Nearly warm white-light emission of silicon-rich amorphous silicon carbide

Abstract

Nearly warm white-light photoluminescence of silicon-rich amorphous silicon carbide (Si-rich a-Si_xC_1−x) films grown by middle-temperature plasma-enhanced chemical vapor deposition is demonstrated. The argon diluted silane, methylsilane and methane environment and low RF plasma power regime assist the achievement of extremely low oxygen incorporation and relatively high tenability of the composition ratio x. From XPS analysis, the decomposed Si_2p orbital electron related energy distribution reveals plentiful Si–Si bonds, gradually reduced Si–C bonds and suppressed Si–O and C–Si–O bonds at a deposition temperature of 650 °C. The Si nanocrystals (Si-ncs) with estimated grain sizes of 2.9 ± 0.3 nm contribute to structural and optical properties. Moreover, the Si–C stretching mode is slightly red-shifted from 806 to 785 cm⁻¹ with increasing deposition temperatures from 450 °C to 650 °C, and the reduced strength and quantity of the Si–C bonds are also observed. The luminescent a-Si_xC_1−x film can only be obtained when deposition occurs at 650 °C since the middle-temperature synthesis induces fractional precipitation of Si-ncs. The post-annealing at 1100 °C for 90 minutes is optimized to re-grow the Si-rich Si_xC_1−x matrix, which leads to the crystallization of Si_xC_1−x with self-aggregated Si-ncs. Such annealed a-Si_xC_1−x films with warm white-light photoluminescence can serve as potential solid-state phosphorous materials for future white-light light-emitting applications.

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Introduction

Amorphous silicon carbide (a-Si_xC_1−x) has gained much attention in recent years owing to its superior characteristics of widely tunable energy bandgap (1.7–3.5 eV), high electron saturation velocity (2 × 10⁷ cm s⁻¹), and high breakdown electric field (2.5 × 10⁶ V cm⁻¹).¹ The non-stoichiometric a-Si_xC_1−x film has also emerged as a potential candidate for high-temperature, high-frequency, high-radiation and high-power electronic and optical devices in view of its prominent mechanical properties and chemical stability.² In most studies, three main types of crystalline SiC materials such as 3C, 4H, and 6H species were emphasized. The 3C–SiC is denoted as β-SiC with ABC stacking sequence and zinc blende structure, whereas 4H–SiC and 6H–SiC with hexagonal structure belongs to α-SiC family.^3–5 The bandgap energies of stoichiometric 3C–SiC, 4H–SiC and 6H–SiC are 2.36, 3.23, and 3 eV, respectively.^6–8 During past years, the 3C–SiC epilayer with oriented plane of (100) synthesized on the Si substrate shows 20% mismatch on lattice constant and 8% mismatch on thermal expansion coefficient between each other.³ The promising utilization of SiC film in numerous kinds of optoelectronic and photonic devices has included the p-i-n thin film light emitting diodes (LED),⁹ the thin film transistors,¹⁰ and the tandem solar cells.¹¹ Versatile synthesizing methods for a-Si_xC_1−x film were also implemented, such as hot-wire chemical vapor deposition (HWCVD),¹² catalytic chemical vapor deposition (Cat-CVD),¹³ and plasma-enhanced chemical vapor deposition (PECVD).^14–17 However, the hydrogen-free and middle-temperature PECVD synthesis of Si-rich a-Si_xC_1−x film with buried Si nanocrystals (Si-ncs) has yet to be investigated. In this work, the a-Si_xC_1−x films grown by using PECVD with a specific recipe of argon (Ar) diluted silane (SiH₄), methane (CH₄) and methylsilane (SiH₃CH₃) based precursor gases are demonstrated, which is a particularly performance with changing the deposition from low to middle temperature for flexibly detuning composition ratio x. Owing to formation of plentiful Si–Si bonds and the suppression of oxygen invasion, the thin a-Si_xC_1−x films form better crystallinity with larger photoluminescence (PL) intensity to enable a higher radiative recombination efficiency. The X-ray photoelectron spectroscopy (XPS), X-ray diffraction (XRD), Fourier transform infrared spectroscopy (FTIR), and PL results of the a-Si_xC_1−x with varied composition ratios when growing at different deposition temperatures are analyzed and discussed. With optimized synthesis condition, the Si-rich environment further enhances the Si–Si bond strength and Si-ncs self-aggregation in Si_xC_1−x matrix.

Experimental

The a-Si_xC_1−x films were deposited on p-type Si (100) substrate by PECVD system with CH₄, Ar diluted SiH₄ and SiH₃CH₃. The composition ratios x is defined as x = Si/(Si + C). Prior to the synthesis, the Si wafers were additionally dipped in buffered oxide etchant for 5 min to remove surface native oxide. The gaseous flow of SiH₄ was 65 sccm and the fluence ratio defined as g = CH₄/(SiH₄ + SiH₃CH₃ + CH₄ + Ar) was fixed at 60%. The deposition pressure was 0.16 torr and deposition temperature changed from 450 °C to 650 °C at an increment of 100 °C, and a subsequent post-annealing treatment was performed at 1100 °C for 90 minutes. During the XPS analysis in vacuum chamber at around 10⁻⁶ torr, the a-Si_xC_1−x sample was excited by Mg K_α line at 1253.6 eV, and the Si_2p, C_1s, O_1s orbital electrons can be emitted from the sample surface for diagnosis of composition ratio and O/Si ratio of a-Si_xC_1−x films. The crystallinity of the middle-temperature as-grown and annealed a-Si_xC_1−x samples was determined by grazing incidence XRD with a Cu K_α radiation source (λ = 1.540562 Å) at a voltage of 40 kV and a current of 200 mA. The FTIR spectroscopy (Thermo Nicolet NEXUS470) with a spectral resolution of 4 cm⁻¹ was utilized to determine the bonding geometries in a-Si_xC_1−x network. At last, the room-temperature continuous-wave PL was performed by using a GaN laser diode pumping with an average power of 30 mW at 405 nm. The PL ranged from 400 to 900 nm was resolved by using a monochromator with a 3000 groove per mm grating.

Results and discussions

The surface images of a-Si_xC_1−x samples obtained by varying the deposition temperature from 450 °C to 650 °C are shown in Fig. 1. Fig. 2 and 3 demonstrate the typical broad-scan XPS spectra of as-grown a-Si_xC_1−x films and corresponding C/Si and O/Si composition ratios, respectively.

Fig. 1 Photographs on surface images of the Si-rich a-Si_xC_1−x deposited at temperature of 450 °C (left), 550 °C (middle), and 650 °C (right).

Fig. 2 Typical broad-scan XPS spectra of as-grown a-Si_xC_1−x samples deposited with fluence ratio g of 60% at various deposition temperatures of 450 °C (upper), 550 °C (middle) and 650 °C (lower).

Fig. 3 Composition and O/Si ratios as a function of deposition temperature calculated by XPS analysis.

In Fig. 2, the Si concentration is observed to increase from 64.7% to 71.6% as high temperature circumstance provides sufficient thermal energy on SiH₄ species to dissociate more Si–H bonds, which also transfers larger kinetic energy to decomposed Si atoms so as to achieve more increasing Si–Si bonds in the a-Si_xC_1−x films. In addition to the suppression on oxygen content from 7.6% to 5.5%, the carbon content also decreases from 27.7% to 22.8%.

At higher temperature, the CH₄ molecule is easier to be decomposed into carbon and hydrogen atoms because of its higher dissociation energy. As expected, the dissociated carbon atoms easily combine with residual oxygen atoms in chamber so as to form gaseous carbon oxide (CO) or carbon dioxide (CO₂) molecules, which eventually reduce the quantity of oxygen atoms in the Si_xC_1−x film. As a supporting evidence, Dasgupta also observed similar phenomenon of reduced carbon and oxygen contents in the Si_xC_1−x film¹⁸ grown with increasing substrate temperatures, as confirmed by XPS analysis.

Fig. 3 shows the composition ratio x of the a-Si_xC_1−x films varied from 0.7 to 0.76 owing to the increased Si concentration in a-Si_xC_1−x films. To prevent the oxygen invasion in a-Si_xC_1−x films, the reduced O/Si ratio from 0.12 to 0.07 is observed with increasing the deposition temperature from 450 °C to 650 °C. With the reduction of oxygen concentration and the enhancement of Si concentration, the crystalline Si and Si-rich a-Si_xC_1−x coexist in the as-grown a-Si_xC_1−x films grown at deposition temperature of 650 °C. Note that the Si, C and O atomic contents depended on deposition temperature are clearly shown in Table 1. As mentioned above, the composition ratio x is tunable to facilitate the deposition of more Si-rich a-Si_xC_1−x films and to enable more precipitation of Si-ncs after annealing a-Si_xC_1−x network at 1100 °C. More importantly, the Si content in a-Si_xC_1−x films is directly proportional to deposition temperature during PECVD growth.¹⁹ Dasgupta et al. reported that the desorption rate at higher temperatures (>250 °C) can be enlarged, and the hydro-carbon radicals may gather to refuse C atoms from easy incorporation into the a-Si_xC_1−x films.¹⁸

Table 1 Atomic content percentage, Si/C, O/Si and composition ratio at various deposition temperatures measured by XPS

C	Si (at%)	C (at%)	O (at%)	Si/C	O/Si	Composition ratio x
450 °C	64.7	27.7	7.6	2.34	0.12	0.70
550 °C	67.9	25.8	6.2	2.63	0.09	0.72
650 °C	71.6	22.8	5.5	3.14	0.07	0.76

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With the XPS analysis on the binding energies and counts of the photoelectrons at different core orbits in Si and C atoms, the composition and stoichiometry of the a-Si_xC_1−x can be verified. The finite interacting depth between X-ray photon and a-Si_xC_1−x film is only around 5 nm, which coincides with the penetration depth of oxygen atoms into a-Si_xC_1−x surface, and the exact bonding information among Si, C, and O atoms near the surface of a-Si_xC_1−x films can be accessed. The XPS detected energy of Si_2p orbital electron can be affected by the Si–Si, Si–C, C–Si–O and Si–O with different binding energies of 99.5–99.8 eV, 100.3–100.5 eV, 101.7 eV, and 103.5 eV, respectively.^20–23 As shown in Fig. 4(a)–(c), the XPS spectra of electron energy distribution at Si_2p core levels are appropriately fitted by four Gaussian lines to analyze different binding components including Si–Si, Si–C, C–Si–O and Si–O bonds. It is observed that all the signals detailed in the XPS spectra exhibit broadened and asymmetric shapes, indicating that the orbital electrons are contributed by Si and C atoms for different bonding orders and geometries.⁹

Fig. 4 Si_2p orbital electron related XPS spectra for as-grown a-Si_xC_1−x samples grown with fluence ratio g = 60% at deposition temperatures of (a) 450 °C, (b) 550 °C and (c) 650 °C. Inset: the XPS intensities of fitted Gaussian peaks for different bonds.

The XPS analyses (by electron energy peaks of Si_2p core levels) of a-Si_xC_1−x samples grown at deposition temperatures varied from 450, 550 and 650 °C show a corresponding composition ratio of x = 0.7, 0.72 and 0.76, respectively. The residual oxygen related few Si–O and C–Si–O bonds can still be detected since the surface of the as-grown a-Si_xC_1−x film is easily oxidized. Most importantly, the XPS intensity of Si_2p electrons related to Si–Si bond is significantly enlarged with raising deposition temperature, since the middle-temperature synthesis process facilitates the growth and nucleation of Si atoms. In contrast, the signal correlated with the Si–C bond is slightly reduced with increasing deposition temperature from 450 °C to 650 °C, which is straightforward in view of the production of the more Si–Si bonds to compete with bonds between Si and C atoms, as shown in Fig. 5(a). In the meantime, the XPS intensities related to Si–O and C–Si–O bonds are considerably reduced by increasing the deposition temperature due to the suppressed oxygen incorporation. That is, the oxygen atoms are prohibited to be absorbed into the a-Si_xC_1−x films. Accordingly, the more Si-rich a-Si_xC_1−x films can be well prepared by raising deposition temperature from 450 °C to 650 °C.

Fig. 5 (a) The deposition temperature dependent intensity of Si–Si and Si–C bonds resolved from Si_2p core electron related XPS. (b) The deposition temperature dependent intensity of C–Si and C–C sp² bonds resolved from C_1s core electron related XPS.

In Fig. 6(a)–(c), the XPS spectra of electrons from C_1s core levels are well fitted by three Gaussian components, since the orbital electron energies in C_1s orbit for C–Si, C–C (sp²) and C–C (sp³) bonds are distinctive with corresponding values of 283.1–283.4 eV, 284.8 eV, and 285.4 eV, respectively.^21–24 In Fig. 6(b), the XPS intensity of C–Si bond is moderately reduced with increasing deposition temperature. Moreover, the XPS intensities of C–C (sp²) and C–C (sp³) bonds become rather trivial when increasing the deposition temperature from 450 °C to 650 °C. In general, the specific deposition condition at low RF plasma power and high temperature during PECVD process is expected to create the non-perfect dissociation environment with plentiful CH₃ radicals. According to chemical reaction, the CH₃ radical can easily combine with H to form CH₄ and then escape from the deposited films such that the C atoms are less incorporated into a-Si_xC_1−x matrix. Conversely, CH and CH₂ radicals can be prominently enhanced under high RF plasma deposition, which essentially scale down the probability of CH₄ reduction so as to facilitate the deposition of C atoms onto the substrate for a-Si_xC_1−x synthesis.

Fig. 6 C_1s orbital electron related XPS spectra for as-grown a-Si_xC_1−x grown with fluence ratio g = 60% at deposition temperatures of (a) 450 °C, (b) 550 °C and (c) 650 °C. Inset: the XPS intensities of fitted Gaussian peaks for different bonds.

With increasing deposition temperature from 450 °C to 650 °C at an increment of 100 °C, the XRD analysis shows the appearance of Si (111)-orientation (only at 650 °C) at azimuth angle of 28.5°,²⁵ as shown in Fig. 7. According to the JCPDS card number of silicon (JCPDS card no. 27-1402)²⁶ and silicon carbide (JCPDS card no. 29-1129),²⁷ the two significant peaks observed from the as-grown a-Si_xC_1−x film synthesized with g = 60% at deposition temperature of 650 °C are at azimuth angles of 28.5° and 35.6° corresponding to (111)-oriented Si and (111)-oriented 3C–SiC matrices, respectively. Nevertheless, other peaks at 47.3 and 56.1° related to (220)- and (311)-oriented Si are not observed. In view of insufficient thermal energy to enforce the precipitation of Si-ncs, the XRD peak of as-grown a-Si_xC_1−x films deposited at 450 °C and 550 °C are not found, and a large quantity of amorphous phase occurs when synthesizing the Si_xC_1−x at lower temperature deposition. According to the Scherrer formula,²⁸ the nano-scale Si grain size g_n can be calculated by g_n = βλ/[Δ(2θ)cos θ] with β denoting as a correction factor (usually 0.9), λ the incident X-ray wavelength around 0.154 nm, and Δ(2θ) the FWHM (in radian) of Si (111) peak chosen to be 3.06 rad. This gives the estimated grain size of Si-ncs as g_n = 2.95 ± 0.3 nm.

Fig. 7 XRD spectra of as-grown a-Si_xC_1−x films synthesized with g = 60% at deposition temperature varying from 450 °C to 650 °C. References: JCPDS cards no. 27-1402 for Si and no. 29-1129 for SiC.

As shown in Fig. 8(a), the SiH–(OCH₃)₃, Si–CH₃, Si–O and Si–C stretching modes are found by FTIR analysis of the a-Si_xC_1−x films grown at different deposition temperatures. In top part of Fig. 8(b), the absorption band at 1093 cm⁻¹ can be related to Si–O stretching mode, and its absorbance decreases with increasing deposition temperature from 450 °C to 650 °C due to the suppression of oxygen incorporation.

Fig. 8 (a) The FTIR spectra (b) the Si–CH₃, Si–O stretching (top part) and Si–C stretching (bottom part) signals in magnified FTIR spectra of as-grown a-Si_xC_1−x films deposited at various temperatures.

Furthermore, the FTIR signal is red-shifted from 1250 to 1230 cm⁻¹, as attributed to the bonding transformation from Si–CH₃ stretching to bending mode in view of the generation of compressive stress. With increasing the deposition temperature from 450 °C to 650 °C, the Si–C stretching is slightly red-shifted from 806 to 785 cm⁻¹ due to the reducing strength of Si–C bonds and its absorbance is decreased with reduced quantity of Si–C bonds, as shown in the lower part of Fig. 8(b). The evolution of Si–C bonding geometries is correlated with those obtained from XPS analysis in view of the decline of Si–C bonding component. The existence of SiH(OCH₃)₃ signal at various deposition temperatures is mainly attributed to (1) the insufficient RF power which hardly dissociates the SiH–(OCH₃)₃ bonds; and (2) the residual oxygen atoms which participate into the synthesis of a-Si_xC_1−x formation.

In addition, the SiH–(OCH₃)₃, Si–CH₃, Si–O and Si–C stretching modes are also analyzed at various deposition temperatures for annealed Si_xC_1−x films at 1100 °C in Fig. 9(a). In upper part of Fig. 9(b), the FTIR peak absorbance related to Si–O stretching mode for the a-Si_xC_1−x sample grown at 650 °C decreases more than others obtained at lower deposition temperatures, since more crystalline SiC formation presences to prevent the bonding between Si and O atoms.

Fig. 9 (a) The FTIR spectra and (b) the Si–CH₃, Si–O stretching (top part) and Si–C stretching (bottom part) signals in magnified FTIR spectra of a-Si_xC_1−x films deposited at different temperatures and annealed at 1100 °C for 90 minutes.

It means that the less oxygen atoms can invade the a-Si_xC_1−x matrix at higher deposition temperature even after subsequent annealing at 1100 °C for 90 minutes. In our case, the structure of the Si-rich Si_xC_1−x film can be re-grown by increasing the deposition temperature to enhance its crystallinity. From XRD analysis, the Si_xC_1−x film grown with at the deposition temperature of 650 °C reveals (111)-oriented Si-ncs and (100)-oriented 3C–SiC matrix. When enlarging the crystallinity of Si_xC_1−x film by increasing the deposition temperatures, the dangling bonds in Si_xC_1−x films can be effectively reduced. This phenomenon also decreases the probability of Si–O and C–O bonds, which can also resist the oxygen invasion. Therefore, the oxygen is suppressed at high temperatures. On the other hand, the FTIR peak of Si–C stretching mode is slightly red-shifted from 817 to 792 cm⁻¹ because the annealing treatment prefers the induction of Si-ncs precipitation and indirectly influences the Si–C bonding procedure, as shown in lower part of Fig. 9(b).

Even though the a-Si_xC_1−x film deposition is slower than those synthesized at lower temperature, but compact and better quality of the a-Si_xC_1−x film with more buried Si-ncs can be formed at middle-temperature synthesis at 650 °C. The deposition rate is also confirmed by scanning electron microscope (SEM) graphs at deposition temperature from 450 °C to 650 °C, as confirmed from Fig. 10(a)–(c). Hence, the Fig. 10(d) shows that the deposition rate of Si_xC_1−x films grown with increasing the deposition temperature from 450 °C to 650 °C (at an increment of 100 °C) is decreased from 12.5 to 11.5 nm min⁻¹. From XPS results, high temperature facilitates the formation of Si–Si bonds and suppresses Si–C bonds. The Si–Si bond length of around 2.35 Å is larger than that of Si–C at 1.87 Å. As expected, the thickness of a-Si_xC_1−x films enlarged with decreasing deposition temperature depends on the quantity of Si–C and Si–Si bonds in a-Si_xC_1−x films.

Fig. 10 SEM micrographs of as-grown Si_xC_1−x films deposited at (a) 450 °C, (b) 550 °C and (c) 650 °C. (d) The deposition rate of a-Si_xC_1−x film as a function of temperatures.

The optical bandgap of a-Si_xC_1−x films obtained at deposition temperature varied from 450 °C to 650 °C is analyzed by PL, as shown in Fig. 11(a). The PL wavelength at 646 nm for as-grown a-Si_xC_1−x sample is dominated due to the contribution of Si-ncs self-assembled at deposition temperature as high as 650 °C, whereas no PL spectra are observed for a-Si_xC_1−x samples deposited at temperatures of 450 °C and 550 °C owing to the insufficient temperature which cannot provide appropriate thermal energy to precipitate Si-ncs in Si_xC_1−x matrix. For comparison, Künle et al. also reported a PL peak at 1.85 eV (670 nm) from as-deposited Si-rich a-Si_0.8C_0.2 matrix.²⁹ Löper's work suggested that the PL peak at 656 nm results from the Si-ncs after annealing at 1100 °C for 30 minutes.³⁰ From XRD result, few Si-ncs are found at deposition temperature lower than 650 °C, meaning that middle-temperature deposition is required to self-aggregate Si-ncs, and corresponding diameter of Si-ncs can be calculated 3.02 nm from the empirical formula of λ = 1.24/[1.12 + 3.73d_Si^−1.39] given by Delerue,³¹ in which λ is the PL peak wavelength and d_Si is the diameter of Si-ncs. In comparison, the mean size of d_Si = 3.02 nm for Si-ncs estimated by Scherrer formula shows a good agreement with that of d_Si = 2.95 nm calculated from Delerue's equation.

Fig. 11 (a) The PL spectra of as-grown a-Si_xC_1−x films at different deposition temperature. (b) The PL spectra of post-annealed a-Si_xC_1−x films at 1100 °C from 30 to 120 minutes. (c) PL wavelength and intensity of the a-Si_xC_1−x films annealed at different durations.

Moreover, the Fig. 11(b) compares the PL spectra of as-grown and annealed a-Si_xC_1−x films at deposition temperature of 650 °C. As the post-annealing treatment introduces at 1100 °C from 30 to 120 minutes at an increment of 30 minutes, the intensity of PL increases from 21.3 count per nm to 46.1 count per nm due to the enhanced Si–Si bonding procedure, which is mainly caused by the increased precipitation of Si-ncs embedded in the a-Si_xC_1−x films when annealing time lengthens from 30 to 90 minutes. However, the PL conversely decreases to 41.4 count per nm after 120 min annealing in view of the oxygen re-oxidization or the Si-nc density reduction by size enlargement. As shown in Fig. 11(c), the blue-shifted phenomenon of PL peak is also observed as the annealing time increases from 30 to 120 minutes as the regrowth of Si-rich Si_xC_1−x matrix initiates concurrently. The optimized PL are obtained by annealing the a-Si_xC_1−x films at 1100 °C for 90 minutes, and their surface images are shown in left part of Fig. 12. In Fig. 12(a)–(e), the black and red lines represent the experimental data and the fitted curve, respectively, whereas blue and green dashed lines represent the fittings of Si-rich a-Si_xC_1−x matrix and Si-ncs under as-grown and annealed condition, respectively. It is seen that the emitted PL can vary its color from orange to purely white by lengthening the annealing duration from 30 to 120 minutes.

Fig. 12 The surface images (left), the PL curves and patterns, and the fitted spectra by using two Gaussian components for a-Si_xC_1−x samples at (a) as-grown and annealed at 1100 °C for (b) 30, (c) 60, (d) 90 and (e) 120 minutes.

In Fig. 13(a), the nearly warm white-light photoluminescence intensity of Si-rich a-Si_xC_1−x is enlarged by lengthening the annealing time from 30 to 120 minutes, which improves the regrowth of Si_xC_1−x grain. Concurrently, the broadband PL of Si-ncs also increase by annealing from 30 to 90 minutes but conversely decreases after 120 min treatment, as the re-oxidation and/or Si-nc reduction process occurs in a-Si_xC_1−x films after long-term annealing. In general, Si-ncs related broadband PL is diminished after thermal treatment at 1100 °C for longer than 3 hours since the regrowth of crystallized Si_xC_1−x matrix is activated then. The Fig. 13(b) demonstrates that the wavelength of PL peak shifts with changing the matrix of Si-rich Si_xC_1−x from as-grown to annealed condition. As the Si-rich Si_xC_1−x matrix regrows from amorphous to crystalline under annealing treatment, the SiC related PL peak is blue-shifted due to the enhanced SiC crystallinity. On the other hand, the size of Si-ncs is slightly enlarged due to the sufficiently long annealing time which accumulates sufficient thermal energy to further precipitate large Si-ncs. Therefore, the bond formations of Si–C and Si–Si bonds are compromised in post-annealed a-Si_xC_1−x films. As confirmed by PL analysis, the annealing condition at 1100 °C for 90 minutes is optimized to regrow the crystallized Si-rich Si_xC_1−x matrix and to optimize the self-aggregation of Si-ncs. Such a condition results in the nearly warm white-light PL emitted from the annealed a-Si_xC_1−x film to facilitate its potential as the solid-state phosphorous material for future white-light LED applications.

Fig. 13 (a) PL intensities and (b) PL wavelength of Si-rich Si_xC_1−x and Si-ncs as a function of annealing time.

Table 2 summarizes the peak wavelength and tuning range of PL emitted from different SiC materials reported in recent years. At early stage, the PL peaks of the SiC with specific structural phases were observed at fixed wavelength.^32–34 Lau's group reported that the PL peak wavelength of the microcrystalline SiC (μc-SiC) is located at 688 nm.³² Feng et al. observed the PL of 3C–SiC nanowires at peak wavelength of 450 nm.³³ Rossi and co-workers synthesized the 6H–SiC nanoparticles to further blue-shift its peak PL to 354 nm.³⁴ In order to broaden the wavelength range of PL, the SiC materials were modified by adding dopants, changing structural phases, and improving crystallinity.^35–38 Addamiano employed boron to dope the crystalline SiC with different phases for detuning their PL peak wavelengths from 535 nm to 625 nm.³⁵ Kumbhar et al. changed the composition ratios of amorphous Si_xC_1−x film by adding C₂H₂ fluence to slightly red-shift the PL from 643 nm to 651 nm.³⁶ Fan et al. fabricated the SiC nanocrystals to increase its PL wavelength from 385 nm to 415 nm.³⁷ Lin's group further synthesized the two-dimensional ultrathin SiC with its PL wavelength broadening from 420 nm to 475 nm.³⁸ From abovementioned reports, the PL peak wavelength range of the Si_xC_1−x film can only be changed by several tens nm with post treatments including ion-implantation or nano-engineering. To significantly enlarge the PL wavelength range and tunability, the composition ratio of the Si_xC_1−x is considerably changed during synthesis.⁸ Xu et al. transformed the Si_xC_1−x film from the stoichiometric to C-rich condition so as to greatly broaden the PL wavelength tunability from 633 nm to 451 nm.³⁹ In our work, the Si/C composition ratio of the Si-rich Si_xC_1−x film is largely detuned from 2.34 to 4.14 such that its PL peak wavelength can be widely red-shifted from 595 nm to 645 nm.

Table 2 Reports about PL of Si_xC_1−x in recent years

Materials	PL peak wavelength	Ref.
μC-SiC	688 nm	32
3C–SiC nanowires	450 nm	33
6H–SiC nanoparticles	354 nm	34
4H–SiC:B	535 nm	35
6H–SiC:B	580 nm
33R–SiC:B	590 nm
15R–SiC:B	605 nm
21R–SiC:B	625 nm
α-SiC	643–651 nm	36
SiC-nc	385–415 nm	37
2D ultrathin SiC	420–475 nm	38
Standard/C-rich SiC	540–620 nm	39
Si-rich a-SixC1−x:Si-nc	646–595 nm	This work

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Conclusions

The nearly warm white-light emission of Si-rich a-Si_xC_1−x grown by mid-temperature and low-plasma PECVD is demonstrated. The composition ratio x of the Si-rich a-Si_xC_1−x film enlarges from 0.7 to 0.76 with increasing the deposition temperature from 450 °C to 650 °C, which enriches Si–Si bonds and slightly suppresses the SiC formation. This phenomenon also favors the self-assembly of Si-ncs and the reduction of O/Si ratio in the Si_xC_1−x grown at 650 °C. The enhanced crystallinity of Si_xC_1−x film grown at high deposition temperature decreases the dangling bonds to reduce the formation of Si–O and C–O bonds during oxygen invasion. The Si-rich a-Si_xC_1−x film grown at 650 °C exhibits the (111)-oriented Si-ncs with average size of 2.9 ± 0.3 nm, as sufficient thermal energy is accessed to enforce the Si-ncs precipitation. For the as-grown Si-rich a-Si_xC_1−x grown at 650 °C, the intensity of PL at peak wavelength of 646 nm is enhanced up to 21.3 count per nm owing to the contribution of self-aggregated Si-ncs. When annealing at 1100 °C to 90 minutes, the Si–Si bonds in Si_xC_1−x increase to improve the precipitation of Si-ncs, which further enhance the corresponding PL intensity to 46.1 count per nm. The blue-shifted PL peak with lengthened annealing time is mainly attributed to the regrowth of Si-rich Si_xC_1−x matrix, which enables such annealed a-Si_xC_1−x films emitting the warm white-light PL. Such a Si-rich a-Si_xC_1−x can potentially serve as the solid-state phosphorous material for future white-light LED applications.

Acknowledgements

The authors thank the Ministry of Science and Technology, Taiwan, R.O.C., for financially supporting this research under grants MOST-103-2221-E002-042-MY3 and MOST-104-2221-E-002-117-MY3.

Notes and references

1. J. Wan, M. A. Capano and M. R. Melloch, Solid-State Electron., 2002, 46, 1227–1230.
2. T. Itoh, S. Tanaka, J. F. Li, R. Watanabe and M. Esashi, J. Microelectromech. Syst., 2006, 15, 859–863.
3. C. W. Liu and J. C. Sturm, J. Appl. Phys., 1997, 82, 4558–4565.
4. M. Bhatnagar, B. J. Baliga, H. R. Kirk and G. A. Rozgonyi, IEEE trans. Electron Devices, 1996, 43, 150–156.
5. C. Schroder, W. Heiland, R. Held and W. Loose, Appl. Phys. Lett., 1996, 68, 1957–1959.
6. Y. Goldberg, M. E. Levinshtein and S. L. Rumyantsev, Properties of Advanced Semiconductor Materials GaN, AlN, SiC, BN, SiC, SiGe, ed. M. E. Levinshtein, S. L. Rumyantsev and M. S. Shur, John Wiley & Sons, Inc, 2001, p. 93.
7. C. Persson and U. Lindefelt, J. Appl. Phys., 1997, 86, 5036–5039.
8. C. Persson, U. Lindefelt and B. E. Sernelius, J. Appl. Phys., 1997, 86, 4419–4427.
9. K. H. Munekata and H. Kukimoto, Appl. Phys. Lett., 1983, 42, 432–434.
10. Y. Xu, Z. Hu, H. Diao, Y. Cai, S. Zhang, X. Zeng, H. Hao, X. Liao, E. Fortunato and R. Martins, J. Non-Cryst. Solids, 2006, 352, 1972–1975.
11. K. H. Wu, Y. K. Fang, W. T. Hsieh, J. J. Ho, W. J. Lin and J. D. Hwang, Electron. Lett., 1998, 34, 2243–2244.
12. B. P. Swain and R. O. Dusane, Mater. Lett., 2006, 60, 2915–2919.
13. A. R. Shimkunas, P. E. Mauger, L. P. Bourget, R. S. Post, L. Smith, R. F. Davis, G. M. Wells, F. Cerrina and R. B. Mcintosh, J. Vac. Sci. Technol., B: Microelectron. Nanometer Struct.--Process., Meas., Phenom., 1991, 9, 3258–3261.
14. R. J. Prado, T. F. DAddio, M. C. A. Fantini, I. Pereyra and A. M. Flank, J. Non-Cryst. Solids, 2003, 330, 196–215.
15. Y. Kurokawa, S. Tomita, S. Miyajima, A. Yamada and M. Konagai, J. Appl. Physiol., 2007, 46, L833–L835.
16. H.-Y. Tai, C.-H. Cheng and G.-R. Lin, IEEE J. Sel. Top. Quantum Electron., 2014, 20, 8200507.
17. C.-H. Cheng, C.-L. Wu, C.-C. Chen, L.-H. Tsai, Y.-H. Lin and G.-R. Lin, IEEE Photonics J., 2012, 4, 1762–1775.
18. A. Dasgupta, S. C. Saha and S. Ray, J. Mater. Res., 1999, 14, 2254–2259.
19. Y. J. Huran, B. Zatko, I. Hotovy, J. Pezoldt, A. P. Kobzev and N. I. Balalykin, Czech J. Phys., 2006, 56, B1207–B1211.
20. C.-H. Cheng, J.-H. Chang, C.-I. Wu and G.-R. Lin, RSC Adv., 2015, 5, 36262–36269.
21. W. J. Zong, T. Sun, D. Li, K. Cheng and Y. C. Liang, Int. J. Mach. Tool. Manufact., 2008, 48, 1678–1687.
22. A. Santoni, J. Lancok, V. R. Dhanak, S. Loreti, G. Miller and C. Minarini, Appl. Phys. A, 2005, 81, 991–996.
23. W. K. Choi, T. Y. Ong, L. S. Tan, F. C. Loh and K. L. Tan, J. Appl. Phys., 1998, 83, 4968–4973.
24. G.-R. Lin, T.-C. Lo, L.-H. Tsai, Y.-H. Pai, C.-H. Cheng, C.-I. Wu and P.-S. Wang, J. Electrochem. Soc., 2011, 159, K35–K41.
25. D. Song, E. C. Cho, Y. H. Cho, G. Conibeer, Y. Huang, S. Huang and M. A. Green, Thin Solid Films, 2008, 516, 3824–3830.
26. JCPDS card no. 27-1402 JCPDS card no. 29-1129.
27. JCPDS card no. 27-1402 JCPDS card no. 29-1129.
28. L. V. Azaroff, Elements of X-ray Crystallography, ed. L. V. Azaroff, McGraw-Hill, 1968, p. 552.
29. M. Künle, T. Kaltenbach, P. Löper, A. Hartel, S. Janz, O. Eibl and K.-G. Nickel, Thin Solid Films, 2010, 519, 151–157.
30. P. Löper, M. Künle, A. Hartel, J. C. Goldschmidt, M. Peters, S. Janz, M. Hermle, S. W. Glunz and M. Zacharias, Quantsol 2009, March 8th–14th (2009).
31. C. Delerue, G. Allan and M. Lannoo, Phys. Rev. B: Condens. Matter Mater. Phys., 1993, 48, 11024–11036.
32. S. P. Lau, J. M. Marshall and L. R. Tessler, J. Non-Cryst. Solids, 1996, 198–200, 907–910.
33. D. H. Feng, T. Q. Jia, X. X. Li, Z. Z. Xu, J. Chen, S. Z. Deng, Z. S. Wu and N. S. Xu, Solid State Commun., 2003, 128, 295–297.
34. A. M. Rossi, T. E. Murphy and V. Reipa, Appl. Phys. Lett., 2008, 92, 253112.
35. A. Addamiano, J. Electrochem. Soc., 1964, 111, 1294.
36. A. Kumbhar, S. B. Patil, S. Kumar, R. Lal and R. O. Dusane, Thin Solid Films, 2001, 2001(395), 244–248.
37. J. Y. Fan, H. X. Li, Q. J. Wang, D. J. Dai and P. K. Chu, Appl. Phys. Lett., 2011, 98, 081913.
38. S. S. Lin, J. Phys. Chem. C, 2012, 116, 3951–3955.
39. J. Xu, T. Ma, K. Chen, J. Du, X. Huang and D. Feng, J. Non-Cryst. Solids, 1998, 227–230, 470.

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