A. Structural and optical properties

Table I summarizes the compositions of representative SiC _x O _y thin films deposited at various O₂ flow rates and pressures, as determined by XPS and RBS measurements. The compositional evolution of the SiC _x O _y thin films was plotted in a Si–C–O ternary diagram (Fig. 2) along with the SiC _x O_2(1−x) (0 < x < 1) line, with the latter representing stoichiometric silicon oxycarbide without any excess carbon. The upper and lower limits in SiC _x O_2(1−x) correspond to SiC and SiO₂, for respectively, x = 1 and x = 0. ^{Reference Kroll21,Reference Bréquel, Parmentier, Walter, Badheka, Trimmel, Masse, Latournerie, Dempsey, Turquat, Chomel, Le Neindre-Prum, Jayasooriya, Hourlier, Kleebe, Soraru, Enzo and Babonneau22}

TABLE I.

Elemental composition and stoichiometry of representative thermal CVD-grown SiC _x O _y samples at different deposition process parameters (precursor flow rate: 10 sccm; Ar flow rate: 450 sccm; and deposition temperature: 800 °C).

FIG. 2.

Si–C–O ternary diagram with the compositions of AD silicon oxycarbide (SiC _x O _y ) materials as they were determined by XPS and RBS measurements. The solid line shows the composition of the stoichiometric silicon oxycarbide, SiC _x O_2(1−x) (0 < x < 1), along with the positions of the SiO₂ and SiC, for respectively, x = 0 and x = 1.

As is shown in Fig. 2, the composition of the thermal CVD grown SiC _x O _y was properly tailored to closely follow the pure stoichiometric oxycarbide formula over the whole range of the synthesis process parameters. This is in contrast to nonstoichiometric compositions reported for SiC _x O _y samples with SiC/SiO₂ phases and free carbon, as it is shown in Fig. 1 of Ref. Reference Bréquel, Parmentier, Walter, Badheka, Trimmel, Masse, Latournerie, Dempsey, Turquat, Chomel, Le Neindre-Prum, Jayasooriya, Hourlier, Kleebe, Soraru, Enzo and Babonneau22. The SiC _x O_2(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 SiC _x O _y materials increases. The latter is also proposed owing to the linear increase of the [Si–C]/[Si–O] bond–area ratio, calculated by FTIR analysis, with increasing C/O content ratio [Fig. 3(a)].

FIG. 3.

(a) The [Si–C]/[Si–O] bond–area ratio plotted as a function of the C/O content ratio showing a linear increase with increasing C/O; (b) The FTIR absorptions of the Si–O–C bonding configuration around ∼437 cm⁻¹ of AD SiC _x O _y for different C contents. For comparison purposes the FTIR absorption of a thermally-grown SiO₂ control sample is also displayed (dotted gray line). The peak positions of the AD SiC _x O _y appear red-shifted compared to the SiO₂ control, due to the incorporation of less electronegative C. The increase of the C content in SiC _x O _y results in the decrease in the Si–O–C intensity; (c) The FTIR absorption spectra of the AD SiC _x O _y and SiO₂ control in the 550–1600 cm⁻¹ range. The inset displays the TO Si–O vibration frequency versus C content.

Figure 3(b) shows the FTIR data of SiC _x O _y thin films at ∼440 cm⁻¹, attributed to Si–O–C bonding groups as it has been reported elsewhere, ^{Reference Gallis, Nikas, Suhag, Huang and Kaloyeros3,Reference Wolfe, Hinds, Wang, Lucovsky, Ward, Xu, Nemanich and Maher23–Reference Kim, Hwang, Kim, Kim and Lee25} which is caused by the ≡Si–O–Si≡ rocking mode due to out of plane motion of O in O_(3−x)C _x ≡Si–O–Si≡C _y O_(3−y) and is virtually unaffected by the bond angle variation. ^{Reference Tolstoy, Chernyshova and Skryshevsky24,Reference Lehmann, Schumann and Hubner26} Additionally, the important role of Si–O–C bonding in the observed visible luminescence from SiC _x O _y was demonstrated, as its emitted luminescence intensity increases in direct correlation with the increased density of Si–O–C. ^{Reference Nikas, Gallis, Huang, Kaloyeros, Nguyen, Stesmans and Afanas'ev27} The replacement of two bridging oxygen atoms by one carbon atom, as dictated by the SiC _x O_2(1−x) stoichiometry, results in the observed decrease in the Si–O–C intensity. In samples with C content higher than 25 at.%, the density of Si–O–C bonding groups is significantly suppressed, suggesting that the backbone network associated to the ≡Si–O–Si≡ bonding in SiC _x O _y transitioned toward more SiC-like structures.

Figure 3(c) shows the FTIR spectra of SiC _x O _y in the 550–1600 cm⁻¹ range. The spectra are characterized by two major absorption bands and a shoulder assigned to the Si–C stretching and Si–O transverse and longitudinal stretching vibration modes, centered at, respectively, ∼800, ∼1000, and ∼1150 cm⁻¹. ^{Reference Gallis, Nikas, Eisenbraun, Huang and Kaloyeros19,Reference Gallis, Nikas, Huang, Eisenbraun and Kaloyeros20} FTIR studies also showed the absence of C=C bonds at ∼1644 cm⁻¹, which would result from the sp ²-carbon molecular orbitals. ^{Reference Gallis, Nikas, Eisenbraun, Huang and Kaloyeros19} The substitution of O with C atoms is reflected to the increase of the Si–C and subsequent decrease in the Si–O absorption band intensities. Furthermore, it was found that the Si–O vibration frequency red-shifts in a linear fashion with increasing C content [inset of Fig. 3(c)], as the less electronegative C atoms replaces O atoms, suggesting that the Si–O–Si bond angle between tetrahedra decreases. ^{Reference Tolstoy, Chernyshova and Skryshevsky24,Reference Kim, Hwang, Kim, Kim and Lee25} Consequently, the mass density of the as-deposited (AD) SiC _x O _y is expected to increase with increasing C content, which is in agreement with our previous findings. ^{Reference Gallis, Nikas, Eisenbraun, Huang and Kaloyeros19}

The structural evolution of the AD silicon oxycarbide materials along the SiC _x O_2(1−x) stoichiometry was also reflected in their optical properties. For example, the E ₀₄ gap corresponds to the energy where the absorption coefficient (α) is equal to 10⁴ cm⁻¹ (α(E ₀₄) = 10⁴ cm⁻¹), and the index of refraction (n). Figure 4(a) shows the E ₀₄ and the Tauc E _g gaps that were determined from optical absorption measurements using the Tauc's plots, as a function of the [Si–C]/[Si–O] bond–area ratio. Both E ₀₄ and E _g were found to decrease with the increasing [Si–C]/[Si–O] ratio. This behavior can be explained by considering the smaller splitting energy difference between the bonding and anti-bonding electronic states in Si–C bonds in comparison to Si–O bonds. ^{Reference Sun28} Thus, the incorporation of additional carbon enhances the bond density difference between Si–C and Si–O, as represented by the [Si–C]/[Si–O] ratio, thereby further decreasing the splitting of their corresponding bonding and antibonding states.

FIG. 4.

(a) E ₀₄ and Tauc E _g energy values for different SiC _x O _y materials versus the [Si–C]/[Si–O] bond ratio. The E _g was determined based on the Tauc's equation, (αhν)^1/2 = B ^1/2(hν − E _g), where α is the absorption coefficient, hν is the photon energy and B is the slope, which is inverse proportional to the width of the tail band. The solid lines depict the linear [Si–C]/[Si–O] dependence; (b) the index of refraction at 800 nm of different SiC _x O _y given as a function of the [Si–C]/[Si–O] bond ratio. The solid line is the linear fit of the displayed data.

Furthermore, an examination of the index of refraction (n) of SiC _x O _y showed a linear increase of n with increasing [Si–C]/[Si–O]. The linear behavior of n versus [Si–C]/[Si–O] over the range analyzed can be partly associated with an increase of mass density in SiC _x O _y with increasing [Si–C]/[Si–O]. ^{Reference Gallis, Nikas, Eisenbraun, Huang and Kaloyeros19} Additionally, the increase of the C content results in an increased molecular polarizability and decreased molecular weight of SiC _x O _y , both contributing to the increase of the index of refraction, as dictated by the Lorenz–Lorenz equation. ^{Reference Gallis, Nikas, Huang, Eisenbraun and Kaloyeros20}

By virtue of the relationship of the carbon-over-oxygen-dependent index of refraction and the tunable optical gap of SiC _x O _y materials, properly tailored SiC _x O _y NW arrays were fabricated. Specifically, 400 nm-periodically-ordered SiC_0.34O_1.52 and SiC_0.25O_1.70 NW arrays were used due to their high density of Si–O–C bonding. Moreover, large SiC _x O _y NW structures (3 × 3 mm arrays of 70 nm-width NWs) were also fabricated on intrinsic Si substrates to perform bonding configuration analysis. The normalized absorbance FTIR spectra of both the AD SiC_0.34O_1.52 thin film and the as-fabricated NWs are shown in Fig. 5. It was observed that the bonding configuration of the SiC_0.34O_1.52 NW system was maintained, along with its structural characteristics, as inferred by the relative intensities between Si–C and Si–O bond densities, which remain the same as in the case of the AD thin film counterpart. The latter can be also observed in the unchanged bonding configuration of C=O related vibration mode at ∼1900 cm⁻¹. A slight relative increase of the Si–O vibration mode at ∼1200 cm⁻¹ may be attributed to a very small oxidation induced on the surface of the as-fabricated NWs.

FIG. 5.

Normalized FTIR absorbance spectra of 70 nm-width SiC_0.34O_1.52 NW array and its thin film counterpart. In the inset an SEM image of the 3 × 3 mm² array used for the measurement is shown, which consists of 36 blocks of 490 × 490 μm² NW arrays with 10 μm spacing between the blocks.

B. PL properties

Figure 6(a) shows the normalized room-temperature PL excitation (PLE) and emission spectra of SiC_0.34O_1.52 NW arrays. The PLE spectrum monitored at the peak luminescence emission (∼550 nm) spans the range of 300–550 nm, from near-UV to blue/green region of the spectrum. The PL emission spectrum of SiC_0.34O_1.52 NWs exhibits broadband characteristics ranging from blue to deep red, that is, from 470 to ∼740 nm, with a full width at half maximum (FWHM) of ∼0.5 eV.

FIG. 6.

(a) Room-temperature PLE spectra of 120 nm-thick SiC_0.34O_1.52 thin film and its corresponding 70 nm-width SiC_0.34O_1.52 NWs (400 nm pitch) measured at its emission peak; (b) room-temperature PL spectrum of the as-fabricated SiC_0.34O_1.52 NWs excited at λ_exc = 457 nm. The inset shows the camera photo of the observed strong visible PL under blue light excitation and its corresponding optical photo. The dimension of the studied NW array structure is 50 μm × 500 μm and the PL photo was taken with a common camera with the use of an optical setup equipped with a 50× lens and a 475 nm long-wave pass filter.

To better understand the visible light emission in SiC _x O _y NWs, their PL and PLE properties were explored in conjunction with data from their thin films counterparts, which were also used as baseline for comparison purposes. The analysis revealed almost identical PLE emission spectra between NWs and thin films [Fig. 6(a)], indicating that there is no change in the excitation path and emission origin of the observed luminescence in SiC_0.34O_1.52 NWs. This is also in agreement with the infrared (IR) absorption data comparing that for NWs with their thin films counterparts. There was no significant change in the bonding configuration (e.g., [Si–C]/[Si–O] bond–area ratio) in the NW IR spectrum upon nanofabrication. The observed visible light emission from SiC _x O _y NWs is therefore likewise assigned to radiative recombination of photo-excited carriers in electronic states and/or band tail states associated with C–Si/Si–O–C bonding groups. ^{Reference Nikas, Gallis, Huang, Kaloyeros, Nguyen, Stesmans and Afanas'ev27} This finding is in contrast to PL emission from structural defects (e.g., Si-related oxygen deficiency centers and/or neutral oxygen vacancies) typical seen in oxides. ^{Reference Nikas, Gallis, Huang, Kaloyeros, Nguyen, Stesmans and Afanas'ev27,Reference Skuja29}

Due to the amorphous nature of the SiC _x O _y NWs, their broadband PL characteristics may be thereby attributed to the inhomogeneous constitution of the electronic states and the band tails related to C–Si/Si–O–C bonding, i.e., each contributes with a slightly different energy resembling small differences in the local environment of the C–Si/Si–O–C bonding. The broadband emission spectrum, which is determined by static disorder, has been observed in amorphous SiC _x luminescent materials. ^{Reference Tessler and Solomon30}

C. Thermal-oxidation and photoexcitation power-density stability

To further understand the properties of luminescent nanostructured materials, a study of their photo and thermal properties is required. Thus, in addition to PL emission and photo-excitation measurements, the PL performance of the SiC_0.34O_1.52 NW system was studied as a function of temperature, varying from room temperature to 250 °C, in either oxygen or argon ambient. To study saturation processes, the NWs were exposed to high intensity blue laser radiation (up to ∼800 W/cm²).

First, thermal treatments did not cause any change in the chemical structure of the SiC_0.34O_1.52 nanostructured material, as observed in IR absorption measurements. With respect to subsequent PL measurements of the NWs, their emission intensity dependence on the annealing temperature did not reveal any change (as seen in Fig. 7) indicating that there is no thermal-oxidation degradation in our nanostructures. The oxidation rates and the activation energies of SiC _x O _y materials are similar to those for SiC, whereas the diffusion of oxygen into the matrix, with simultaneous CO diffusion outwards, typically governs the oxidation rate. ^{Reference Chollon31}

FIG. 7.

PL spectra of the AD and 2 h-oxygen-annealed SiC_0.34O_1.52 NW at various temperatures (excitation wave length λ_exc = 457 nm). The inset shows the integrated PL intensity versus annealing temperature.

The reaction of the oxidation for SiC _x O _y can be described by the following equation: ${\rm{Si}}{{\rm{C}}_x}{{\rm{O}}_{2{{\left( {1 - x} \right)}_{{\rm{(s)}}}}}} + {{3x} \over 2}{{\rm{O}}_{{2_{{\rm{(g)}}}}}} \to {\rm{Si}}{{\rm{O}}_{{2_{{\rm{(s)}}}}}} + x{\rm{C}}{{\rm{O}}_{{\rm{(g)}}}}$ . Hence, the SiC _x O _y material (with higher carbon content) appears to be significantly more resistant to oxidation annealing, for example, compared to its SiO₂-like counterpart (with very low carbon content). ^{Reference Gallis, Nikas, Suhag, Huang and Kaloyeros3} This behavior is attributed primarily to the fact that the higher carbon SiC _x O _y has a higher density of Si–C bonds (as seen in Fig. 3), which is in turn translates into higher mass density than that of SiO₂-like, and thus more effectively preventing oxygen diffusion inwards. This compelling behavior may provide a pathway for further controlling and enhancing the thermal stability of visible light emission at elevated temperatures (>500 °C) from such nanostructured materials through optimal engineering of Si–C/Si–O bonds in the matrix. In addition to the fact that the emission intensity is temperature independent, the lineshape of the observed visible light emission remains also unchanged (Fig. 7), showing good light emission stability for the SiC _x O _y nanostructured materials produced herein.

Second, the power dependence results at high continuous wave (cw)-excitation levels revealed that the integrated PL of SiC_0.34O_1.52 NWs remained linear over the investigated probing laser intensity range (100–800 W/cm²), as depicted in Fig. 8. It is worth noting that there was no change in the peak position and the lineshape of the PL in the NWs, revealing good emission stability at high excitation intensity.

FIG. 8.

Room-temperature PL spectra of the 70 nm-width SiC_0.34O_1.52 NWs for different cw excitation probing power-densities (excited at the λ_exc = 457.9 nm of an Ar ion laser). The inset shows the linear behavior of the integrated PL intensity as a function of the excitation power density.

This provides further evidence supporting the proposed origin of the visible emission in SiC _x O _y NWs and the absence of defect-related localized emission, as the latter is typically associated with PL saturation or deviation from the linearity (e.g., slope < 1) at high probing power intensity (e.g., due to state filling of the localized states). ^{Reference Tayagaki, Fukatsu and Kanemitsu18,Reference Grieshaber, Schubert, Goepfert, Karlicek, Schurman and Tran32} Moreover, the influence of nonradiative AR on the overall luminescence efficiency of SiC _x O _y NWs will need to be considered in conjunction with the development of such nanostructures to achieve an improved PL efficiency. The fast AR becomes important under high excitation power densities and also when electron–hole pairs are localized in a small volume, as in Si nanocrystals with long radiative lifetimes, causing PL intensity saturation (bulk Si and Si-based materials). ^{Reference Tayagaki, Fukatsu and Kanemitsu18,Reference Kovalev, Diener, Heckler, Polisski, Kunzner and Koch33}

The linear behavior of the PL on the power density may be considered as indicative of a suppression of AR in SiC _x O _y NWs. The very short radiative lifetime of a few hundreds of picoseconds in the SiC _x O _y material produced herein (associated data will be presented in a subsequent report) may imply that fast nonradiative processes cannot dominate at high probing power densities. A decrease of the AR rate [(τ_nr)⁻¹] has been observed in Si NW structures, ^{Reference Guichard, Kekatpure, Brongersma and Kamins15} which is imposed by bimolecular annihilation so that the AR rate is characterized by a linear term (∼n), instead of quadratic term ∼n ², of the carrier density (n). To obtain more insight into the PL-linearity-on-excitation-power behavior, dedicated studies including temperature dependence PL transient and time-resolved emission spectroscopy measurements will be carried out on SiC _x O _y NWs with different width dimensions (e.g., sub-15 nm).