mode discussed above. The IR region from 600 to 1300 cm−1 is broader compared to that in the SiC-like films, and its deconvolution shows the presence of four peaks centered at 663, 816, 1002, and 1114 cm−1 attributed to Si‐C‐H, Si‐C stretching, and to the transverse optical (TO) and longitudinal optical (LO) asymmetric Si‐O-stretching modes, respectively. Compared to the SiC-like films, the position of the Si‐C absorption peak shifted from 764 to 816 cm−1, owing to the addition of more electronegative O atoms in the network [47].
The FTIR absorption spectrum of the SiO2-like film is characterized by an intense Si‐O‐Si mode (rocking), centered at ~440 cm−1: a Si‐O mode (bending) located at ~815 cm−1, and an intense Si‐O vibration mode (stretching) at ~1100 cm−1. The hydrogen-related vibration modes for C‐H and O‐H are observed at, respectively, ~2900 cm−1 and in the 3100–3700-cm−1 range.
The findings outlined above describe the evolution of the as-deposited films from silicon carbide-like to silicon dioxide-like films as the amount of C decreases. Regarding the annealed samples up to 700°C, the IR absorption behavior remains similar to the case of their as-deposited counterparts, and it is worth mentioning that changes took place at annealing temperatures beyond 900°C [44−46].
In the case of the SiC-like sample annealed at 900°C (Figure 4a), the Si‐C- and Si‐O-stretching modes show minor changes. However, both hydrogen-related modes appear with reduced intensities, as expected from the NRA results shown in Figure 1b. After the 1100°C annealing, the Si‐C absorption band increased drastically in intensity and its line shape changed from Gaussian to mix of Gaussian and Lorentzian, suggesting the presence of longer-range order (Lorentzian). Additionally, all hydrogen-related vibration modes are no longer present in the films owing to hydrogen desorption.
Similarly, for the 900°C-annealed Si‐C‐O a small intensity increase of the Si‐C-stretching mode was observed, while further annealing at 1100°C led to an overall absorption intensity increase accompanied with a blue shift of the Si‐O stretching (Figure 4b). Finally, the annealing studies on the SiO2-like material (Figure 4c) revealed a significant intensity increase of the Si‐O‐Si-rocking and the Si‐O-stretching modes.
The bond density is directly proportional to the area of the IR band absorption, and can be estimated as in reference [37] using the inverse absorption cross section found in literature (3 × 1019 cm−2 for Si‐C [48], 1.4 × 1020 cm−2 for Si‐H [49], and 1.35 × 1021 cm−2 for C‐H [50, 51]). The Si‐C bond density for the as-deposited SiC-like sample is ~2.2 × 1023 cm−3 while for the as-deposited Si‐C‐O sample is ~5.5 × 1022 cm−3. The dependence of Si‐C bond density with annealing temperature for both SiC-like and Si‐C‐O is presented in Figure 5a. It shows a constant concentration up to 700°C annealing temperature. At higher temperatures, the Si‐C bond concentration increases owing to the densification of the materials and hydrogen desorption, which contributes to the increased availability of Si and C atoms formerly bonded to hydrogen. Indeed, as seen in Figure 5b, the bond density of the Si‐H and C‐H bonds decreases with annealing temperature
As shown in Figure 5c, the total concentration of hydrogen atoms as determined by the NRA measurements is greater than the total content of hydrogen as calculated by FTIR. Each H-related bond corresponds to one H atom; therefore, the H-related bond density corresponds to the H atomic density. The total atomic density is determined by RBS measurements. This finding suggests that some H atoms are incorporated during CVD growth and are not chemically bound to other elements. These non-bonded H atoms may be present in the form of molecular hydrogen formed during the decomposition of the precursor [37].
3.2. FTIR characterization of silicon oxycarbide nanowires
Large SiCxOy NW structures were fabricated on intrinsic Si substrates in order to perform bonding configuration analysis [22]. The normalized absorbance FTIR spectra of both the as-deposited SiC0.34O1.52 thin film and the as-fabricated NWs are shown in Figure 6. It was found that the bonding configuration of the SiC0.34O1.52 NW system was maintained after nanofabrication as the relative intensities of Si‐C and Si‐O absorption bands remained the same with respect to the as-deposited thin film. A slight absorption increase of the Si‐O shoulder at ~1200 cm−1 may be due to surface oxidation induced on the as-synthesized NWs.
3.3. XPS characterization of silicon oxycarbide thin films
The information about the bonding configuration in silicon oxycarbide thin films extracted by the means of FTIR analysis was also independently assessed by XPS studies [37]. The XPS analysis focused on examining the electronic environment (chemical bonding) of the Si 2p, C 1s, and O 1s core energy.
In the case of the as-deposited samples, the Si 2p spectrum of the SiC-like matrix (Figure 7) is composed of a center peak at 100.3 eV assigned to Si‐C bonds and two shoulders centered at 99.2 and 101.2 eV assigned to Si‐H- and Si‐C‐O-type bonding [37, 52]. In the case of the Si‐C‐O material, the Si 2p peak broadened and shifted to higher binding energies. The peak deconvolution showed the presence of three components centered at 100.8, 102.1, and 103.2 eV which are attributed to the Si‐C, Si‐C‐O, and Si‐O bonds, respectively [37, 46, 52]. This result agrees with the FTIR findings where, for both Si‐C- and Si‐O-stretching modes, the vibration frequencies increased with increasing O content (incorporation of more electronegative atom).
In the case of the as-deposited SiO2-like material, the Si 2p spectrum shifted to even higher binding energy and yielded two peaks centered at 102.3 and 103.5 eV which are related to two different Si‐O electronic configurations [37, 53, 54].
4. Optical properties of silicon oxycarbide materials
The structural evolution of the as-deposited silicon oxycarbide materials along the SiCxO2(1−x) stoichiometry was also reflected in their optical properties. The evolution of the index of refraction and the optical gap as a function of the C content of the materials and upon annealing treatments will be addressed in the following subsections.
4.1. Refractive index
The refractive index (
The FTIR measurements in Figure 3 show a decrease in the Si‐O‐Si bond angle with C addition, which is expected considering the difference in the electronegativity between C and O. The mass density and Si atomic content (Si atoms possess higher electronic polarizability than O) of samples with higher carbon concentration is larger as opposed to SiO2-like samples [56]. It is therefore suggested that the increased index of refraction with increasing [Si‐C]/[Si‐O] ratio is due to the variations in bond angles, larger mass density, and higher Si content. Further increase of the refractive index of all three classes of silicon oxycarbide materials is also observed upon post-deposition annealing beyond 900°C (Figure 8b). This behavior is expected considering the densification of SiCxOy materials upon annealing as shown in Figure 1c [22, 37].
4.2. Optical gap
The observed increase in the refractive index
Similar to
In addition, as seen in Figure 9a,
As seen in Figure 9b, both the
Furthermore, the decrease in optical gap with increased annealing temperatures seems to correlate well with the loss of hydrogen. As seen in Figure 9c, hydrogen reduction is accompanied by a decrease in the energy band gap in the annealed SiCxOy films. The presence of Si‐H bonds (~2000 cm−1) can be an indication of dangling-bond passivation in the material, while the presence of other hydrogen-bonding configurations may be responsible for forming recombination centers and increasing the degree of structural disorder [61, 62].
In the as-deposited high C content samples, the Si‐H
4.3. Structural disorder and dangling bonds
Regarding the electronic structure of amorphous materials, it is common to expect the presence of band-tail states and localized defect states. These states exist due to the structural disorder in materials and may have a significant effect on the material’s performance even at low concentrations. Therefore, it is important to elaborate on the degree and impact of structural disorder in the CVD-grown SiCxOy systems.
As it was discussed earlier, thermal annealing causes lowering of the optical gap owing to the increased optical absorption observed in the SiCxOy materials grown by CVD. The latter is true even at photon energies well below the optical gaps. The enhanced sub-band-gap absorption is a result of an increased density of band-tail states and localized defect states [31]. One of the proposed mechanisms responsible for the increased density of band-tail states upon annealing is the annealing-induced enhanced bond-angle disorder due to structural reconfiguration and/or strain relaxation [65]. In this context, the FTIR full-width half maximum (FWHM) values for both the Si‐O- and Si‐C-stretching modes in the TCVD SiCxOy films, with the exception of the 1100°C-annealed SiC-like sample, increased after annealing, suggesting that thermal treatments indeed enhance bond-angle distortion (see Figure 4(b and c)).
In the case of SiC-like material annealed at 1100°C, the FTIR spectrum shown in Figure 4a suggests that a significant amount of crystallization takes place resulting in the Lorentzian line shape of the infrared absorption band. This suggests that the bond-angle disorder is dramatically reduced. However, the optical absorption for the SiC-like films annealed at 1100°C is further increased compared to its as-deposited counterparts [38]. This suggests that, in addition to bond-angle variation, a different type of disorder is also present. Such behavior is known for amorphous covalent materials where both topological and compositional disorders are present simultaneously [66].
Furthermore, the deviation of the CVD-grown SiCxOy films from the purely stoichiometric SiCxO2(1−x) shown in the ternary diagram of Figure 1 suggests that there is a small excess of carbon that can form homonuclear bonds upon annealing. Also, this deviation increases for high C concentration materials, toward the SiC-like class regime. The electronic states associated with the homonuclear bonds would exist as localized states within the gap due to their relatively weak bond strength [67, 68].
In this context, electron paramagnetic resonance (EPR) studies on SiCxOy materials, grown by CVD, showed the presence of unpaired electrons (dangling bonds) [5]. The same study proposed that one of the major EPR signals may be originated from (≡Si-)3C• radicals or associated defects with different backbone atoms bonded to C atom, such as in C‐Si‐O configuration [69]. The SiCxO2(1−x) stoichiometry trend of SiCxOy (SiCxO2(1−x)) suggests that two divalent oxygen atoms are replaced by one tetravalent C, further supporting the presence of (≡Si-)3C• radicals in our films, originating from oxygen incorporation into (≡Si-)4C structures. Consequently, the density of such radicals is expected to increase following thermal oxidation of (≡Si-)4−
5. Photoluminescence properties of silicon oxycarbide thin films and nanowires
5.1. Visible luminescence emission from SiCxOy thin films
The room-temperature photoluminescence spectra for SiCxOy samples with different C concentrations under excitation at 300 nm are shown in Figure 10a. The spectra are characterized by a broad emission in the whole visible range (350–800 nm). The photoluminescence excitation (PLE) intensity (Figure 10b) shows the presence of a shoulder at low excitation energies (<3.5 eV) and a steep increase at high excitation energies (>~3.7 eV). This was fit linearly and the intercept of the fitted line at the photon energy abscissa was taken as the
Based on optical, FTIR characterizations, passivation experiments and electron paramagnetic resonance measurements, defect-related mechanisms and small sp2-carbon clusters that could be attributed to white luminescence from SiCxOy thin films were excluded [3, 5]. For example, structural defects typical seen with EPR measurements in silicon oxides, which cannot be explicitly controlled by material processing and are not directly related to the stoichiometry and material characteristics, such as Si-related neutral oxygen vacancies or non-bridging oxygen-hole centers, were ruled out. Instead, the observed intense white luminescence originates from the recombination of photogenerated carriers between the energy bands and at their tail states associated with the Si‐O‐C/Si‐C bonds, as indicated by the direct correlation between the integrated luminescence intensity and the Si‐O‐C bond density (see Figure 11) [3].
On this, the integrated red and green PL emission bands were added and their normalized integrated values were plotted along with the intensity of the red-shifted Si‐O-rocking (related to Si‐O‐C bond density) mode as a function of C content (Figure 11) [3]. A strong correlation between the emitted luminescence and the Si‐O‐C bond density in SiCxOy was revealed. This suggests that the emitted luminescence can be directly associated with the Si‐O‐C structure in the materials [5]. Additionally, the PL/PLE analysis revealed a strong similarity in the PLE behavior for the green/blue band emissions between the SiC control and SiCxOy, suggesting that C‐Si/C‐Si‐O bonding may be also responsible for the excitation path of the observed luminescence in SiCxOy.
5.2. Band-tail recombination model
Representative forming gas-annealed SiCxOy samples were additionally studied with means of PL and PLE analyses as presented in Figure 12 [22].
The evolution of the PL peak position in SiCxOy was supported by the band-tail states recombination model, typical of amorphous materials [71]. Upon excitation, the photogenerated carriers thermalized to lower energies associated with band edges (hopping edge) before they recombined radiatively (energy plateau in Figure 12b). The PL peak position increased with the excitation energy up to the
PL dynamics experiments showed a fast decrease of the PL intensity, suggesting the existence of fast recombination mechanisms in SiCxOy, and the PL decay spectra followed a stretched exponential law [23]. These findings further support a band-tail states recombination model, in which carriers recombine by tunneling between spatially separated conduction and valence band-tail states. Due to the diffusivity/tunneling of the photogenerated carriers during thermalization in the band-tail states before they recombine and the inhomogeneous constitution of the band-tail states related to C‐Si/Si‐O‐C bonding, where each state contributes with a slightly different PL lifetime, a distribution of lifetimes is expected, hence the stretched exponential behavior. This is inherent in disordered semiconductors, such as amorphous Si:H, C:H, SiCx, and SiNx, due to the broad and variable spatial density of these states [72–74].
To further elucidate on the band-tail states recombination model in SiCxOy, the PL luminescence decay at different emission energies was additionally investigated. The time evolution of the PL line shape is presented in Figure 13a. It can be seen that during the first 1 ns of the luminescence decay, the integrated intensity substantially decreased and a red shift, equal to Δ
Furthermore, the PL lifetimes increased as the emission energy was decreased, in agreement with electron-hole (e-h) hopping within lower-energy tail states, as the rate of thermalization decreases significantly due to the rapidly decreasing density of band-tail states [75]. Similar behavior was observed in amorphous semiconductors where the luminescence lifetime increase is attributed to e-h hopping [74]. Upon FG annealing, the average lifetimes exhibited higher values compared to their as-deposited counterparts (Figure 13c). Furthermore, the integrated PL intensity of the FG-annealed films increased significantly along with a blue shift in peak position (e.g., six times for the sample with 21 at.% C with 0.2 eV blue shift). This change in lifetimes and PL intensity can be attributed primarily to the passivation of non-radiative recombination centers present in lower-energy portion of the band-tail states of the as-deposited films.
The increase in the luminescence lifetime in SiCxOy with low
Additionally, it is expected that in the case of high
5.3. Visible luminescence from SiCxOy nanowires
Figure 14a and b shows the normalized room-temperature PLE and PL emission spectra of SiC0.34O1.52 nanowire arrays. To better understand the visible light emission in SiCxOy NWs, their PL and PLE properties were explored in conjunction with data from the thin film of the same composition. The PL emission spectrum of SiC0.34O1.52 NWs exhibits broadband characteristics ranging from blue to deep red, while the PLE spectrum monitored at the peak luminescence emission (~550-nm) spans from near-UV to blue/green regions of the spectrum.
PLE analysis suggested that the observed luminescence from SiCxOy NWs is related to radiative recombination of photo-excited carriers in band-tail states associated with C‐Si/Si‐O‐C-bonding groups [5, 22]. Furthermore, a supplementary mechanism, in addition to the proposed band-tail states recombination process, may be needed in order to take into account the reduced dimensionality of SiCxOy NW. In the case of the NW structure with spatially confined volume, the statistics of the lowest energy states due to Si‐C bonds may be excluded [77]. Furthermore, by nanostructuring the recombination volume is reduced, thus, the tunneling probability of a carrier between adjacent states at similar energy is expected to decrease. Consequently, the carriers will be exposed to a smaller number of non-radiative sites [23, 77, 78], resulting in enhanced PL efficiency, PL blue shift, and faster lifetimes with a tighter distribution as observed in SiC0.34O1.52 (12 at.% C) (Figure 14c). However, it is important to emphasize that the effects of NW-related surface recombination and optical scattering may not be ruled out.
5.4. Photo-stability upon thermal oxidation and excitation power density
In addition to PL and PLE studies, the luminescence performance of the SiC0.34O1.52 NW arrays was investigated as a function of oxidation temperature and excitation power density [22].
The oxidation treatments up to 250°C (Figure 15a) did not cause any change in the composition and bonding configuration of the SiC0.34O1.52 material. Consequently, the emission intensity and line shape, from SiC0.34O1.52 nanostructured arrays, did not reveal any changes, suggesting the absence of luminescence degradation due to thermal oxidation. This stability can be attributed to the similarity of the oxidation rates and activation energies of SiCxOy materials to those of SiC [79]. Hence, SiCxOy materials (with higher carbon content) appear to be significantly more resistant to oxidation annealing, compared to its SiO2-like counterparts (with very low carbon content) [3].
The excitation power dependence study (Figure 15b) showed a linear behavior in the integrated PL with respect to the power density (up to ~800 W/cm2). Similar to the thermal-oxidation study, the peak position and the line shape of the PL in the nanowires remained unchanged, suggesting good emission stability at high excitation intensity [22]. The latter is more supporting evidence suggesting the absence of defect-related localized emission in the proposed origin of the visible luminescence from SiCxOy nanowires, as no PL saturation at high powers was observed (e.g., due to state filling of the localized states) [80, 81].
6. Concluding remarks
The optical and luminescence properties of silicon oxycarbide thin films and nanostructured (e.g., NW) arrays are correlated to their synthesis routes, structural properties, and bonding configuration. The composition of the chemically CVD-grown SiCxOy thin films approximate the SiCxO2(1−x) (0 < x < 1) stoichiometry. The index of refraction increases linearly as the [Si–C]/[Si–O] bond-area ratio increases, accompanied by a linear decrease of the optical gap. The white (red, green, and blue) emission can be achieved simultaneously from the same SiCxOy film following a single-deposition process, without the complications encountered in the case of using nanocrystals (e.g., Si, SiC). In particular, it was determined that the white PL emission of SiCxOy thin films is strong enough to be seen with the naked eye under bright room conditions. Based on the PL, FTIR, and EPR analyses, typical structural defects in oxides were ruled out as the mechanism for white luminescence from SiCxOy. Instead, the observed intense visible luminescence originates from the recombination of photo-generated carriers between the energy bands and at their tail states associated with Si‐O‐C/Si‐C bonds. In this regard, the potential advantages offered by our proposed approach of SiCxOy thin films and NWs range from color tunability, thermal/photo-stability to enhanced light extraction efficiency and from cost reduction to environmental considerations. To this end, these compelling behaviors may provide a pathway for further controlling and enhancing the thermal stability and PL yield of white light emission from such films and nanostructured materials through optimal engineering of Si‐O‐C/Si‐C bonds in the matrix.
Acknowledgments
The authors would like to thank Dr. Mengbing Huang, of the Colleges of Nanoscale Sciences and Engineering at SUNY POLY, for his contribution to the work presented herein.
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