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Crystallization mechanism of silicon quantum dots upon thermal annealing of hydrogenated amorphous Si-rich silicon carbide films
Thin Solid Films 552 (2014) 18–23
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Crystallization mechanism of silicon quantum dots upon thermal annealing of hydrogenated amorphous Si-rich silicon carbide films Guozhi Wen a,b, Xiangbin Zeng a,⁎, Wugang Liao a, Chenchen Cao a a b
School of Optical and Electronic Information, Huazhong University of Science and Technology, Wuhan, Hubei 430074, China School of Electronic and Electrical Engineering, Wuhan Polytechnic University, Wuhan, Hubei 430023, China
a r t i c l e
i n f o
Article history: Received 28 December 2012 Received in revised form 25 November 2013 Accepted 2 December 2013 Available online 8 December 2013 Keywords: Silicon Silicon carbide Quantum Dots Bonding Configuration Crystallization Mechanism Plasma-Enhanced Chemical Vapor Deposition Transmission Electron Microscopy
a b s t r a c t We have investigated the crystallization process of silicon quantum dots (QDs) imbedded in hydrogenated amorphous Si-rich silicon carbide (a-SiC:H) films. Analysis reveals that crystallization of silicon QDs upon thermal annealing of the samples can be explained in terms of bonding configuration and evolution of microstructure. The precursor gases were dissociated via electron impact reactions in the plasma-enhanced chemical vapor deposition, where the hydrogenated silicon radicals and reactive SiHn species lead to the formation of primary Si nuclei. With increasing annealing temperature, the breaking of SiHn bonds and decomposition of Si-rich SiC were progressively enhanced, allowing the formation of crystalline silicon QDs inside the a-SiC:H matrix. The results help clarify a probable mechanism for the growth of silicon QDs and provide the possibility to optimize the microstructure of silicon QDs in a-SiC:H films. © 2013 Elsevier B.V. All rights reserved.
1. Introduction As an unlimited and pollution-free energy source device, solar cells have been paid much attention recently. However, the reported efficiencies of conventional Si solar cells are lower than the ideal values [1–3]. The power-generating costs are, in this way, higher than that of thermal power generation systems or hydroelectric systems. To circumvent the Shockley–Queisser limit, multiple energy threshold approaches, such as all-silicon tandem cells using silicon QDs, are proposed [4,5]. There is a unique quantum size effect in QDs [6], which reveals that the bandgap of QDs can be properly tuned, depending on the dot dimension distribution. For all-silicon tandem solar cells using silicon QDs, the higherenergy photons can be absorbed at higher bandgaps, while the lowerenergy photons can be absorbed at the lower bandgaps. Thus silicon QDs, especially those with diameters less than the Bohr radius of exciton (~5 nm) in amorphous dielectric semiconductors, have been a subject of intensive research efforts [7–9]. Present efforts on silicon QDs are mainly allocated in preparation of silicon QDs embedded in a variety of semiconductor materials, such as non-stoichiometric silicon-oxide (SiOx), silicon-nitride (SiNx) and silicon-carbide (SiCx). The tunneling barrier for electrons and holes between adjacent silicon QDs in amorphous SiCx (2.5 eV) is lower than that in SiOx (9.0 eV) or SiNx (5.3 eV), where carriers can be easily transported and greater tunneling currents can be expected [5]. ⁎ Corresponding author. Tel./fax: +86 27 87544760. E-mail address: [email protected] (X. Zeng). 0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.tsf.2013.12.001
Amorphous Si-rich SiC films comprising silicon QDs are considered to be good candidates for the fabrication of future third generation Sibased photovoltaic devices [10]. Several different preparation techniques of silicon QDs in amorphous SiC matrix are under development. Considerable reports have been published about the formation of silicon QDs, as well as about exploration of their microstructure, photoluminescence and electrical properties [11]. In spite of the importance and availability of these materials, the key points for a full understanding about the structural evolution or crystallization process of silicon QDs upon thermal annealing of Si-rich a-SiC:H films have been less documented, and the discussion is still open [12–14]. 2. Experimental details Non-stoichiometric a-SiC:H samples were fabricated by decomposition of H2-diluted 10% silane (SiH4) and pure methane (CH4) gas mixtures on (100) crystal silicon wafers and quartz (SiO2) plates simultaneously. The radio frequency of the plasma-enhanced chemical vapor deposition (PECVD) system was 13.56 MHz. The vacuum chamber had a base pressure of less than 2.0 × 10−4 Pa. Deposition was carried out at a working pressure of 106.7 Pa. The power density and the growth temperature were fixed at 160 mW·cm−2 and 200 °C, respectively. The flow rate of SiH4 was maintained at 50 sccm and CH4 at 30 sccm. After deposition, the samples were cut into smaller parts. To explore the crystallization or growth process of silicon QDs in the SiC host matrix, each part was annealed subsequently in a quartz tube
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furnace with a 99.999% N2 atmosphere for 30 min at different temperatures. The chemical bonding configuration behaviors were deduced from Fourier transform infrared absorption measurements (FTIR, VERTEX 70). The chemical composition analyses were carried out with X-ray photoelectron spectroscopy (XPS, VG MultiLab 2000), using a monochromatic Al Kα (1486.5 eV) X-ray source and a hemispherical energy analyzer. The X-ray source power was 300 W. The instrument resolution was 0.47 eV. The working pressure was below 6.67 × 10−8 Pa. The analyzed area of the sample was about 0.36 mm2. Before detections, the samples were sputtered using a beam of 3 kV × 2 μA Ar+ bombardments for 180 s. The spectra were collected at 25 eV pass energy and the narrow-scan peaks were fitted with Thermo Avantage software. The binding energy values were calibrated using the contaminant carbon C1s = 284.6 eV. A standard Smart background was used for fitting the spectra. The crystalline phase and crystallinity were characterized by Raman scattering (HORIBA Jobin Yvon LabRAM Spectrometer HR 800 UV) in backscattering configuration. The laser light came from an Ar+ ion laser with a wavelength of 514.5 nm. The crystalline phase was further characterized by grazing incidence X-ray diffraction, using a Philips X'Pert Pro (XRD, PANalytical PW3040/60) at a voltage of 40 kV and a current of 40 mA, with Cu Kα radiation (λ = 1.540562 Å). The glancing angle between the incident X-ray and sample surface was 1.0°. A semiquantitative analysis of QD size was performed by X'Pert HighScore software. The direct observations of the microstructure were performed by transmission electron microscopy (TEM, JEM-2100F) and selected area electron diffraction (SAED, JEM-2010 HT) operated at 200 kV. The specimen for TEM and SAED analyses was reduced down to some tens of nanometers by a standard mechanical thinning technique and subsequent Ar+ ion milling polishing. 3. Results and discussion Direct physical evidence of the formation of silicon QDs in the nonstoichiometric a-SiC:H samples was displayed by a high resolution TEM. Fig. 1(a) shows a representative plan-view image of the sample annealed at 1050 °C, where a high density of isolated black spots could be identified, as well as coalescence of parts of the adjacent spots in a homogenous host matrix. Fig. 1(b) depicts the high resolution image of one perfect crystalline QD, where the lattice fringes could be
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observed clearly. The lattice spacing measured equals to approximately 3.17 Å. Fig. 1(c) displays the SAED pattern, where there are three diffraction rings and obvious diffuse aureoles. The diffraction rings reveal that the structure is polymorphous for the 1050 °C annealed nonstoichiometric SiC sample [15]. Radii of the rings are 3.12, 2.11 and 1.64 nm, respectively, which match well with the lattice parameters of the (111), (220) and (311) planes of c-Si [16]. Therefore, the black spots could be attributed to crystalline silicon QDs. The uniformity of the diffraction rings reveal that the number density of crystallized silicon QDs is very large, the size is very small and the orientations are random in the sample. The diffuse aureoles in vicinity of the 3.12 and 2.11 nm rings could be assigned to the presence of a large amount of amorphous silicon nanoclusters, whose average size is smaller than that of crystalline silicon QDs. At the same time, there are faint rings around 2.77 nm and several feeble bright spots dispersed in a circle with a radius of about 2.53 nm in the pattern, which could be ascribed to the amorphous carbon and crystalline SiC nanoparticles, respectively [17]. These feeble spots and faint rings reveal that only a small portion of crystalline SiC or carbon particles exist in the sample [18]. In order to investigate the crystallization process of silicon QDs, chemical bonding configuration behaviors of the as-grown samples and the samples annealed at temperatures of 750, 900, 1050 and 1200 °C were characterized by infrared absorption spectrometry. The absorption of substrate was eliminated by a bare silicon wafer. As Fig. 2 shows, two prominent absorption bands are observed for the asgrown sample. The prominent band around 2085 cm− 1 could be assigned to SiHn (n = 1,2,3) and/or C\SiH stretching vibration mode [19]. The other prominent band extending from 500 to 1200 cm− 1 could be ascribed to superposition of the following four absorption components: SiHn (n = 1,2,3) rocking and wagging mode near 655 cm−1 [20], Si\C non-hydrogen rocking and/or wagging mode or Si\CH3 stretching vibration mode near 780 cm−1 [21], SiH2 bending vibration mode at 800–900 cm−1 [22], CHn (n = 1,2,3) wagging or bending vibration mode near 1000 cm−1 [23]. When the annealing temperature is elevated from ambient circumstance to 750 °C, the SiHn and/or C\SiH mode around 2085 cm−1 disappear completely, the SiHn mode near 655 cm− 1 and CHn near 1000 cm−1 diminish and the Si–C or Si–CH3 mode near 780 cm−1 increases gradually. When the annealing temperature is further elevated to 1050 °C, the Si\C or Si\CH3 peak near 780 cm−1 is shifted toward
Fig. 1. (a) Typical plan-view high resolution TEM image of the amorphous non-stoichiometric SiC:H sample annealed at 1050 °C, (b) Image of one individual perfect crystalline silicon QD, and (c) Pattern of SAED.
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Fig. 2. Infrared absorption spectra for the as-grown sample and the samples annealed at temperatures of 750, 900, 1050 and 1200 °C, respectively. The spectrum was recorded in a range of 400–4000 cm−1 with resolution set at 4 cm−1.
a higher wave number of 828 cm− 1, the CHn peak near 1000 cm− 1 is shifted toward 1078 cm− 1, and both intensities are enhanced obviously. The above evolution of chemical bonding configurations induces the occurrence of phase separation and growth of Si QDs in its amorphous SiC host matrix. As it is well known, there are plenty of Si\Si, SiHn, C\SiH, Si\C and Si\CH3 radicals in the as-grown sample. When the annealing temperature increases to 750 °C, Si\H bonds are broken and hydrogen atoms are released from the SiHn and/or C\SiH entities, which is consistent with the disappearance of the mode around 2085 cm−1 and the reduction of the mode near 655 cm−1. When the annealing temperature reaches at 900 °C, the non-stoichiometric SiC radicals are dissociated as Si and SiC radicals: Si1 − xCx → x(SiC) + (1 − 2x)Si [24], in consistent with the enhancement of the Si\C vibration mode near 780 cm−1. These phenomena demonstrate that Si and SiC phase separation takes place at temperature as low as 900 °C. Silicon atoms with dangling bonds are separated out from SiHn and nonstoichiometric SiC, resulting in the increase in Si\Si and Si\C bond concentrations and the growth of silicon QDs in the amorphous host matrix. On the other hand, when the temperature increases to 1050 °C, the blue-shift of the peak at 782 cm−1 could be ascribed to passivation of some silicon atoms in the surface vicinity. Molecular N2 would be dissociated by absorbing the high thermal energy, a number of N atoms enter into the matrix to react with silicon dangling bonds: Si + N2 → SixNy [25]. The simultaneous blue-shift of the peak at 1000 cm−1 could be attributed to oxidation of some surface silicon atoms, where the oxygen atoms come from the ambient atmosphere [26]. It is well known that silicon crystallization process is related to elemental chemical composition ratios and bonding states of the samples before and after annealing. Besides the expected elemental Si and C, elemental O was also detected in each wide-scan spectrum. Elemental N was only found in the 1050 °C and 1200 °C annealed samples. Fig. 3 displays the narrow-scan XPS investigation of Si2p, C1s, N1s and O1s photoelectron emission peaks of the 1050 °C annealed sample. As Fig. 3 shows, the Si2p spectrum could be decomposed into four Gaussian components located at binding energies of 99.2, 99.7, 101.9 and 103.2 eV, which could be ascribed to Si\Si, Si\C, Si\N and Si\O bonds, respectively. The C1s spectrum could be decomposed into three components centered at 282.4, 284.3 and 285.7 eV, corresponding to C\Si, C\C and C\N bonding, respectively. The N1s spectrum could be decomposed into two components located at 397.6 and 398.5 eV, corresponding to N\Si and N\C bonding. The O1s spectrum could be decomposed into three components located at 530.8, 531.9 and 532.7 eV. The 530.8 eV bond and 531.9 eV bond could be ascribed to O\Si of SiOx (x b 2) while the 532.7 eV bond to SiO2 [27–29].
The quantitative compositions of the sample are estimated from the integrated intensities under the Si2p, C1s, N1s and O1s peaks. The atomic ratio of silicon to carbon (Si/C) could be calculated by the following equation: RSi / C = X / Y = AxSy / AySx, where X and Y represent Si and C, respectively. Ax is the area under the peak of Si2p and Ay is that of the C1s. Sx is the sensitivity factor of Si and Sy is that of C [30]. The area under the peak of Si2p is 41342.71 CPS·eV and that of C1s is 14618.46 CPS·eV. The sensitivity factors of Si, C, N, and O are 0.17, 0.205, 0.38, and 0.63, respectively [31]. From the data presented, we could roughly estimate the relative ratios of (Si\C)/(C\Si) = 3.41, (C\N)/(N\C) = 1.59, (Si\N)/ (N\Si) = 5.41, and (Si\O)/(O\Si) = 1.93. Therefore, we could conclude that the non-stoichiometric a-SiC:H sample is in Si-rich feature. The appearance of the peak around 99.2 eV indicates the precipitation of silicon nanoclusters, suggesting that silicon phase separation from amorphous SiC phase in the host matrix is taking place. The C\C bonding at 284.6 eV demonstrates that there are carbon particles in the film. Si\N at 101.9 eV and C\N at 285.7 eV appear as a result of the passivation of dangling bonds by N atoms aforementioned. By the same token, O\Si at 530.8, 531.9 and 532.7 eV appears as a result of the passivation of dangling bonds by O atoms. Fig. 4 manifests the Raman scattering measurements obtained from the as-grown sample and from the samples annealed at 900, 1050 and 1200 °C. Before measurement, the system was calibrated with a single crystal Si wafer, which had a clear peak at around 520 cm − 1. For the as-grown sample, the spectrum exhibits one broad peak at 480.9 cm− 1, which is due to scattering by the transverse optical (TO) mode of Si\Si and/or Si\C vibrations in the amorphous phase [26]. After the 900 °C annealing treatment, the spectrum features a pronounced peak at 515.2 cm− 1 and a highly asymmetrical low-energy tail extending down to 360 cm − 1 . The asymmetry to lower wavenumbers is an indication of the coexistence of amorphous and crystalline silicon phase in the matrix [18,32,33]. When the temperature is elevated to 1050 °C, the spectrum is characterized by a sharper asymmetrical peak, shifting to about 516.7 cm−1 with the full width at half maximum (FWHM) becoming narrower. Three models have been proposed for the growth of silicon nanoclusters embedded in the amorphous SiC matrix [34]. These models consider that atomic hydrogen dissociates the precursor gases via electron impact reactions. The abundant charged hydrogenated silicon radicals and reactive SiHn species lead to the formation of primary amorphous Si nuclei in the as-grown film, which is characterized by a broad Raman peak at 480 cm− 1. Thus, the appearance of a similar band indicates the existence of some Si clusters in the as-grown film. When the annealing temperature increases, the diffusion of the reactive precursors is enhanced [35]. The abundant atomic Si with dangling bonds would precipitate to form new nuclei in the disordered amorphous matrix. The high number density of new nuclei together with primary clusters would absorb other silicon atoms, diffusing from the SiC host matrix, to form amorphous silicon nanoclusters. In turn, certain amount of the amorphous silicon nanoclusters would grow to form crystalline silicon particles in randomly orientation. This gives place to the phase transition from amorphous Si-rich SiC:H to amorphous silicon clusters. Subsequently, the transition from amorphous silicon clusters to crystalline silicon clusters undergoes by increasing the thermal annealing temperature [32,36,37]. Ostwald ripening effect operates probably in the samples. Therefore, the amorphous silicon clusters are not crystallized in quantity [38]. The residual amorphous clusters and crystalline silicon clusters coexist in the amorphous SiC matrix. By taking into account the effect of both phonon confinement and elastic strain upon Raman scattering of small crystals, the crystallite size correspondent to special shift could be extracted from the measured profiles [39]. Compared with silicon monocrystalline, the shift of the sample annealed at 1050 °C may be caused by an average size of crystallite around 3.0–3.7 nm. The shift in wavenumbers from 515.2 to 516.7 cm− 1 suggests that the crystalline particle dimension increases. Furthermore, all the dominant Raman spectra are deconvoluted
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a) Si 2p
b) C1s
c) N1s
d) O1s
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Fig. 3. Narrow-scan survey spectra of (a) Si2p, (b) C1s, (c) N1s and (d) O1s of the Si-rich a-SiC:H sample annealed at 1050 °C (open circles), together with the Gaussian curves fitted (solid dots).
in the frequency range of 360–580 cm−1 using the least-squares routine to evaluate the degrees of crystallinity. A representative deconvolution of the spectrum is depicted in the inset of Fig. 4, where three independent Gaussian components could be obtained: a broad band near 474.9 cm−1 corresponding to amorphous Si or Si\C phase, a narrow band around 514.8 cm−1 corresponding to silicon crystal phase,
Fig. 4. Micro-Raman spectra measured on the as-grown sample and the samples annealed at temperatures of 900, 1050 and 1200 °C in a range of 100–1200 cm−1, respectively. Inset shows a representative deconvolution of the sample annealed at 900 °C in a range of 360–580 cm−1.
and an intermediate one in the vicinity of 504.7 cm−1 corresponding to the amorphous silicon clusters with smaller size. The crystal volume fraction Xc could be estimated from the sum of the integrated intensities of the individual crystalline and intermediate bands divided by the total sum of the integrated intensities of the individual crystalline, intermediate and amorphous bands [18,40]. It is found that the crystal fraction is about 40.1% for the 900 °C annealed sample and increases to 42.3% for the 1050 °C annealed one. With increasing annealing temperature, the nucleation process is enhanced and more Si atoms are involved in precipitation. This implies that high annealing temperatures can promote the crystallization of silicon and improve the crystalline fraction by the formation of more ordered structures in the amorphous phase. In addition, when the samples are annealed at temperatures above 900 °C, an increasing broad peak centered at about 956 cm−1 is observed. There are two components centered at 934.9 and 970.3 cm−1 when the spectrum is deconvoluted by the Gaussian function. Considering that amorphous SiC vibration density of states in Raman spectra is up to about 900 cm−1, the maximum optical phonon energy of any of the crystalline polytypes of SiC is about 972 cm−1 and one very feeble peak is observed at 796 cm−1. The origin of this Raman peak could be attributed to Si clusters (two-phonon process) with a small quantity of SiC particles [41]. Therefore, the Raman bands at 956 cm− 1 (see Fig. 4) could be explained by changes in the Si\C and Si\Si bonding states from amorphous to crystalline with increasing annealing temperature. The intensity of this peak increases with increasing Si and SiC nanoclusters volume fraction in the films [24]. The structural evolution of the samples was further investigated by grazing incident XRD. Fig. 5 illustrates XRD patterns of the same samples as those characterized by Raman scattering measurements. The
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4. Conclusions
Fig. 5. X-ray diffraction patterns of the as-grown sample and the samples annealed at temperatures of 900, 1050 and 1200 °C in a range of 20–80°, respectively.
Crystalline silicon QDs with average diameters of 3–4 nm are synthesized upon thermal annealing of hydrogenated Si-rich a-SiC:H samples. The growth process of crystalline silicon QDs is investigated through the evolution of bonding configurations and microstructure characterizations. The results reveal that the charged hydrogenated silicon radicals and reactive SiHn species lead to the formation of primary Si clusters in the as-deposited film. High temperature treatments could promote the separation of atomic Si from SiHn bonds and Si-rich SiC radicals. The abundant Si atoms, together with dangling bonds, would agglomerate to form new silicon nuclei and grow into amorphous clusters, even into crystal silicon QDs. Thus, a transition from amorphous to crystalline phase undergoes. With increasing thermal annealing temperature, the amorphous silicon and the crystalline silicon QDs coexist in the amorphous SiC matrix. Only a small portion of crystalline SiC or carbon particles are there in the sample. These findings help elucidate the probable crystallization mechanisms of silicon QDs and tailormake the growth parameters to determine more finely controllable QD size for the next generation photovoltaics. Acknowledgements
diffraction pattern was obtained by changing the position of the counter. Only two weak diffraction protuberances appeared at 2θ = 31.8° and 34.0° for the as-grown sample. After the 900 °C annealing treatment, in addition to the enhanced 31.8° and 34.0° peaks, other clearly resolved peaks were detected at 2θ = 28.4°, 36.9°, 47.4°, 54.9° and 56.3°, respectively. When the annealing temperature was further increased to 1050 °C, the peaks at 31.8°, 34.0°, 36.9° and 54.9° declined or disappeared and the amplitudes of the peaks at 28.4°, 47.4° and 56.3° increased obviously. Moreover, the amplitudes of these three peaks declined slightly again for the 1200 °C annealed sample. The three diffraction peaks at 28.4°, 47.4°, and 56.3° could be attributed to (111), (220) and (311) planes of crystalline silicon [16]. Silicon radicals are in amorphous nature in the as-grown film. The peak at 31.8° could be assigned to carbon nanoparticles [42] and the peaks at 34.0°, 36.9°, 54.9° to SiC nanoparticles [17]. When the annealing temperature increases to 900 °C, crystalline silicon QDs are produced and their sizes and degree of crystallinity are enhanced. Supposing that the crystalline silicon QDs have a spherical shape, the average size could be estimated by Scherrer's formula with the FWHM of the corresponding peaks [24,25]. With temperature increases from 900 to 1050 °C, the average size of silicon QDs increase from 3.1 to 3.6 nm for the (111) peak, and 3.4 to 4.0 nm for the (220) peak. These values are consistent with the Raman measurements. The decline of the diffraction peaks could be attributed to a composite effect of nitrogenation and oxidation of the silicon QDs at 1200 °C, which suppress the Si QDs from further growing or even cause a shrinkage of the silicon QD's domain. It is worthwhile to mention that diffraction peaks associated with SiC or carbon nanoparticles have been observed in all of the samples. The two feeble and broad peaks at 31.8° and 34.0° reveal that a few SiC and carbon clusters are synthesized in the as-deposited film. When the annealing temperature increases to 900 °C, the dissociation of Si-rich SiC radicals leads to the increase in Si\C bond concentrations and the rearrangement in the amorphous matrix. Therefore, parts of the amorphous carbon and SiC particles grow into crystallized ones. The feeble peaks seem to imply that there is only a small amount of crystalline SiC and carbon particles. The carbon atoms are mostly incorporated in the formation of Si\C and C\H bonds in the amorphous network [26]. When the annealing temperature increases further to 1050 °C, certain amount of N atoms is incorporated into the host matrix and reacts with the amorphous SiC and carbon particles. At the same time, the size of SiC clusters augment slightly, the increasing peaks of crystalline silicon particles suppress those of crystalline carbon and SiC particles [24]. This results in the decline or disappearance of the peaks of SiC and carbon particles.
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Crystallization mechanism of silicon quantum dots upon thermal annealing of hydrogenated amorphous Si-rich silicon carbide films Guozhi Wen a,b, Xiangbin Zeng a,⁎, Wugang Liao a, Chenchen Cao a a b
School of Optical and Electronic Information, Huazhong University of Science and Technology, Wuhan, Hubei 430074, China School of Electronic and Electrical Engineering, Wuhan Polytechnic University, Wuhan, Hubei 430023, China
a r t i c l e
i n f o
Article history: Received 28 December 2012 Received in revised form 25 November 2013 Accepted 2 December 2013 Available online 8 December 2013 Keywords: Silicon Silicon carbide Quantum Dots Bonding Configuration Crystallization Mechanism Plasma-Enhanced Chemical Vapor Deposition Transmission Electron Microscopy
a b s t r a c t We have investigated the crystallization process of silicon quantum dots (QDs) imbedded in hydrogenated amorphous Si-rich silicon carbide (a-SiC:H) films. Analysis reveals that crystallization of silicon QDs upon thermal annealing of the samples can be explained in terms of bonding configuration and evolution of microstructure. The precursor gases were dissociated via electron impact reactions in the plasma-enhanced chemical vapor deposition, where the hydrogenated silicon radicals and reactive SiHn species lead to the formation of primary Si nuclei. With increasing annealing temperature, the breaking of SiHn bonds and decomposition of Si-rich SiC were progressively enhanced, allowing the formation of crystalline silicon QDs inside the a-SiC:H matrix. The results help clarify a probable mechanism for the growth of silicon QDs and provide the possibility to optimize the microstructure of silicon QDs in a-SiC:H films. © 2013 Elsevier B.V. All rights reserved.
1. Introduction As an unlimited and pollution-free energy source device, solar cells have been paid much attention recently. However, the reported efficiencies of conventional Si solar cells are lower than the ideal values [1–3]. The power-generating costs are, in this way, higher than that of thermal power generation systems or hydroelectric systems. To circumvent the Shockley–Queisser limit, multiple energy threshold approaches, such as all-silicon tandem cells using silicon QDs, are proposed [4,5]. There is a unique quantum size effect in QDs [6], which reveals that the bandgap of QDs can be properly tuned, depending on the dot dimension distribution. For all-silicon tandem solar cells using silicon QDs, the higherenergy photons can be absorbed at higher bandgaps, while the lowerenergy photons can be absorbed at the lower bandgaps. Thus silicon QDs, especially those with diameters less than the Bohr radius of exciton (~5 nm) in amorphous dielectric semiconductors, have been a subject of intensive research efforts [7–9]. Present efforts on silicon QDs are mainly allocated in preparation of silicon QDs embedded in a variety of semiconductor materials, such as non-stoichiometric silicon-oxide (SiOx), silicon-nitride (SiNx) and silicon-carbide (SiCx). The tunneling barrier for electrons and holes between adjacent silicon QDs in amorphous SiCx (2.5 eV) is lower than that in SiOx (9.0 eV) or SiNx (5.3 eV), where carriers can be easily transported and greater tunneling currents can be expected [5]. ⁎ Corresponding author. Tel./fax: +86 27 87544760. E-mail address: [email protected] (X. Zeng). 0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.tsf.2013.12.001
Amorphous Si-rich SiC films comprising silicon QDs are considered to be good candidates for the fabrication of future third generation Sibased photovoltaic devices [10]. Several different preparation techniques of silicon QDs in amorphous SiC matrix are under development. Considerable reports have been published about the formation of silicon QDs, as well as about exploration of their microstructure, photoluminescence and electrical properties [11]. In spite of the importance and availability of these materials, the key points for a full understanding about the structural evolution or crystallization process of silicon QDs upon thermal annealing of Si-rich a-SiC:H films have been less documented, and the discussion is still open [12–14]. 2. Experimental details Non-stoichiometric a-SiC:H samples were fabricated by decomposition of H2-diluted 10% silane (SiH4) and pure methane (CH4) gas mixtures on (100) crystal silicon wafers and quartz (SiO2) plates simultaneously. The radio frequency of the plasma-enhanced chemical vapor deposition (PECVD) system was 13.56 MHz. The vacuum chamber had a base pressure of less than 2.0 × 10−4 Pa. Deposition was carried out at a working pressure of 106.7 Pa. The power density and the growth temperature were fixed at 160 mW·cm−2 and 200 °C, respectively. The flow rate of SiH4 was maintained at 50 sccm and CH4 at 30 sccm. After deposition, the samples were cut into smaller parts. To explore the crystallization or growth process of silicon QDs in the SiC host matrix, each part was annealed subsequently in a quartz tube
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furnace with a 99.999% N2 atmosphere for 30 min at different temperatures. The chemical bonding configuration behaviors were deduced from Fourier transform infrared absorption measurements (FTIR, VERTEX 70). The chemical composition analyses were carried out with X-ray photoelectron spectroscopy (XPS, VG MultiLab 2000), using a monochromatic Al Kα (1486.5 eV) X-ray source and a hemispherical energy analyzer. The X-ray source power was 300 W. The instrument resolution was 0.47 eV. The working pressure was below 6.67 × 10−8 Pa. The analyzed area of the sample was about 0.36 mm2. Before detections, the samples were sputtered using a beam of 3 kV × 2 μA Ar+ bombardments for 180 s. The spectra were collected at 25 eV pass energy and the narrow-scan peaks were fitted with Thermo Avantage software. The binding energy values were calibrated using the contaminant carbon C1s = 284.6 eV. A standard Smart background was used for fitting the spectra. The crystalline phase and crystallinity were characterized by Raman scattering (HORIBA Jobin Yvon LabRAM Spectrometer HR 800 UV) in backscattering configuration. The laser light came from an Ar+ ion laser with a wavelength of 514.5 nm. The crystalline phase was further characterized by grazing incidence X-ray diffraction, using a Philips X'Pert Pro (XRD, PANalytical PW3040/60) at a voltage of 40 kV and a current of 40 mA, with Cu Kα radiation (λ = 1.540562 Å). The glancing angle between the incident X-ray and sample surface was 1.0°. A semiquantitative analysis of QD size was performed by X'Pert HighScore software. The direct observations of the microstructure were performed by transmission electron microscopy (TEM, JEM-2100F) and selected area electron diffraction (SAED, JEM-2010 HT) operated at 200 kV. The specimen for TEM and SAED analyses was reduced down to some tens of nanometers by a standard mechanical thinning technique and subsequent Ar+ ion milling polishing. 3. Results and discussion Direct physical evidence of the formation of silicon QDs in the nonstoichiometric a-SiC:H samples was displayed by a high resolution TEM. Fig. 1(a) shows a representative plan-view image of the sample annealed at 1050 °C, where a high density of isolated black spots could be identified, as well as coalescence of parts of the adjacent spots in a homogenous host matrix. Fig. 1(b) depicts the high resolution image of one perfect crystalline QD, where the lattice fringes could be
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observed clearly. The lattice spacing measured equals to approximately 3.17 Å. Fig. 1(c) displays the SAED pattern, where there are three diffraction rings and obvious diffuse aureoles. The diffraction rings reveal that the structure is polymorphous for the 1050 °C annealed nonstoichiometric SiC sample [15]. Radii of the rings are 3.12, 2.11 and 1.64 nm, respectively, which match well with the lattice parameters of the (111), (220) and (311) planes of c-Si [16]. Therefore, the black spots could be attributed to crystalline silicon QDs. The uniformity of the diffraction rings reveal that the number density of crystallized silicon QDs is very large, the size is very small and the orientations are random in the sample. The diffuse aureoles in vicinity of the 3.12 and 2.11 nm rings could be assigned to the presence of a large amount of amorphous silicon nanoclusters, whose average size is smaller than that of crystalline silicon QDs. At the same time, there are faint rings around 2.77 nm and several feeble bright spots dispersed in a circle with a radius of about 2.53 nm in the pattern, which could be ascribed to the amorphous carbon and crystalline SiC nanoparticles, respectively [17]. These feeble spots and faint rings reveal that only a small portion of crystalline SiC or carbon particles exist in the sample [18]. In order to investigate the crystallization process of silicon QDs, chemical bonding configuration behaviors of the as-grown samples and the samples annealed at temperatures of 750, 900, 1050 and 1200 °C were characterized by infrared absorption spectrometry. The absorption of substrate was eliminated by a bare silicon wafer. As Fig. 2 shows, two prominent absorption bands are observed for the asgrown sample. The prominent band around 2085 cm− 1 could be assigned to SiHn (n = 1,2,3) and/or C\SiH stretching vibration mode [19]. The other prominent band extending from 500 to 1200 cm− 1 could be ascribed to superposition of the following four absorption components: SiHn (n = 1,2,3) rocking and wagging mode near 655 cm−1 [20], Si\C non-hydrogen rocking and/or wagging mode or Si\CH3 stretching vibration mode near 780 cm−1 [21], SiH2 bending vibration mode at 800–900 cm−1 [22], CHn (n = 1,2,3) wagging or bending vibration mode near 1000 cm−1 [23]. When the annealing temperature is elevated from ambient circumstance to 750 °C, the SiHn and/or C\SiH mode around 2085 cm−1 disappear completely, the SiHn mode near 655 cm− 1 and CHn near 1000 cm−1 diminish and the Si–C or Si–CH3 mode near 780 cm−1 increases gradually. When the annealing temperature is further elevated to 1050 °C, the Si\C or Si\CH3 peak near 780 cm−1 is shifted toward
Fig. 1. (a) Typical plan-view high resolution TEM image of the amorphous non-stoichiometric SiC:H sample annealed at 1050 °C, (b) Image of one individual perfect crystalline silicon QD, and (c) Pattern of SAED.
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Fig. 2. Infrared absorption spectra for the as-grown sample and the samples annealed at temperatures of 750, 900, 1050 and 1200 °C, respectively. The spectrum was recorded in a range of 400–4000 cm−1 with resolution set at 4 cm−1.
a higher wave number of 828 cm− 1, the CHn peak near 1000 cm− 1 is shifted toward 1078 cm− 1, and both intensities are enhanced obviously. The above evolution of chemical bonding configurations induces the occurrence of phase separation and growth of Si QDs in its amorphous SiC host matrix. As it is well known, there are plenty of Si\Si, SiHn, C\SiH, Si\C and Si\CH3 radicals in the as-grown sample. When the annealing temperature increases to 750 °C, Si\H bonds are broken and hydrogen atoms are released from the SiHn and/or C\SiH entities, which is consistent with the disappearance of the mode around 2085 cm−1 and the reduction of the mode near 655 cm−1. When the annealing temperature reaches at 900 °C, the non-stoichiometric SiC radicals are dissociated as Si and SiC radicals: Si1 − xCx → x(SiC) + (1 − 2x)Si [24], in consistent with the enhancement of the Si\C vibration mode near 780 cm−1. These phenomena demonstrate that Si and SiC phase separation takes place at temperature as low as 900 °C. Silicon atoms with dangling bonds are separated out from SiHn and nonstoichiometric SiC, resulting in the increase in Si\Si and Si\C bond concentrations and the growth of silicon QDs in the amorphous host matrix. On the other hand, when the temperature increases to 1050 °C, the blue-shift of the peak at 782 cm−1 could be ascribed to passivation of some silicon atoms in the surface vicinity. Molecular N2 would be dissociated by absorbing the high thermal energy, a number of N atoms enter into the matrix to react with silicon dangling bonds: Si + N2 → SixNy [25]. The simultaneous blue-shift of the peak at 1000 cm−1 could be attributed to oxidation of some surface silicon atoms, where the oxygen atoms come from the ambient atmosphere [26]. It is well known that silicon crystallization process is related to elemental chemical composition ratios and bonding states of the samples before and after annealing. Besides the expected elemental Si and C, elemental O was also detected in each wide-scan spectrum. Elemental N was only found in the 1050 °C and 1200 °C annealed samples. Fig. 3 displays the narrow-scan XPS investigation of Si2p, C1s, N1s and O1s photoelectron emission peaks of the 1050 °C annealed sample. As Fig. 3 shows, the Si2p spectrum could be decomposed into four Gaussian components located at binding energies of 99.2, 99.7, 101.9 and 103.2 eV, which could be ascribed to Si\Si, Si\C, Si\N and Si\O bonds, respectively. The C1s spectrum could be decomposed into three components centered at 282.4, 284.3 and 285.7 eV, corresponding to C\Si, C\C and C\N bonding, respectively. The N1s spectrum could be decomposed into two components located at 397.6 and 398.5 eV, corresponding to N\Si and N\C bonding. The O1s spectrum could be decomposed into three components located at 530.8, 531.9 and 532.7 eV. The 530.8 eV bond and 531.9 eV bond could be ascribed to O\Si of SiOx (x b 2) while the 532.7 eV bond to SiO2 [27–29].
The quantitative compositions of the sample are estimated from the integrated intensities under the Si2p, C1s, N1s and O1s peaks. The atomic ratio of silicon to carbon (Si/C) could be calculated by the following equation: RSi / C = X / Y = AxSy / AySx, where X and Y represent Si and C, respectively. Ax is the area under the peak of Si2p and Ay is that of the C1s. Sx is the sensitivity factor of Si and Sy is that of C [30]. The area under the peak of Si2p is 41342.71 CPS·eV and that of C1s is 14618.46 CPS·eV. The sensitivity factors of Si, C, N, and O are 0.17, 0.205, 0.38, and 0.63, respectively [31]. From the data presented, we could roughly estimate the relative ratios of (Si\C)/(C\Si) = 3.41, (C\N)/(N\C) = 1.59, (Si\N)/ (N\Si) = 5.41, and (Si\O)/(O\Si) = 1.93. Therefore, we could conclude that the non-stoichiometric a-SiC:H sample is in Si-rich feature. The appearance of the peak around 99.2 eV indicates the precipitation of silicon nanoclusters, suggesting that silicon phase separation from amorphous SiC phase in the host matrix is taking place. The C\C bonding at 284.6 eV demonstrates that there are carbon particles in the film. Si\N at 101.9 eV and C\N at 285.7 eV appear as a result of the passivation of dangling bonds by N atoms aforementioned. By the same token, O\Si at 530.8, 531.9 and 532.7 eV appears as a result of the passivation of dangling bonds by O atoms. Fig. 4 manifests the Raman scattering measurements obtained from the as-grown sample and from the samples annealed at 900, 1050 and 1200 °C. Before measurement, the system was calibrated with a single crystal Si wafer, which had a clear peak at around 520 cm − 1. For the as-grown sample, the spectrum exhibits one broad peak at 480.9 cm− 1, which is due to scattering by the transverse optical (TO) mode of Si\Si and/or Si\C vibrations in the amorphous phase [26]. After the 900 °C annealing treatment, the spectrum features a pronounced peak at 515.2 cm− 1 and a highly asymmetrical low-energy tail extending down to 360 cm − 1 . The asymmetry to lower wavenumbers is an indication of the coexistence of amorphous and crystalline silicon phase in the matrix [18,32,33]. When the temperature is elevated to 1050 °C, the spectrum is characterized by a sharper asymmetrical peak, shifting to about 516.7 cm−1 with the full width at half maximum (FWHM) becoming narrower. Three models have been proposed for the growth of silicon nanoclusters embedded in the amorphous SiC matrix [34]. These models consider that atomic hydrogen dissociates the precursor gases via electron impact reactions. The abundant charged hydrogenated silicon radicals and reactive SiHn species lead to the formation of primary amorphous Si nuclei in the as-grown film, which is characterized by a broad Raman peak at 480 cm− 1. Thus, the appearance of a similar band indicates the existence of some Si clusters in the as-grown film. When the annealing temperature increases, the diffusion of the reactive precursors is enhanced [35]. The abundant atomic Si with dangling bonds would precipitate to form new nuclei in the disordered amorphous matrix. The high number density of new nuclei together with primary clusters would absorb other silicon atoms, diffusing from the SiC host matrix, to form amorphous silicon nanoclusters. In turn, certain amount of the amorphous silicon nanoclusters would grow to form crystalline silicon particles in randomly orientation. This gives place to the phase transition from amorphous Si-rich SiC:H to amorphous silicon clusters. Subsequently, the transition from amorphous silicon clusters to crystalline silicon clusters undergoes by increasing the thermal annealing temperature [32,36,37]. Ostwald ripening effect operates probably in the samples. Therefore, the amorphous silicon clusters are not crystallized in quantity [38]. The residual amorphous clusters and crystalline silicon clusters coexist in the amorphous SiC matrix. By taking into account the effect of both phonon confinement and elastic strain upon Raman scattering of small crystals, the crystallite size correspondent to special shift could be extracted from the measured profiles [39]. Compared with silicon monocrystalline, the shift of the sample annealed at 1050 °C may be caused by an average size of crystallite around 3.0–3.7 nm. The shift in wavenumbers from 515.2 to 516.7 cm− 1 suggests that the crystalline particle dimension increases. Furthermore, all the dominant Raman spectra are deconvoluted
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a) Si 2p
b) C1s
c) N1s
d) O1s
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Fig. 3. Narrow-scan survey spectra of (a) Si2p, (b) C1s, (c) N1s and (d) O1s of the Si-rich a-SiC:H sample annealed at 1050 °C (open circles), together with the Gaussian curves fitted (solid dots).
in the frequency range of 360–580 cm−1 using the least-squares routine to evaluate the degrees of crystallinity. A representative deconvolution of the spectrum is depicted in the inset of Fig. 4, where three independent Gaussian components could be obtained: a broad band near 474.9 cm−1 corresponding to amorphous Si or Si\C phase, a narrow band around 514.8 cm−1 corresponding to silicon crystal phase,
Fig. 4. Micro-Raman spectra measured on the as-grown sample and the samples annealed at temperatures of 900, 1050 and 1200 °C in a range of 100–1200 cm−1, respectively. Inset shows a representative deconvolution of the sample annealed at 900 °C in a range of 360–580 cm−1.
and an intermediate one in the vicinity of 504.7 cm−1 corresponding to the amorphous silicon clusters with smaller size. The crystal volume fraction Xc could be estimated from the sum of the integrated intensities of the individual crystalline and intermediate bands divided by the total sum of the integrated intensities of the individual crystalline, intermediate and amorphous bands [18,40]. It is found that the crystal fraction is about 40.1% for the 900 °C annealed sample and increases to 42.3% for the 1050 °C annealed one. With increasing annealing temperature, the nucleation process is enhanced and more Si atoms are involved in precipitation. This implies that high annealing temperatures can promote the crystallization of silicon and improve the crystalline fraction by the formation of more ordered structures in the amorphous phase. In addition, when the samples are annealed at temperatures above 900 °C, an increasing broad peak centered at about 956 cm−1 is observed. There are two components centered at 934.9 and 970.3 cm−1 when the spectrum is deconvoluted by the Gaussian function. Considering that amorphous SiC vibration density of states in Raman spectra is up to about 900 cm−1, the maximum optical phonon energy of any of the crystalline polytypes of SiC is about 972 cm−1 and one very feeble peak is observed at 796 cm−1. The origin of this Raman peak could be attributed to Si clusters (two-phonon process) with a small quantity of SiC particles [41]. Therefore, the Raman bands at 956 cm− 1 (see Fig. 4) could be explained by changes in the Si\C and Si\Si bonding states from amorphous to crystalline with increasing annealing temperature. The intensity of this peak increases with increasing Si and SiC nanoclusters volume fraction in the films [24]. The structural evolution of the samples was further investigated by grazing incident XRD. Fig. 5 illustrates XRD patterns of the same samples as those characterized by Raman scattering measurements. The
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4. Conclusions
Fig. 5. X-ray diffraction patterns of the as-grown sample and the samples annealed at temperatures of 900, 1050 and 1200 °C in a range of 20–80°, respectively.
Crystalline silicon QDs with average diameters of 3–4 nm are synthesized upon thermal annealing of hydrogenated Si-rich a-SiC:H samples. The growth process of crystalline silicon QDs is investigated through the evolution of bonding configurations and microstructure characterizations. The results reveal that the charged hydrogenated silicon radicals and reactive SiHn species lead to the formation of primary Si clusters in the as-deposited film. High temperature treatments could promote the separation of atomic Si from SiHn bonds and Si-rich SiC radicals. The abundant Si atoms, together with dangling bonds, would agglomerate to form new silicon nuclei and grow into amorphous clusters, even into crystal silicon QDs. Thus, a transition from amorphous to crystalline phase undergoes. With increasing thermal annealing temperature, the amorphous silicon and the crystalline silicon QDs coexist in the amorphous SiC matrix. Only a small portion of crystalline SiC or carbon particles are there in the sample. These findings help elucidate the probable crystallization mechanisms of silicon QDs and tailormake the growth parameters to determine more finely controllable QD size for the next generation photovoltaics. Acknowledgements
diffraction pattern was obtained by changing the position of the counter. Only two weak diffraction protuberances appeared at 2θ = 31.8° and 34.0° for the as-grown sample. After the 900 °C annealing treatment, in addition to the enhanced 31.8° and 34.0° peaks, other clearly resolved peaks were detected at 2θ = 28.4°, 36.9°, 47.4°, 54.9° and 56.3°, respectively. When the annealing temperature was further increased to 1050 °C, the peaks at 31.8°, 34.0°, 36.9° and 54.9° declined or disappeared and the amplitudes of the peaks at 28.4°, 47.4° and 56.3° increased obviously. Moreover, the amplitudes of these three peaks declined slightly again for the 1200 °C annealed sample. The three diffraction peaks at 28.4°, 47.4°, and 56.3° could be attributed to (111), (220) and (311) planes of crystalline silicon [16]. Silicon radicals are in amorphous nature in the as-grown film. The peak at 31.8° could be assigned to carbon nanoparticles [42] and the peaks at 34.0°, 36.9°, 54.9° to SiC nanoparticles [17]. When the annealing temperature increases to 900 °C, crystalline silicon QDs are produced and their sizes and degree of crystallinity are enhanced. Supposing that the crystalline silicon QDs have a spherical shape, the average size could be estimated by Scherrer's formula with the FWHM of the corresponding peaks [24,25]. With temperature increases from 900 to 1050 °C, the average size of silicon QDs increase from 3.1 to 3.6 nm for the (111) peak, and 3.4 to 4.0 nm for the (220) peak. These values are consistent with the Raman measurements. The decline of the diffraction peaks could be attributed to a composite effect of nitrogenation and oxidation of the silicon QDs at 1200 °C, which suppress the Si QDs from further growing or even cause a shrinkage of the silicon QD's domain. It is worthwhile to mention that diffraction peaks associated with SiC or carbon nanoparticles have been observed in all of the samples. The two feeble and broad peaks at 31.8° and 34.0° reveal that a few SiC and carbon clusters are synthesized in the as-deposited film. When the annealing temperature increases to 900 °C, the dissociation of Si-rich SiC radicals leads to the increase in Si\C bond concentrations and the rearrangement in the amorphous matrix. Therefore, parts of the amorphous carbon and SiC particles grow into crystallized ones. The feeble peaks seem to imply that there is only a small amount of crystalline SiC and carbon particles. The carbon atoms are mostly incorporated in the formation of Si\C and C\H bonds in the amorphous network [26]. When the annealing temperature increases further to 1050 °C, certain amount of N atoms is incorporated into the host matrix and reacts with the amorphous SiC and carbon particles. At the same time, the size of SiC clusters augment slightly, the increasing peaks of crystalline silicon particles suppress those of crystalline carbon and SiC particles [24]. This results in the decline or disappearance of the peaks of SiC and carbon particles.
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