Controlled Growth, Microstructure, and Properties of Functional Si

Article pubs.acs.org/JPCC

Controlled Growth, Microstructure, and Properties of Functional Si Quantum Dot Films via Plasma Chemistry and Activated Radicals Bibhuti Bhusan Sahu,* Yongyi Yin, and Jeon Geon Han* School of Advanced Materials Science and Engineering, Center for Advanced Plasma Surface Technology (CAPST), NU-SKKU Joint Institute for Plasma Nano Materials, Sungkyunkwan University, Suwon, 440-746, Korea ABSTRACT: Control of plasma and radical generation and associated energy deposition near the growing thin films are still the main challenges in materials fabrication in the plasma-assisted deposition of Si quantum dot (QD) thin film. To control and enhance the material’s performance concerning film properties and application durability, we prepare 2.6 nm sized Si QDs with a fully ordered structure and entrapped them in amorphous silicon nitride using advanced dual frequency capacitively coupled plasmas. Raman and XRD analyses consistent with the high-resolution transmission electron micrographs reveal that the QD size can be controlled and altered from ∼2.6 to 4.0 nm simply by changing the operating pressure, which affects the film’s crystallinity in a broad range from 60% to 72% and the resulting microstructure. Further, a broad visible range ∼ 1.8−3.0 eV photoluminescence, with intense intensity and narrow to broad widths, is observed from Si QDs films. It is also seen that the observed photoluminescence featured is due to the quantum confinement effect within the QD material. Data reveal that the film properties are controllable by modifying a change in the plasma properties and radical parameters. The radio frequency and ultrahigh frequency dual frequency plasmas at low operating pressures have produced a very high atomic density of H and N radicals and a very high plasma density at low electron temperature, which are critically necessary and favorable to the control of film growth, nucleation, and other film properties. It is also seen that the deposition energy plays a significant role for the resulting microstructure and the QD size. The high luminescent yields in the visible range with a PL lifetime of ∼0.75 ns and size-tunable low-temperature deposition with plasma and radical control enable these QD materials as a good candidate for light emitting applications. Additionally, a plausible mechanism is foreseen for the QD film formation. of the QD size and film properties is the crucial goal during the deposition process to make a device like films. It may be of note that most of the existing literature is focused on the deposition of Si QDs embedded in silicon carbide (SiC), silicon oxide (SiO2), and silicon nitride (SiNx:H) via thermally activated phase separation (or high-temperature annealing) and precipitation of Si NCs in the dielectric matrix.5,8−10,12−17 Additionally, a few reports are present on the low-temperature deposition of Si QDs embedded in an amorphous SiNx:H matrix.8,9,12 Probably, this is because, compared with SiNx:H and SiO2 matrixes, the length and polarity of the Si−Si bond in the SiC matrix decrease, which brings about the complexity of phase separation and growth of Si nanocrystals in the SiC matrix.17,18 Compared to Si QDs embedded in SiO2 and SiNx:H matrixes, the Si QDs embedded in the SiC matrix have various advantages because of the lower barrier height of SiC (∼2.5 eV) relative to SiNx:H (∼4.2 eV) and SiO2 (∼7.9 eV).14,19 Some of the advantages may be listed as the formation of mini-bands between Si QDs, the higher Bloch carrier mobility, the increased tunneling probability (varies exponentially) between neighboring Si QDs, and

1. INTRODUCTION Silicon (Si) nanocrystals (NCs) and Si quantum dots (QDs) have recently shown the immense attention because of their compatibility with the Si technology that makes Si NCs and QDs very attractive for the fabrication of microelectronics, optoelectronics, and energy producing devices like solar cells.1−3 Particularly, Si QDs embedded in a dielectric matrix have emerged as the material for devices such as nextgeneration photovoltaic cells, thin-film transistors (TFTs), light emitting diodes (LEDs), and other applications.4−11 When the size of the Si QDs in the confined system approaches the Si Bohr diameter (∼10 nm), the overlapping of the electron−hole wave function is significantly enhanced, which, in turn, enhances the likelihood of radiative electron−hole recombination.5,7−9 Due to this reason, nanocrystalline Si-based nanostructured materials have a distinct advantage over bulk Si that exhibits an indirect band-gap semiconductor, which, in turn, necessitates phonon interactions while absorbing or emitting photons.8,9 Further, note that the band gap (Eg) of bulk Si is ∼1.12 eV, which can be much higher for Si QDs depending on the size of the QD. Because of QD’s enlarged Eg compared to the bulk Si, an intense visible photoluminescence (PL) can be possible at room temperature.8 Thus, by using properly sized Si QDs, it is possible to cover the maximum visible length of the electromagnetic spectrum.9 In this sense, simultaneous control © XXXX American Chemical Society

Received: March 15, 2017 Revised: April 18, 2017

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DOI: 10.1021/acs.jpcc.7b02430 J. Phys. Chem. C XXXX, XXX, XXX−XXX

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The Journal of Physical Chemistry C

Figure 1. (a) Side view of the DF PECVD reactor showing advanced RF and UHF DF plasma source used for this study. (b) Cross-sectional view of the experimental system showing various plasma diagnostic tools. (c) RF compensated LP used for the measurement of plasma parameters in RF environment. (d) VUVAS radical monitoring diagnostic for determining atomic H and N radicals. (e) Radical monitoring system for measuring the atomic radicals.

others.20−22 However, the deposition process of Si QDs in the SiC matrix suffers the limitation that it requires a very high deposition temperature along with the postannealing at a temperature ∼ thousand of degree. Nevertheless, SiNx:H is being utilized as a material with a higher dielectric constant and a greater strength over a wide temperature range than that of SiO2 and SiC along with a low tunneling barrier.18,23 Moreover, it has an excellent antireflection and surface passivation properties, which are very important for the device like solar cell’s performance.24,25 Additionally, for film growth and nanostructure control, there are numerous options of the assembly of nanoscale matters using various precursors in different states. In reality, one can arrange and assemble (e.g., by precipitating) solid nanoparticles using precursors in the gaseous, liquid, or plasma states. Further, depending on the specific way of organization, the atomic arrangement of nanoscale objects can also be different. It can result in differences in the structural and chemical properties of the objects produced. Typically, in plasma assisted deposition processes, the plasma produces an excess of different precursor species (e.g., atoms or atomic radicals, molecules, and molecular fragments) by electron impact, in a huge number of different energetic states (excited, metastable, ionized, and ground states) compared to solid, liquid, and gaseous precursors. Consequently, these plasma species/precursors are thereby very likely to congregate into solids with quite different microstructures that, in turn, will result in the different properties of these nanosolids. However, when dealing with nanomaterials like QDs, greater precision is

necessary, simply due to the key requirement to confine matters to nanometer dimensions. Therefore, plasma-produced species would lead to quite different outcomes in comparison to other environments. Plasma enhanced chemical vapor deposition (PECVD) by radio frequency (RF) capacitively coupled plasmas (CCP) is one of the most popular deposition methods for the preparation of these QD materials.8,9,15,26 For example, to synthesize the nanophase materials nanocrystalline Si (nc-Si) in a SiNx:H matrix, the plasma discharges are normally operated at a high power density (typically ≫ 100 mW cm−2) with the gas precursors highly diluted with hydrogen (e.g., the ratio, R = H2/ (SiH4 + NH3) > 10−100).26 Additionally, high-energetic ions produced under such adverse conditions quite often cause significant film degradation, particularly, by inducing defects, dislocations, and dangling bonds in the film.26,27 Further, an excessive dilution of the precursors with H2 and a hightemperature deposition give rise to low deposition rates (typically a few to tens of nm per min).8,9,28 However, the low-rate deposition and an elevated deposition temperature along with postdeposition annealing are certainly undesirable for the cost-effective QD synthesis and high-throughput fabrication of devices like films for solar cells, LEDs, TFTs, etc.29 Recently, nonequilibrium, low-pressure, high-density dual frequency (DF) CCP based processes have been very attractive and effective in the deposition of device-quality nc-Si and SiNx:H at low substrate temperatures and with high deposition rates.30−33 This is because, compared to commonly used CCPs B


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2.2. Plasma Chemistry Characterization. Figure 1b displays the schematic of the LP assembly, which is used to measure a current−voltage characteristic in the plasmas. Figure 1c shows the construction and components of the LP. A tungsten wire of diameter ∼ 0.5 mm, inserted into the stainless steel (ss) tube from one end of the capillary glass tube, constitutes the probe; about 2 mm of the wire protrudes from the glass capillary, to become the probe tip. The inner conductor of the coaxial line carried by the ss tube is connected to the tungsten tip at one end and the BNC connector at the other end to complete the electrical connection. To use the LP in the RF environment, an RF compensated LP35 is used. We apply a bias voltage to the probe from the power supply to get a probe current, and we acquire the relevant current−voltage characteristic in a computer by changing the bias voltages. Various parameters like plasma density or electron/ion density (n0), electron temperature (Te), and plasma potential (Vp) are then determined from the LP characteristic.36 2.3. Measurement of Deposition Energy on the Substrate. The energetic and thermal conditions at the substrate can be influenced by different plasma species like neutrals, electrons, ions, etc. These plasma activated species establish various elementary processes like chemical reactions, excitation, dissociation, ionization, radiation, recombination, etc. Thus, for plasma-based applications like thin-film deposition, the effect of energy per incoming plasma species/ particle to the substrate and the particle flux is crucial. Further, the product of these two factors constitutes the amount of deposited energy or energy flux to the substrate. This energy flux can be due to the contributions of the kinetic and potential energy of the particles along with the radiation from the adjoining environment plasma and the chamber wall. To measure the energy flux deposited on the substrate from the plasma, we have used the CP as the dummy substrate. Figure 1d presents the geometrical construction of the CP. The probe is basically a thin copper plate (of thickness ∼ 0.1 mm with a diameter of 20 mm) as the virtual substrate along with a thermocouple to measure the continuous change in its temperature. The measurement of the energy flux is based on the principle of the rate of change of the probe temperature. This variation is measured by taking the temperature characteristic while plasma is on (or during the heating process) and switched off (or during the cooling process).34 The details of the electronic circuitry, the data acquisition, and the analysis procedure can be found in the literature.34 2.4. Monitoring Atomic N and H Radicals. We have used the technique of vacuum ultraviolet absorption spectroscopy (VUVAS), as shown in Figure 1e, to monitor the atomic N and H radicals in the plasmas. The two basic components of the VUVAS system are a micro hollow cathode lamp as the light source and a vacuum ultraviolet (VUV) monochromator. These are symmetrically assembled about the plasma chamber. We generate a hollow-cathode discharge to form the VUV light in the light source using an on−off modulated dc power (∼10 Hz). The VUV light emerging from the light source is aligned and focused parallel through a MgF2 lens and is allowed to pass through the VUV monochromator through a MgF2 lens placed on another side of the chamber. Then, the VUV light is transformed to the visible light using the sodium salicylate scintillator equipped with the monochromator and captured by the photomultiplier tube (PMT). The intensity of the light signal is traced as a voltage signal in the oscilloscope. We track the intensity of the 121.1 and 120.6 nm VUV emission lines,

and inductively coupled plasmas (ICPs) operated at 13.56 MHz, the DF CCP source is quite capable of producing highdensity plasmas and atomic H and N radicals at low pressures and providing greatly controlled ionic species and neutral fluxes on the depositing substrate, which is necessary for the efficient crystallization of the microstructure network.12,14,33 In this contribution, we report the possibility of controlled nanophase growth in nc-Si/SiNx:H films, which leads to the development of self-assembled Si QDs embedded in the amorphous SiNx:H matrix at a very low temperature. Emphasis is put on controlling the plasma chemistry and energy of the plasma species for engineering the composition, structure, and QD size. The high growth rate deposition of self-assembled Si QDs is achieved using effective atomic H and N radicals and a high-density ions/electrons from nonequilibrium plasmas of SiH4 and NH3 operated with relatively low H2 dilution. Additionally, we have systematically and extensively investigated the effect of different H2 flow rates on the structural, compositional, and optical properties of the Si QD films. Further, a possible physical mechanism is proposed to elucidate the observed experimental results for a better understanding of film growth and associated film properties.

2. EXPERIMENTAL SECTION 2.1. Low-Temperature Growth of Si QDs in Advanced Dual Frequency CCP Reactor. The Si QD thin films are deposited in a nonthermal plasma reactor equipped with two capacitively coupled plasmas sources (CCPs): a radio frequency (RF) and ultra-high-frequency (UHF) plasma sources, operated at frequencies of 13.56 and 320 MHz, respectively. Figure 1a,b shows the sketches of the advanced DF PECVD system. For the RF CCP plasma, a commercial Dressler generator with an L-type matching network is used. The UHF source is a combination of two parallel copper tubes with an electrode spacing ≈ 4 cm, placed in between the two RF electrodes (with a spacing of ≈4 cm), with the ceramic feedthru inserted from the side port.30,31 A low power signal generator supplies a signal to a power amplifier to generate continuous-wave (CW) power up to 300 W. The output of the power amplifier is connected to the copper tubes (that acts as an antenna of length ∼ λ/2)34 through a coaxial cable. Here, λ ≈ 93.8 cm corresponds to the free space wavelength at the excitation frequency of 320 MHz. A total power ∼ 400 W with equal input powers of ∼200 W to the plasma via RF and UHF sources are continuously monitored by the RF generator and the power meter, respectively. In the present studies, the main gas precursors are silane (SiH4) and ammonia (NH3) diluted in H2. The flow rates of SiH4 and NH3 are, respectively, maintained at ∼4 and 1 sccm. We change the operating pressure in the reactor from 20 to 70 mTorr, by varying the gas flow of H2. The amount of H2 addition corresponds to the mixing ratio R = H2/(SiH4 + NH3) ∼ 2−14. The deposition temperature is maintained at 230 °C during the film growth. No postannealing process is employed after the deposition of Si QD films of thickness ∼ 1200 nm on the p-type Si(100) substrates, quartz, and SiO2 coated glasses. Additionally, Figure 1b shows the systematic and integrated diagnostic systems used to study the in situ plasma properties during the deposition. The various diagnostics used are Langmuir probe (LP), calorimetric probe (CP), optical emission spectroscopy (OES), and vacuum ultraviolet absorption spectroscopy (VUVAS). All the measurements are undertaken at the substrate location. C


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in the film. The Ti:Sapphire laser source generates an approximately 30 fs pulse at a repetition rate of ∼80 MHz. The laser is integrated with a Hamamatsu Streak camera for this measurement.

respectively, for the determination of H and N atom radical density. The details of the VUVAS system and the procedure for the determination of the density of N and H radicals are present in the literature.30,35 2.5. OES Diagnostics and Gas Temperature Measurements. In plasmas, neutral atoms and molecules are excited to their higher electronic states through collisions with electrons. The relaxation of these excited species to their lower energy states can be the source of light via emitted photons. We use an Acton Spectra Pro 500i Spectrometer to measure the optical emission intensities of various plasma species. The spectrometer has a resolution of ∼0.02 nm with an ∼10 μm wide entrance slit and a 1200 grooves mm−1 grating. The spectrometer is connected to a PIMAX ICCD camera (from Princeton Instrument) and interfaced with a computer (PC). We use window-based software “WinSpec32TM” to acquire the OES data. In the H2/NH3/SiH4 plasmas processes, there will be the creation of atomic H, atomic and molecular nitrogen, molecular H2, and other plasma and gas precursors species by the electron impact dissociation. Particularly, various excited lines from H (Hα, Hβ), H2 (Fulcher, GOBO), SiH, and N2, can be observed in the process plasma.12,30 Due to electron impact interactions, these species get heated, which can influence the thermal mobility in plasmas and, therefore, the gas temperature (Tg). Taking into account the degeneracy of the upper Fulcher energy level along with the transition strength, we express the line intensity to a relative population of the upper level.37,38 Then, from the slope of the variation of the relative population with the ground state energy (that gives a linear dependence), one can determine the value of Tg. The details of the Tg evaluation procedure can be found in ref 38. 2.6. Material Characterization. For systematic analysis and comparison, we consider the thickness of the deposited films to be constant ∼ 1200 nm. We use a JEOL-JSM2010 transmission electron microscope operating at 200 kV to acquire the high-resolution transmission electron (HRTEM) micrographs. The XRD measurements are performed using a Bruker-AXS Micro-diffractometer having a 2.2 kW sealed Cu Xray. The Raman spectra of the deposited films are acquired using the Renishaw inVia Micro-Raman spectrophotometer, which is equipped with a 514 nm Ar+ laser as the excitation source. The laser source is operated with a power density ∼ 2 mW cm−2 at the room temperature in a backscattering geometry. We use an ALPHA infrared spectrometer (Bruker Optics) to measure the Fourier transform infrared spectroscopy (FTIR) in a broad range of wavenumbers. We utilize the Al Kα (1486.6 eV) radiation as an X-ray source for the XPS measurements of the samples over a 4 mm × 5 mm area. The sample surfaces are sputter cleaned before the measurement using the ∼2 keV Ar ions to eliminate the contaminants. XPS survey spectra are measured at a pass energy of ∼20 eV with a step of 0.50 eV. Further, we use the binding energy (∼284.5 eV) of the C 1s peak to calibrate the possible energy shift because of accumulated charges. We measure the room temperature photoluminescence (PL) spectra by applying a He−Cd laser source at the ∼325 nm ultraviolet (UV) excitation wavelength. The spectra are recorded using a TRIAX 320 monochromator combined with the cooled Hamamatsu (R928) photomultiplier detector. Also, the PL lifetime is measured by utilizing a pulsed Ti:Sapphire laser (which has an excitation wavelength of 360 nm (energy ∼ 3.4 eV) at a laser power ∼ 100 mW) to check the efficacy of the optical transition

3. RESULTS AND DISCUSSION This section focuses on the detailed study of the possibility of controlled deposition of Si nanophase in the amorphous SiNx:H matrix by varying the working pressure by changing the gas flow rate of H2 in the precursor gas mixtures of SiH4 and NH3. A wide range of standard characterization tools including HRTEM, Raman, XRD, FTIR, XPS, and PL have been used for this purpose. Additionally, a detailed investigation by in situ monitoring of plasma characteristics, radicals, and energy deposition on the substrate by the plasma species is undertaken using various diagnostics for the plasma chemistry analysis. 3.1. Film Properties. To study the microstructure of the deposited QD films at the atomic scale, we perform the HRTEM measurements. Figure 2a,b shows typical HRTEM

Figure 2. HRTEM images of samples made at working pressures of (a) 30 mTorr and (b) 40 mTorr. (c) HRTEM image with high resolution of (a) showing crystalline Si QDs. The inset in the figure shows the ring patterns for the transmission electron diffraction from crystalline Si QDs. (d) The size distribution histogram of the Si QDs observed in (a) with imaging area ∼ 55 nm × 55 nm.

images of the samples deposited at pressures of 30 and 40 mTorr, respectively. Since the atomic density of Si is higher than that of the SiNx:H, the signature of black dots in the image is attributed to Si QDs. The HRTEM observation of the samples shows that the QD size of the sample prepared at a pressure of 40 mTorr (Figure 2b) is somewhat bigger than that prepared at 30 mTorr (Figure 2a). It is apparent from these figures that the Si QDs with a fairly uniform nanograin size are distributed over the whole scanned region. Figure 2c presents the HRTEM image with a higher resolution. Also, Figure 2c clearly shows lattice fringes represented by the shaded lines (enclosed by two yellow colored lines). Note that the QD size measured from this image approximately amounts to ∼2.6−2.7 nm. The inset, in Figure 2c, shows the ring patterns for the transmission electron diffraction from crystalline Si QDs in the D


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at a low pressure. To investigate further, the crystalline volume fraction (XC) of the films is determined using the method of deconvolution.31,39 Figure 4b presents a typical deconvolution of a Raman spectrum into three Gaussian components to determine the qualitative information, for the films prepared at 20 mTorr. These Gaussian peaks, as shown in Figure 4b, are, respectively, considered at ∼515−520, ∼510, and 480 cm−1 corresponding to the nc-Si, the ultra-nanocrystalline, and the amorphous (a-Si) components. The relevant portions of the individual components are then evaluated.31,39 The fraction XC is determined as XC = (I2 + I3)/(I1 + I2 + I3). Additionally, the ultra-nanocrystalline volume fraction (Xnc) is calculated from Xnc = I2/(I1 + I2 + I3). In these calculations, the quantities I1, I2, and I3 are, respectively, the integrated intensities of the components of peaks at wavenumbers ∼ 480, 510, and 515− 520 cm−1, respectively. The component sited in the range of ∼495−510 cm−1 is commonly considered as the grain boundary.31,39 Here, it is considered as the fraction Xnc with the signature of tiny crystallites with sizes much less than a nm surrounded by the grain boundaries. Figure 4c presents the overall crystallinity of the films prepared at various pressures. It is apparent from Figure 4c that XC slowly decreases from ∼72% to 68% as pressure increases from 20 to 50 mTorr, beyond which it slightly increases to ∼70% for the films prepared at a pressure of 60 mTorr and then, it drops sharply to ∼61% for the deposited films at 70 mTorr. Except for the films prepared at 40 mTorr, the tendency of variation of Xnc is very much similar to that of XC with a change in the range of 21−38% over the whole operating span of operating pressure from 20 to 70 mTorr. Note that Figure 4c also shows the plot of the intensity ratio r = I⟨220⟩/I⟨111⟩ at various pressures. Here, the parameters I⟨220⟩ and I⟨111⟩ represent the integrated intensities of the peaks relevant to the crystal planes ⟨220⟩ and ⟨111⟩ of the XRD spectra (Figure 3). The overall variation of this ratio gives the information on the change in the crystallographic texture of the films prepared as a function of working pressure. The XRD peaks (Figure 3) relevant to ⟨111⟩ orientations are, in general (except at a pressure of 70 mTorr), prominent. This feature indicates the preferential growth of Si crystallites along the ⟨111⟩ crystallographic orientation. However, careful observation of Figure 3 reveals that the relative intensity of the ⟨220⟩ peak, and hence, the intensity ratio r (Figure 4c), gradually changes with working pressure. The nature of variation of r is very much similar to that of XC. The sharp decrease in the value of r at a pressure of 70 mTorr suggests that the network gradually shifts from a highly nanocrystalline structure toward an amorphous dominated network where the ⟨311⟩ diffraction peak virtually disappears. The ⟨111⟩ orientation can arise from random nucleation while the ⟨220⟩ orientation is due to the growth of thermodynamically preferred grains.39 Also, the XRD spectra (Figure 3) with prominent (220) peaks indicate the formation of ⟨220⟩ oriented Si nanocrystallites within the SiNx:H matrix. Additionally, we correlate the HRTEM (Figure 2) and XRD (Figure 3) data with the Raman analysis (Figure 4a,c) to obtain the information about the QD structures. From the individual Raman spectrum, the average QD size is determined as dSi ≈ (1/3)·exp(−π2)/[(ωd − ω0)2 + (δ0/2)2].8,9 In this estimation, ωd represents the frequency of the nc-Si like mode for an NC of size dSi. The other parameters ω0 and δ0 are, respectively, 520 and 3.5 cm−1 for nc-Si. Figure 4d shows the tendency in the variation of the Si QD size qualitatively for the films deposited

SiNx:H matrix. Figure 2d illustrates the size distribution histogram of the QDs observed in Figure 2a. It is evident from the size distribution that the size of >65% of the Si nanocrystals falls within a narrow range ∼ 2−3 nm, which demonstrates the excellent uniformity of the QD in the entire scanned region. Additionally, the average size of the Si QDs determined from this histogram is 2.65 nm. Further, to study the crystal structure of the nanophase ncSi/SiNx:H films, we measure the XRD spectra of the samples deposited at various pressures (or gas mixing ratios, R) as shown in Figure 3. The figure shows various crystal planes of

Figure 3. XRD spectra of the samples prepared at different pressures at various H2 dilution ratios.

nc-Si within the films. The relevant peaks located at 2θ = 28.4, 47.3, and 56.2° are small and broad, representing the crystallographic planes ⟨111⟩, ⟨220⟩, and ⟨311⟩ of nc-Si, respectively.14,30 The occurrence of these three diffraction peaks indicates that the nano-Si crystallites are formed almost at all operating conditions (except 70 mTorr) of pressure. With the change in the working pressure, the intensities of these diffraction peaks are gradually changed while the full widths at half-maximum (FWHM) of all these peaks become smaller and higher. This observation suggests that the crystallite size turns out to be bigger and smaller, and accordingly, the degree of crystallinity is changed. Furthermore, it is worthwhile to mention that no other diffraction peaks attributed to crystalline Si3N4 are detected in the XRD spectra. The crystallinity of the deposited Si QD films is further examined by Raman spectroscopy. Figure 4 presents the Raman spectra of the films deposited on a quartz substrate relevant to the operating conditions of Figure 3. The dominant Raman peaks, in Figure 4a, are located between the wavenumbers ∼ 500 and 520 cm−1. The inset in the figure shows the signature of narrow and broad peaks with the change in the location of the peaks as the films prepared at different working pressures in between 20 and 70 mTorr at an interval of 10 mTorr. This feature of continuous variation indicates that there is a variation in the network structure that is linked with a concurrent change in the microstructure and, therefore, the size of the nano-Si crystallites. Figure 4a also shows that the Raman peak position significantly shifts toward lower wavenumbers, which is located at ∼503 cm−1 for the films deposited at 70 mTorr. This feature indicates that there is a significant change in the microstructure and the crystallinity of this film as compared to others prepared E


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Figure 4. (a) Measured Raman spectra of the Si QD films deposited at various working pressures. For clarity, the inset shows the Raman peak positions, which are labeled by the numbers 1−6. (b) Deconvolution of a typical Raman spectrum of the film prepared at 20 mTorr into various subcomponents. (c) Plot of overall variations of crystalline and ultra-nanocrystalline volume fractions of the films made at various working pressures. The figure also shows the plot of intensity ratio r = I⟨220⟩/I⟨111⟩ relevant to the crystal planes ⟨220⟩ and ⟨111⟩ of the XRD spectra (Figure 3). (d) Change in the QD size determined from Raman, XRD, and HRTEM images. Here, the size corresponds to the average QD size estimated from the Raman peak positions and FWHM of ⟨220⟩ XRD peaks. The figure also shows a similar variation of the ratio of crystalline and ultra-nanocrystalline volume fractions.

Figure 5. (a) FTIR spectra of QD films displaying typical stretching, bending, and wagging modes of amorphous SiNx:H. (b) Change of bonded N in SiNx:H and H content (in at. %), with pressures. (c) Si−H stretching mode absorption coefficient spectrum deconvoluted into components of Si− H monohydride and dihydride components and an intermediate bond-centered hydrogen, Si−H−Si, configuration. (d) Effect of plasma pressure on the microstructure factor and surface passivation index of the network.

F


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6.3 × 1018 cm−2 and NSi = 5 × 1022 cm−3 represent the oscillator strength, and the atomic density of nc-Si.40 Similarly, the bonded H content, f H (in at. %), is estimated from the corresponding absorption peaks relevant to Si−H wagging and N−H stretching modes as f H = Kω/NSi·[∫ (α(ω)/ω)·dω]· 100%.41,42 Here, Kω = 1.6 × 1019 cm−2 and 2.8 × 1020 cm−2 are, respectively, the oscillator strength of the Si−H wagging mode and N−H stretching mode.41,42 The overall variations of f N and f H of the deposited films are shown in Figure 5b at various pressures. The high and low values of f N and f H with changing pressures are due to their relative change in absorption coefficients of the Si−N asymmetric stretching mode and Si−H stretching mode and Si−H wagging mode. As the working pressure varies from 20 to 30 mTorr, f N increases from ∼0.78 to 0.88. It slowly falls to values ∼ 0.80−0.82 at 40−50 mTorr, and then it rapidly decreases to 0.63 with further increase in pressure up to 70 mTorr. Further, except at 50 mTorr, the overall variation of f H is very much similar to that of f N. It is apparent from Figure 5b that both the f N and f H values are simultaneously high at pressures of 30 and 50 mTorr. For these conditions, the QD size (Figure 4d) is much smaller compared to other conditions. Further, at 70 mTorr pressure, there is a significant reduction of f N and f H values, for which the XRD peaks of Si virtually disappear (Figure 3), and the film crystallinity XC (Figure 4c) is also very low. Further, careful inspection of Figures 5b and 4d reveal that the SiNx:H matrix simultaneously exhibiting lower contents of f N and f H has resulted in a relatively bigger QD size. We wish to note that, in the QD plasma processes involving H2/SiH4/NH3, there will be the production of high-density atomic H and N radicals, high-energy electrons, and ions owing to many electron impact reactions with the molecules and molecular fragments. Also, there will be frequent chemical reactions from the precursor gases during the film growth. Consequently, the film growth progression will undergo a twophase process, gas-phase and solid-phase reaction. For a better viewpoint, some of the important reactions can be listed as follows:

at various pressures. Also, the figure presents the QD size (represented by two stars in magenta color) measured from the HRTEM analyses for the comparison with the Raman data. The deconvolution of the line shape of the nc-Si components (Figure 4b) indicates a Gaussian-like distribution. The Si QD size, calculated from this deconvoluted peak positions (see Figure 4a as an example) of the Raman data, shows an overall variation of ∼2−3.2 nm at a pressure variation of 20−70 mTorr. Figure 4d also shows the ratio of Xnc to XC, whose variation is very similar to the variation of QD size. It can be seen from Figure 4c,d indication that the smaller the value of Xnc, the smaller the value of QD size. Additionally, we use the well-known Scherrer equation39 relevant to the ⟨220⟩ crystal plane to determine the nano grain (QD) sizes, which are also shown together in Figure 4d. It is evident that the overall trend of the grain size has a very similar and close variation to that of QD size estimated from Raman data. Further, a careful observation of Figures 2a−c reveals that each Si QD becomes dark and denser than the adjoining region of the amorphous SiNx:H matrix. Comparing our present result with the existing literature8,9,14,33 on Si QDs within an amorphous SiO2, SiC, and SiNx:H matrix, this adjacent layer can be treated as a shell covering the Si QD core. Also, this neighboring layer isolates QDs from the amorphous SiNx:H matrix where they are uniformly distributed and embedded. This amorphous shell performs as a protective layer for confining the nc-Si core and can play a vital role in controlling the QD size and its numbers (see Figure 2a,b). It is worth noting that the HRTEM measured QD sizes (Figure 2) for the samples prepared at 30 and 40 mTorr are in proximity to the values of ∼2.1 and 3.0 nm determined from the Raman analysis. Additionally, these values are consistent with the XRD analyses that determine the dot sizes ∼ 2.7 and 3.4 nm relevant to the peak of the (220) plane in Figure 3. Thus, all the Raman (Figure 4), HRTEM (Figure 2), and XRD (Figure 3) data are well correlated and fairly consistent. As discussed above, the microstructure and the size of QD can be varied by changing the operating conditions. This feature suggests that one can change the film deposition environment and, hence, the resulting film properties by controlling the plasma chemistry by changing the plasma conditions. The changes in plasma characteristics can induce an effect on the film composition or chemical properties as shown in Figure 5. Figure 5a shows the whole FTIR spectra of films prepared at different pressures in the wide wavenumber (ω) range = 600−3800 cm−1. The leading peaks in the spectra are identified as Si−H wagging (at ω ∼ 650 cm−1), Si−N stretching (ω ∼ 800 cm−1), Si−N asymmetric stretching (ω ∼ 875 cm−1), Si−N stretching of H-SiN3 (ω ∼ 1060 cm−1), N−H bending (ω ∼ 1220 cm−1), N−H2 scissors (ω ∼ 1380−1670 cm−1), Si− H stretching (ω ∼ 1900−2200 cm−1), and N−H stretching (ω ∼ 3300−3400 cm−1) modes of SiNx:H.8,32 Note that the crucial parameters of the SiNx:H film are the atomic H and N contents in the deposited films.32 To make a careful observation, insets in Figure 5a also show the sections of the Si−H stretching and N−H stretching modes. The figure shows that the intensities of the Si−N, N−H, and Si−H absorption peaks change progressively, suggesting the variation in the atomic N and H contents, in the network, those are incorporated into the SiNx:H matrix. To study further, the bonded N content, f N, in SiNx:H is calculated from the Si−N asymmetric stretching vibration (∼875 cm−1), as f N = Kω/NSi· [∫ (α(ω)/ω)·dω].32,40 In this estimation, the parameters Kω =

e− + SiH4 → SiH3 + H + e− → SiH 2 + 2H + e− → SiH4* + e−

SiH4* → SiH* + H 2 + H SiH4* → Si* + H 2 + 2H e− + H 2 → H + H + e−

H + SiH4 → SiH3 + H 2 H + NH3 → NH 2 + H 2 NH 2 + SiH3 → SiH 2(NH 2) + H

SiH 2(NH 2) + NH 2 → H + SiH(NH 2)2 SiH(NH 2)2 + NH 2 → H + Si(NH 2)3

Si(NH 2)3 + NH 2 → H + Si(NH 2)4 e− + NH3 → NH3+

e− + NH3 → NH + + H 2 → NH + + H + H− G


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Figure 6. Typical wide scans of XPS spectra of the samples deposited at different pressures. The inset shows the narrow scan XPS spectra for N 1s peaks and various contributions (at. %) in the films.

polyhydride component in the matrix, R = I2100/(I2000 + I2060 + I2100). Here, the factors I2100, I2060, and I2000 are, respectively, the integrated area of the components of peaks relevant to ω ∼ 2100, 2060, and 2000 cm−1. Additionally, the fraction of the Si−H−Si component is termed as the surface passivation index, S = I2060/(I2000 + I2060 + I2100). The variation of R and S are shown in Figure 5d. It is apparent from the Figure 5d that the sharp decrease in the microstructure factor (R), which essentially identifies the fast removal of structural imperfections from the network, corresponds to the rapid increase in the surface passivation index (S) of the matrix at pressure > 60 Torr, which indicates the transformation from highly crystalline toward the amorphous dominated network prepared at a relatively high pressure. The significant variation of R at high pressure is possibly due to the rapid change in the particle mobility diameter, which modifies the film surface of Hterminated Si nanoparticles.44,45 Note that the trend in the variation of R is similar to that of the ratio Xnc/XC size (Figure 4d) except for the films deposited at 30 mTorr. In this sense, the FTIR data (Figure 5) are also consistent with the Raman (Figure 4) and XRD (Figure 3) results. Further, to check the consistency of FTIR data, XPS analysis is performed for a few samples to investigate the elemental composition and bonding states of the films. Figure 6 presents the survey scan of XPS spectra of some samples prepared at pressures of 20−40 mTorr. The leading peaks occurring in the XPS spectra are from the elements Si and N. It seems that the residual surface contamination causes the trace of the elemental peaks of C and O at the chamber or the postdeposition absorbents on the sample surface. The existence of the Si−N bonds can be recognized from XPS spectra (Figure 6) of the Si (2p) and N (1s) peaks. For the Si (2p) spectra, the central peak located at ∼101.7 eV is recognizable to Si−N bonds, whereas, for the N (1s) spectra, the major peak centered at ∼398.8 eV refers to N−Si bonds.46 We further determine the quantitative information on the various elemental composition (X) of the deposited films using the formula as follows46

e− + NH+ → (NH )** → (NH+)* + e− → N + + H + e− → N + H + + e− → N2 + + H + 2e− SiHx + NHx → SiNx + H 2

SiHx + N → SiNx + H 2

In these above reactions, the symbol “*” represents the excited species. It is apparent that there are many sources for the generation of active H and N atomic radicals in the plasmas near the film growing surfaces that can affect the film properties. In this sense, the above-mentioned unusual variation of the f N and f H is attributed to the alteration of the Si−N, Si− H, and N−H bond concentrations in the SiNx:H network owing to the change in the atomic N and H radicals in the process plasmas with the progressive change in the operation conditions. This observed behavior indicates that a balance between atomic N and H atoms in the film is crucial for the formation of QDs in the SiNx:H matrix. In this sense, atomic N and H radical control during the deposition would be very much useful to provoke a change in the QD size. The relative contents of f N and f H in the film is attributed to the degree of dissociation of H2 and NH3 to produce crucial H and N radicals, and also, the ionization in the plasma that produces active H+ ions and electrons those are responsible for the chemical reactions to assist the film growth and nucleation. The role of plasma chemistry will be discussed later. Further, the absorption band in between 1900 and 2200 cm−1 can be deconvoluted to three satellite components corresponding to the monohydride SiH configuration at ω ∼ 2000 cm−1, the dihydride and polyhydride (SiH2)n mode at ω ∼ 2100 cm−1, and an intermediate bond-centered hydrogen, Si−H−Si, referred to be as the hydrides in a platelet-like configuration at ω ∼ 2060 cm−1 that is not usually observable in the case of nc-Si, as shown in Figure 5c.43 Figure 5c shows a representative plot relevant to the film prepared at 20 mTorr. The inset, in Figure 4a, shows that the relevant sections of Si− H stretching modes have broad and sharp absorption peaks. This modification can be attributed to the change in the film microstructure at different working pressures. The variation of the microstructure factor (R) is defined as the fraction of a

n

%X = 100·(AX /SX )/ ∑ (AM /SM ) M=1

(1)

where the parameters AX and AM, respectively, correspond to the areas under the peaks of the particular elements X and M in the XPS spectrum. The other parameters SX and SM in the H


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Figure 7. (a) Measured PL spectra of the deposited films under UV light excitations at room temperature. (b) The variations in PL peak energy and the optical band gap as a function of pressure. (c) The change of integrated area of PL intensities of the measured spectra along with the FWHM with pressure. (d) Normalized PL emission and absorbance with respect to the photon energy displaying emission and absorption. (e) A typical PL decay from the sample prepared at 30 mTorr showing the PL lifetime ∼ 0.75 ns. (f) PL peak energy for Si QDs as a function of dot size. Dotted lines are fitted curves for amorphous Si QDs and crystalline Si QDs. The filled squares are data points obtained for crystalline Si QDs satisfying the quantum confinement model represented by the solid line.

location of PL peak energy of the films prepared at various pressures is shown in Figure 7b. It is apparent from Figure 7b that the PL peak energy remains high (∼2.85 and 2.42 eV) for the working conditions relevant to pressures of 30 and 50 mTorr. Further, the PL peak energy does not change much and remains in between 1.72 and 2.01 eV for the other films prepared at other pressures. Further, Figure 7b also displays the variation of the optical band gap (Eg) of the material estimated from UV−vis spectroscopy of the film. The value of Eg of the deposited films is evaluated using the well-known Tauc’s equation by measuring the reflectance and absorbance in the UV−vis region.8,32 It is evident that the overall variation of Eg (∼1.8−3.05 eV) is very much similar and nearly in the same range to that of the PL peak energy. Additionally, H passivation of dangling bonds in nanoscale silicon is normally known to improve the PL intensity.48 Looking at the data relevant to atomic N and H contents in Figure 5b, one expects that surface passivation of the Si QDs could be the possibility for the observed high PL intensities. In particular, Figure 5b shows that the samples prepared at

denominator are the sensitivity factors of elements X and M, respectively. The sensitivity factors for O, Si, C, and N are 0.63, 0.17, 0.205, and 0.38, respectively.47 The N concentration is calculated from the narrow scan of the XPS spectra relevant to the N (1s) peaks centered at ∼398.8 eV (see the inset in Figure 6). Similar to the FTIR analysis (Figure 5b), the XPS result also confirms the fact that the N concentration increases from ∼17.5 to ∼24.7 at. % as the pressure changes from 20 to 30 mTorr, and then, it decreases to ∼19.7 at. % for the film deposited at 40 mTorr. We further investigate the optical properties to get the additional information. Figure 7a displays the room temperature PL spectra of the deposited films at various conditions of working pressure. The films exhibit very low to strong PL intensities in the visible wavelength range showing different colors upon ultraviolet (UV) laser light excitation. For the ease of reference, the inset, in Figure 7a, also shows the PL spectra as a function of photon energy. It is apparent from Figure 7a that the PL peak energy shifts from the violet (∼2.85 eV) toward red (∼1.72 eV) end of the visible spectrum. The specific I


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conduction band can easily go through a transition between these energy levels and can be delocalized. Further, the excited electron undergoes relaxation into the ground state with the release of light energy in the form of photons, which is known as recombination. Thus, this overall process of the photon mediated promotion and relaxation of electrons leads to the occurrence of the optical transitions. In this way, we take into account the optical excitation and recombination processes when electrons absorb and emit light. The time gap between the excitation and the recombination process is called as the optical recombination time that is referred to as the PL lifetime. Eventually, to ensure a recombination process, an excited electron should interact with a hole in the lower valence band. In a highly confined structure, like a Si QD, this can occur rapidly, and we expect the PL lifetime to be very small. The result of PL lifetime measurement of the film prepared at 30 mTorr is shown in Figure 7e. Several curves are plotted together in Figure 7e. The figure shows the relative comparison between the raw data (measured) on our QD sample and the impulse response function (IRF) of the measuring instrument. The figure also shows that the measured data are a convolution of the PL decay with the IRF. The deconvolution calculations are performed in Origin with the representative peak of the measurement. We subtract the baseline counts from the data for doing the deconvolution. Further, we perform the Fourier transform, a division of the IRF, and the inverse Fourier transform. Then, the deconvoluted measurement is plotted and fitted with a fairly accurate exponential curve (green line in Figure 7e) to find out the PL lifetime ∼ 0.75 ns. Note that a similar result of PL lifetime of Si QDs embedded in a SiO2 matrix has shown a very long lifetime from 6.3 μs to 83 ns.16 Further, we check the efficacy of the quantum confinement effect in the nanocrystalline Si QDs. We correlate the band-gap energy (which is nearly equal to the PL peak energy as shown in Figure 7b) and the size of crystalline Si QDs (Figure 4d), as determined by Raman, XRD, TEM, which are examined using a quantum confinement model.13,15 Taking into account the quantum confinement phenomena occurring in the nc-Si/ SiNx:H network of the Si−N−H complex system, we recall here the effective mass theory.15 According to the effective mass theory, the band-gap energy of a three-dimensionally confined Si QD can be expressed as

pressures of 30 and 50 mTorr have a high concentration of atomic H, which is conducive for the surface passivation. At other pressure conditions (20, 60, and 70 mTorr), surface passivation likely remains incomplete as the available atomic H radicals in the process plasmas are insufficient to satisfy the huge number of dangling bonds available in the film growing network. This aspect will be more cleared in light of the plasma diagnostic result to be discussed later. Further, it is apparent from Figures 7b and 4d that the nature of the variation of PL peak energy and Eg is very much opposite to that of the ratio Xnc/XC and the dot size. The higher value of PL intensity and peak energy for the films prepared at 30 and 50 mTorr corresponds to a low ratio of Xnc/XC ∼ 0.49 (Figure 4d) and a high microstructure factor (Figure 5d) with a higher fraction of atomic N and H (Figure 5b). This feature indicates that the network is dominated by very tiny Si nanocrystallites in the SiNx:H matrix. Along with this feature, the richness of Si (shown by the XPS analysis (Figure 6)) in the sample could be the source of the enhanced PL intensity with the high PL energy or violet shift of the PL band.32 Further, in a semiconductor material, electrons tend to make transitions near the edge of the band gap (Eg). Nevertheless, with QDs, the size of Eg is purely specified by adjusting the sizes of QDs. Further, the photoemission frequency of a QD semiconductor depends on the band gap Eg. In this sense, it is likely to control the output wavelength of the QD precisely. In effect, it is possible to tune the energy gap Eg of a QD, and, so, specify its “color”. It is apparent from Figure 7a,b that, at the pressure of ∼30 mTorr, there is the significant change in the plasma characteristic, which induces a foremost change in the deposition environment and, hence, the film properties. On the both sides of 30 mTorr, i.e., at low and high pressure, the change in the PL peak energy toward a much low value is caused by the change in the color of the spectrum from orange toward red. The distinctive shift in the PL peak energy corresponds to the change in the integrated area of PL intensity and the relevant FWHM as shown in Figure 7c. The overall shape of the curve of FWHM has a similar shape to that of PL peak energy (Figure 7b) with an overall variation from 0.2 to 0.6 eV. This narrow to a broad distribution of visible PL emissions (Figure 7a) with tunable QD size (Figure 4d) can be useful for the successful applications of the QD materials for the device like LEDs. Further, for a qualitative elucidation of the mechanisms of this observed PL along with the connection between the electronic structure and optical absorption of the QD semiconductor, the PL spectrum of the film deposited at 30 mTorr (Figure 7a) is shown in Figure 7d. In Figure 7d, both the optical emission and absorption spectra are plotted in an overlapping manner with the change in the photon energy. In Figure 7d, the PL is normalized to its maximum intensity since our purpose in this figure is to investigate the emitting− absorbing regions. Figure 7d shows that the absorption occurs at an energy ∼ 2.90 eV, and emission takes place at ∼2.85 eV. Additionally, the widths of the emission and absorption spectra are ∼0.50 and 0.45 eV, respectively, which are very close to each other. Note that, in a semiconductor, the band gap (Eg) separates the valence and conduction bands. Photons can introduce energy that can promote an electron to the upper conduction band. Further, the conduction band possesses many vacant energy levels. In this sense, an excited electron in the

E (eV) = E bulk +

C dSi 2

(2)

where Ebulk is the band gap of the bulk crystal silicon = 1.12 eV, C represents the quantum confinement parameter, and dSi corresponds to the diameter of the Si QDs (in nm). As shown in Figure 7f, a least-square fit to the data displays a close fit with the equation, E (eV) = 1.12 + 11.4/dSi2 for the band-gap energy of Si QDs in SiNx:H films deposited in H2/SiH4/NH3 plasmas. It is also reported that the connection between the band-gap energy and QD size could be fitted as E (eV) = 1.16 + 11.8/dSi2 in crystalline Si QDs embedded in a SiNx:H matrix grown by SiH4/N2 plasmas.49 Compared to this reported value,49 the fitted band-gap energy of 1.12 eV found in this work is in agreement with that of bulk crystal Si (∼1.12 eV), and the fitted parameter C = 11.4 is also larger, suggesting that quantum confinement is extremely increased in Si QDs, which are spontaneously grown in SiNx:H films using H2/SiH4/NH3 plasmas than samples grown in SiH4/N2 plasmas. The fitted bulk band-gap energy of 1.12 eV and the large value of the J


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Figure 8. (a) Typical OES spectrum measured at a pressure of 20 mTorr in the DF plasma process showing various emitted lines and their intensities. (b) Variation of gas temperature at various pressures. (c) Hα/SiH and Hβ/SiH intensity ratios showing the possible role of SiH radicals on the QD film properties. (d) Hα/N2 and Hβ/N2 intensity ratios indicating the effect of N2 excitation and their dissociation to form atomic N radical on the film properties. (e) Film growth rate measured at various pressures. (f) Measured plasma parameters, which show the plasma characteristics as a function of pressure. (g) The panel shows the variation of atomic N and H radical density relevant to the conditions of (f). (h) Measured energy deposition flux as a function of pressure. Particularly, the curve shows two minima corresponding to the QD films prepared at 30 and 50 mTorr, which produces the smaller-sized QDs.

is also reported to form energy levels in the band gap of Si QDs, which can obstruct the blue/violet shift of PL with the decreasing Si QD size.51 This feature can also cause electronic states in the Si QD with surface defects that are less sensitive to the quantum confinement. In the case of Si QDs that are spontaneously grown in SiNx:H using H2/SiH4/NH3 plasmas, there can be many resources (as discussed earlier in using chemical reactions), which would provide more H and N in the

quantum confinement parameter (=11.4) can be attributed to the enhanced crystallinity (Figure 4c) and the well-passivated surface of the Si QDs. In the case the crystallinity of Si QD is not high enough or perfect, that is, if the Si QD has a certain degree of structural disorder in the network, the electronic states would be somewhat localized and would be less responsive to the quantum confinement.50 Further, imperfect surface passivation K


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The Journal of Physical Chemistry C SiNx:H film using the RF/UHF dual frequency powers in the PECVD processes. The resulting H produced is believed to enhance the crystallinity of Si QDs via H-induced crystallization, where the presence of H leads to disorder-to-order transitions.52 Additionally, H dissociated from H2, SiH4, and NH3 would effectively passivate the surface of a Si QD. We attribute the formation of nanocrystalline Si QDs and the quantum confinement in the Si QDs to the high crystallinity and surface passivation by atomic H by significant dissociation of molecular gases using dual frequency power to be discussed below. 3.2. Effect of SiH Radicals, Plasma Parameters, Atomic H and N Radicals, and Energy Deposition on Growing Films. As demonstrated above, the film microstructure and associated properties depend crucially on the deposition environment of a given operation condition in PECVD deposition of Si QDs films. In this work, deposition is undertaken at low pressure and low substrate temperature. In the process plasmas, there will be the creation of various plasma species (especially SiH radicals, atomic H and N radicals, ions, and electrons) due to many chemical reactions and electron impact collisions among the atoms/molecules of the precursor gases. To study the effect of radicals, we acquire the plasma emission using the OES method from a viewport (Figure 1b) during the deposition process. Figure 8a presents a typical emission spectrum in the broad wavelength range ∼ 300−700 nm relevant to the pressure of 20 mTorr. The figure shows that the leading emission lines in the spectrum are from SiH radicals, molecular H2 (Fulcher and GoBo), molecular N2, and atomic H (Hα and Hβ).30,35 As explained earlier, for changing the deposition condition, the working pressure is sequentially changed by changing the H2 gas flow rate at fixed flow rates of SiH4 and NH3. Due to electron impact collision with the neutrals, the plasma generated species get heated and affect the thermal mobility in the process plasmas. Figure 8b shows the overall variation of gas temperature (Tg) as the pressure varies from 20 to 70 mTorr. It is apparent from the figure that Tg increases from ∼720 to 880 K as the pressure changes from 20 to 30 mTorr. It remains in between ∼750 and 800 K up to a pressure of 50 mTorr; then, it decreases further to ∼650 K at 60 mTorr, and saturates at ∼650−670 K at higher pressure (70 mTorr). Further, to get more insight of the deposition environment, we track the optical emission intensities of some excited lines relevant to the Hα, Hβ, SiH, and N2 (388.4 nm) excitations. The change in the ratios of Hα/SiH and Hβ/SiH intensities with the change in working pressure is shown in Figure 8c. It is apparent that the overall changes in the Hα/SiH and Hβ/SiH intensity ratios are similar but not systematic. Nevertheless, the comparison between the change of the Xnc/XC fraction and the QD size (Figure 4d) and Hα/SiH intensity ratio (Figure 8c) with pressure (except at 70 mTorr) indicates an opposite trend in the change of their values. Note that the high values of the intensity ratios at 30 mTorr and at 50 mTorr pressure indicate an enhancement in the line intensity of Hα relevant to atomic H and (or) a decrease in the value of SiH intensity. Additionally, it is possible that the decrease in the SiH radical intensity is owing to the creation of a smaller fraction of Si by the electron impact dissociation in the high-density plasmas, which is evident from the low value of Xnc (Figure 4c). In particular, the species SiH represents SiH4 derived species like SiH, SiH2, and SiH3. However, one cannot detect these species using OES since the energy necessary to generate a SiH radical is higher compared

to that of Si, SiH2, and SiH3.53 Further, the enhancement and reduction of the Hα/SiH and Hβ/SiH ratios would correspond to a decrease and increase in the intensity of SiH2, SiH3, and Si, respectively. This observed behavior gives a conjecture that SiH species derived from the molecular dissociation (SiH4) and the Hα/SiH ratio play a crucial role in controlling the network microstructure and QD size. Additionally, Figure 8d presents the variation of Hα/N2 and Hβ/N2 intensity ratios (relevant to the N2 (388.4 nm) emission line) at various pressures. The figure shows that the overall changes of the Hα/N2 and Hβ/N2 intensity ratios are very similar to each other. Careful inspection of the overall variation of the Hα/N2 and Hβ/N2 ratios (Figure 8d) and Xnc/XC fraction with the QD size (Figure 4d) reveals a similar trend in the change of their values. Further, the higher or lower value of the Hα/N2 and Hβ/N2 ratios depends on the relative intensity of Hα, Hβ, and N2 lines. Further, a higher emission intensity of N2 would indicate an enhancement of excited N2, which could dissociate to form more atomic N for the SiNx:H matrix formation. Probably due to this reason, the atomic N content obtained from the FTIR data, in Figure 5b, has an opposite nature of variation to that of Hα/N2 and Hβ/N2 ratios (Figure 8d). The above observations (Figure 8b−d) clearly indicate that the atomic H and N radicals and the plasma characteristics can play a vital role in the film growth and associated film properties. Figure 8e presents the variation of the deposition rate, which is similar to the change of the Tg (Figure 8b) and Hα/SiH ratio (except at a pressure of 70 mTorr) in Figure 8c. In particular, at a higher pressure ≥ 60 mTorr, the values of Tg and Hα/SiH are much lower due to the enhanced collision between the neutrals, which can severely affect the plasma parameters as discussed below. The clear features of radical and plasma formation are further investigated by the LP (Figure 8f) and VUVAS (Figure 8g) measurements, respectively. Figure 8f shows the variation of various plasma parameters at different working pressures. It is apparent from the Figure 8f that the electron temperature (Te) marginally varies in between ∼2.5 and 3.0 eV relevant to a pressure variation of 20−60 mTorr. Then, it rapidly falls to 1.7 eV at high pressure ≥ 60 mTorr. Relevant to the overall variation of Te ∼ 1.7−3.0 eV, the plasma density (n0) changes in between 3.6 × 1011 cm−3 and 1 × 1011 cm−3, for the pressure variation of 20−70 mTorr. Except at 30 and 50 mTorr, the profiles of Te and n0 are quite similar, indicating the electron impact ionization in the plasmas. The plasma potential (Vp) has a variation of ∼(82 ± 5 V) for the pressure variation of 20−70 mTorr. Additionally, the values Te ≪ Vp suggest the electrostatic confinement of the electrons in the plasmas. Further, Vp acquires a maximum value to compensate for the excess loss of electron (low n0 value) from the plasma to the chamber wall or ground at 70 mTorr. Further, Figure 8g presents the corresponding variation of atomic H density (nH) and N density (nN) at various pressures. It is apparent that, except at 70 mTorr, the profiles of nH is following the profile of plasma density (n0) (Figure 8f), indicating the plasma formation from the H radicals by the electronic impact. Further, except at a pressure of 70 mTorr, the overall variations of nH and nN are very much similar. Also, looking at the shape of Te (Figure 8f) curve, one may expect a monotonically decreasing curve of the same except at a pressure of 30 mTorr. However, at this pressure condition, there are enough atomic H atoms available (nH is maximum) for the L


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3.3. Mechanism in Light of the Observed QD Film Properties, Plasma Chemistry, and Energy Deposition. As discussed earlier, we change the deposition conditions by changing the working pressure using the sequential addition of H2 gas at fixed flow rates of SiH4 and NH3. We surmise here that the addition of the H2 to SiH4 and NH3 gases serves two key roles: it favors the creation of Si QDs by dropping the free energy (as evident from the low-energy flux in Figure 8h) of particular crystalline facets; it enhances and maintains the highdensity H radicals (nH) necessary for passivation (Figure 8g) and N content (Figure 6) of the growing films by effective ionization and dissociation by the DF plasma source. Let us now consider the effect of H2 dilution on the film growth (that includes the nucleation, the aggregation/coalescence, etc.) of the Si QDs embedded in the SiNx:H matrix. Taking into account the classic thermokinetic theory, we can express the Gibbs free energy for an arbitrary nanoparticle as14,54,55

plasma ionization (Figure 8g), which is apparent from the high value of n0 at a low Te. As explained earlier, owing to the high value of nH, one would anticipate here the role of H passivation of dangling bonds in nanoscale Si. Further, FTIR analysis (Figure 5b) shows that the profiles of atomic H and N in the film are approximately similar to that of measured profiles of nH and nN (Figure 8g). Additionally, one can see that the PL peak energy (Figure 7b) has a very similar profile to the profile of nN (Figure 8g). This correlation suggests that the band gap of the QD can be tuned by incorporating more atomic N (nH) in the network, which are controlled by the plasma parameters (Figure 8f) during the deposition process. The incorporation of a high fraction of atomic H and N in the growing film can be fulfilled by creating high values of nH and nN in the plasmas by a dedicated plasma source like RF/UHF DF source. The presence of high values of n0 (Figure 8f) can provide a high nH (Figure 8g) by electron impact dissociation, which can maintain the film crystallinity, and a simultaneous higher nN (Figure 8g) enables an N-rich SiNx:H matrix promoting a higher band gap with a smaller size QD. This typical feature is evident from the films prepared at working pressure conditions of 30 and 50 mTorr. Further, due to the formation and presence of a very high n0 (electrons and ions), nH (atomic H), and nN (atomic N), there is the likelihood of many heterogeneous chemical reactions among the plasma species, which are believed to be the feasible pathway for film growth. Additionally, the structural network is mostly created by the energy bombardment and relaxation processes of the plasma species and the adsorbed precursors on the surface, respectively. In this sense, the effect of energy per incoming plasma species or particle and, hence, the particle flux is decisive during the QD film growth. Particularly, the product of energy per incoming particle and particle flux defines the energy flux or the amount of energy deposition on the substrate. Further, there can be several contributions to the energy flux: the potential and kinetic energy of the energetic particles in the plasma apart from the radiation energy from the surrounding atmosphere like chamber walls and plasmas. Nevertheless, these individual contributions of energy are not simply isolated or measured in any given measurement. However, the CP can measure the effect of all contributions of energy, which would provide a better understanding of the impact of energy deposition on the substrate. Figure 8h presents the overall variation of measured energy flux on the substrate during the QD film formation. The figure shows that the profile of energy flux is approximately opposite to that of n0 (Figure 8f) and nN (Figure 8g). Also, it is apparent that the curve of the energy flux has two minima relevant to the operation conditions of the pressures of 30 and 50 mTorr. At other conditions, the energy flux is quite higher than their minimum value. In particular, at 70 mTorr, the energy flux is maximum, which is attributed to the higher ion energy bombardment (= eVp) by the plasma potential (Vp) in the plasma. Comparison of this energy deposition data (Figure 8h) with the film microstructure (Figure 4c) and QD size (Figure 4d), it is apparent that QDs with smaller dot size are observed for the conditions when the energy deposition on the substrate is the minimum. Further, for the condition of 70 mTorr, where the energy flux is very high, we do not observe the QD formation. This indicates that higher energy flux is not conducive for the network crystallinity (Figure 4c) and QD formation (Figure 4d).

G = Gnbulk + Gnsurface + Gnedge + Gncorner

(3)

where “G” represents the total Gibbs free energy (G) by the corresponds to the standard formation of a nucleus, and Gbulk n ” represents the free energy of the formation. The term “Gsurface n contribution to the Gibbs free energy to the surface, which is the sum of the product of two terms: the surface free energy of a facet and the relevant facet surface area. The term “Gedge n ” corresponds to the sum of the product of two terms: free energy to the edge and the lengths of the boundaries between ” represents the sum of the the two facets. The last term “Gcorner n free energy coupled with each corner.14 Further, in a bare Si substrate surface, there exist abundant Si dangling bonds. Meanwhile, in the plasma-based deposition with the H2 gas, a substantial fraction of Si dangling bonds is terminated by reactive atomic H atoms. For the Si−H bond, the surface free energy is ∼ −28 kcal/mol, which is much lower than that of the Si dangling bond (∼42 kcal/mol).56 These data indicate that the free energy G for the H-terminated surface is quite lower than that for the Si dangling bond surface using the above equation.14,55 Additionally, the lower G is predominantly promoting for the nucleation of Si QDs.54 Due to this reason, the degree of crystallinity of the nc-Si/SiNx:H films is maintained for the films prepared as the pressure increases from 20 to 60 mTorr. Further, simulation works have shown that the edges and corners of the nanocrystallites can make a significant contribution to G when a typical crystallite has atoms less than ∼104.54,57 Also, increased hydrogenation of the surface severely reduces the surface free energy on top of the surface of the Si QDs. Due to this reason, the net energy flux (Figure 8h) has two minima for the conditions of working pressure at 30 and 50 mTorr, where the nH is also quite high as compared to other conditions. Consequently, QDs with small sizes showing intense PL are observed for these conditions. We envisage further the effect of gas precursors on the depositing QD film. Note that the precursors are incorporated into the Si QD films by two ways: they can directly land on the film from the gas/plasma phase; they deposit on the substrate first and then travel to the crystallites through surface diffusion.58,59 As we increase the working pressure or H2 flow rate, the surface reactivity of the Si nuclei will be significantly reduced, which can lead to a lower surface diffusion activation energy (Ediff) along with a lower surface desorption energy (Edes).28,60 We can express the diffusion and desorption coefficients as follows M


Article

The Journal of Physical Chemistry C D=

a 2 −Ediff / kBT υe ; 4

μD = υe−Edes / kBT

UHF reactor at low temperature without postannealing. The reduction of the thermal resources by using low-temperature processes could enhance the device viability on inexpensive substrates. By varying the operating conditions by changing the H2 flow rate in a fixed background of SiH4 and NH3 gases, we critically monitor the plasma parameters, SiH radicals, atomic H and N radicals, and the deposition energy on the substrate by using numerous plasma diagnostics tools. The deposited films properties are systematically analyzed using standard analysis techniques, and the observed films properties are then correlated with the plasma chemistry to investigate the possible mechanism of QD film formation. Essentially, this works reports a detailed investigation of QD film processes design, which covers optimization of operation parameters, exploration of plasma chemistry, mechanism, and control of QD formation and associated film properties. Almost spherical Si QDs of average size ∼ 2.6 nm are observed from HRTEM analysis that is consistent with both XRD and Raman analyses. Raman and XRD data reveal that the Si QD size can be controlled and tuned from ∼2.6 to 4.0 nm by changing the plasma chemistry, which severely influences the network microstructure and the crystallinity in a wide range from ∼60% to 72%. Different film analyses reveal that the observed film properties can be tailored by inducing a change in the plasma characteristics, SiH, H and N radical parameters, and energy imparted to the substrate. The DF PECVD at low working pressures has shown a very high ionization and molecular dissociation capability to generate a very high density of H and N radicals with a low electron temperature, which are necessary and favorable for the film growth and nucleation. Notably, our experimental data clearly demonstrate efficient nanophase segregation with highly uniform Si QDs embedded in the amorphous SiNx:H matrix formed at a low pressure. Additionally, it is seen that the smallest-sized QD can be formed for the conditions when the deposition energy is the minimum. A feasible physical mechanism in light of the effect of H2 dilution on the nucleation and growth processes is proposed to explain the observed experimental results. It is also seen that the smallest QD size ∼ 2.6 nm diameter is formed, which has a PL lifetime of ∼0.75 ns when the energy deposition on the substrate is the minimum.

(4)

where the parameters a, υ, kB, and T, respectively, represent the lattice parameter, the attempt frequency, the Boltzmann constant, and the surface temperature, which is equivalent to the gas temperature Tg (Figure 8b) in the plasmas.28,58 As discussed above, the increase in H2 content to the plasma decreases Ediff, which would exponentially enhance D as can be seen from the above equation. Additionally, a higher T ∼ Tg (see Figure 8b for the case of 20−50 mTorr) will also induce enhancement of D. This behavior suggests that the adsorbed radicals on the film surface move much faster, and ultimately find the optimum stacking point in the crystal lattice. Additionally, the increase in H2 fraction also reduces Edes, which, in turn, enhances the rates of desorption as reactive H atoms make collisions with the Si adatoms on the surface and etch/remove them. Consequently, these effects strongly decrease the plasma activated species’ residence times on the surface. This characteristic reduces the rate of precursor incorporation into the growing QDs. As a result, the growth rates of the QD (Figure 8e) decreases at higher H2 dilution, particularly at a pressure ≥ 60 mTorr. Nevertheless, the higher mobility of adsorbed radicals owing to the reduction in Ediff results in pronounced crystallization as evident by the results as shown in Figures 2−4. Under such conditions, the plasma adsorbed species can move fast to the stacking points in the crystal lattice and are also removed quickly without being incorporated into the developing lattice.28 Further, one can foresee the role of atomic N in the confinement of Si QDs in the SiNx:H matrix. Relevant to the higher value of Eg (Figure 7b), one would expect enhanced nitrogenization (atomic N) in the film network. Notably, the profiles of the change of atomic N content of FTIR data (Figure 5b) and measured profile of nN (Figure 8g) favor this assertion. Also, the overall change of the PL peaks (Figure 7a) along with the relevant change in the Raman peak positions (Figure 4a) identifies the modification in the network microstructure and QD size. This overall feature further supported by XRD (Figure 3) and HRTEM studies (Figure 2) altogether indicates the feasibility of the evolution of PL (Figure 7a) from the quantum confinement effect (Figure 7f) within the material.15,49,50 We wish to note that this work demonstrates that the DF PECVD method is suitable for high ionization, excitation of plasmas species, and dissociation of molecular precursors in the plasmas. It is also revealed that the simultaneous existence of high atomic H radicals (nH), N radicals (nN), and high plasma density (n0) (Figure 8) are essential for the development of small-sized QDs (Figure 4d) at a high growth rate (Figure 8e). Further, there is also the requirement of low deposition energy (Figure 8h) on the substrate, which assists crystallization in the network. Additionally, we have stressed that the development progress made in the plasma-based film deposition and in situ plasma diagnostic techniques can be useful for a deeper understanding of Si QD process control and QD film growth.

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Bibhuti Bhusan Sahu: 0000-0002-1012-9559 Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the R&D Program of “Plasma Advanced Technology for Agriculture and Food (Plasma Farming)” via the National Fusion Research Institute of Korea (NFRI) funded by Government funds, the Global Development Research Center GRDC, a program of the Ministry of Science, ICT and Future Planning (MSIP, Grant

4. CONCLUSIONS In this contribution, the Si QDs films with intense visible PL emission from violet to red color tunable over a broad energy range ∼ 3.0−1.7 eV illustrate significant attention toward the fabrication of devices like LEDs. The QD films are prepared using single-step PECVD deposition in an advanced DF RF/ N


Article

The Journal of Physical Chemistry C No. 20120006672, second stage first year). It was also supported by the Korea Institute for Advancement of Technology (KIAT), which is funded by the Ministry of Trade, Industry & Energy (MOTIE, Grant No. N0000590).

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Controlled Growth, Microstructure, and Properties of Functional Si

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