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Synthesis of heterostructured SiC and C–SiC nanotubes by ion irradiation-induced changes in crystallinity
Carbon 95 (2015) 279e285
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Synthesis of heterostructured SiC and CeSiC nanotubes by ion irradiation-induced changes in crystallinity Tomitsugu Taguchi*, Shunya Yamamoto, Katsuaki Kodama, Hidehito Asaoka Quantum Beam Science Center, Japan Atomic Energy Agency, Tokai-mura, Ibaraki-ken 319-1195 Japan
a r t i c l e i n f o
a b s t r a c t
Article history: Received 1 May 2015 Received in revised form 17 August 2015 Accepted 18 August 2015 Available online 20 August 2015
Amorphous SiC nanotubes are successfully synthesized for the first time by the irradiation of polycrystalline SiC nanotubes with 340 keV Siþ ions. A polycrystalline/amorphous heterostructured SiC nanotube, in which polycrystalline SiC and amorphous SiC coexist in the same nanotube, is also synthesized by ion irradiation with a mask in front of the polycrystalline SiC nanotube. According to evaluation by electron energy loss spectroscopy, the plasmon energies of the SiC nanotube change rapidly at the interface between the polycrystalline and amorphous regions. The volume swelling induced by amorphization, as evaluated from the differences in plasmon energies, is approximately 5.0%. This result reveals that more relaxed amorphous SiC nanotubes with higher densities can be produced. The graphitic shells in the carbon layer of CeSiC nanotubes are found to gradually bend to align with the radial directions of the nanotubes by ion irradiation. Since graphite (002) spots in the selected-area electron diffraction pattern are clearly observed even after ion irradiation, the carbon layer in the CeSiC nanotube maintains crystallinity. Moreover, a new multi-walled carbon nanotube with graphitic shells completely parallel to the radial direction of the nanotube is also produced inside the amorphous SiC tubular layer, in the case of large-caliber CeSiC nanotubes. © 2015 Elsevier Ltd. All rights reserved.
Keywords: Silicon carbide Nanotubes Transmission electron microscopy Microstructure Interfaces
1. Introduction Since the discovery of carbon nanotubes (CNTs) in 1991 [1], many researchers have synthesized new one-dimensional (1-D) nanostructured materials such as nanotubes, nanorods, and nanowires for potential applications because of their fascinating geometries and physical properties [2,3]. Some have reported the fabrication of many nanomaterials such as oxide, carbide, and nitride ceramics nanostructures, using CNTs as a template [4e8]. SiC is one of the most important wide-band-gap semiconducting materials for high-temperature and high-power applications [9]. Furthermore, SiC is also an important structural material at high temperatures [10]. Therefore, SiC offers exciting opportunities in both electronic devices and structural materials at high temperature [9,10]. We have previously reported on the formation of CeSiC coaxial nanotubes, in which CNTs were sheathed with SiC. Moreover, single-phase SiC nanotubes were successfully synthesized by heating the CeSiC coaxial nanotubes in air [11,12]. However, many
* Corresponding author. E-mail address: [email protected] (T. Taguchi). http://dx.doi.org/10.1016/j.carbon.2015.08.056 0008-6223/© 2015 Elsevier Ltd. All rights reserved.
grain boundaries were observed in the SiC nanotubes, because of the polycrystalline state of the SiC nanotubes fabricated in the previous studies [11,12]. SiC materials have four possible crystal states: single-crystalline, poly-crystalline, nano-crystalline, and amorphous. It is well-known that SiC materials of different crystal states have different physical properties, such as elastic constants and band gaps [13,14]. Amorphous SiC has attracted considerable interest as a material for constructing high-efficiency silicon solar cells [15]. Amorphous SiC is also applied as a passivation layer for crystalline silicon [16] and as a metal diffusion barrier for copper [17]. In general, the amorphization of a crystal material causes the uniformization of physical properties in the material, since neither grain boundaries nor crystal anisotropy exists. Therefore, the synthesis of amorphous SiC nanotubes without grain boundaries is required to attain uniformity. Bulk SiC materials have been reported to be amorphized by both ion and neutron irradiation [18]. Here we report, for the first time, the synthesis of amorphous SiC nanotubes by ion irradiation. Furthermore, it has been reported that the synthesis of 1-D heterostructures with different materials can produce materials with emphasized and novel electronic and optical properties, compared to the synthesis of 1-D homogeneous nanostructures [19]. Therefore, many 1-D heterostructures, such as axial
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and coaxial heterostructure nanowires and nanotubes, have been synthesized, including SieSiGe, GaPeGaAS, and InPeInAs axial nanowires and TiO2eSiO2, AlNeBN, ans CeSiCeSiO2 coaxial nanotubes [20e25]. However, the synthesis of heterostructured nanotubes with different crystal states of the same material has not yet been reported. Here, we report the successful synthesis of heterostructured SiC nanotubes, in which the polycrystalline and amorphous states coexist in one nanotube. 2. Experimental procedure Commercial CNTs (GSI Creos Corporation, Tokyo, Japan) were used as template materials. CeSiC coaxial nanotubes were synthesized by heating CNTs with Si powder (The Nilaco Corporation, Tokyo, Japan) at 1200 C for 100 h in a vacuum of around 5 104 Pa. The Si powder did not directly contact the CNTs. The amount of Si powder was much larger than that of the CNTs. Singlephase SiC nanotubes were produced by the heat treatment of CeSiC coaxial nanotubes at 600 C for 2 h in air. Both the single-phase SiC nanotubes and CeSiC coaxial nanotubes were ultrasonically dispersed in ethanol for 20 min. The dispersions of nanotubes in ethanol were each deposited on a nickel grid sample holder and dried at room temperature. The nickel grid holder, which held depositions of both single-phase SiC nanotubes and CeSiC coaxial nanotubes, was irradiated with 340 keV Siþ ions at a temperature below 100 C in the Takasaki Ion Accelerators for Advanced Radiation Application (TIARA) facility of the Japan Atomic Energy Agency. Silicon ions were chosen as the irradiation species because they were self-ions of the SiC nanotubes. The projected range of the 340 keV Siþ ions into the SiC target was calculated to be approximately 400 nm by Stopping and Range of Ions in Matter (SRIM) 2008 software [26], which indicated that the irradiated Siþ would completely penetrate the SiC nanotubes. The ion fluence was 2.4 1020 ions/m2, and the corresponding irradiation damage (displacement per atom, dpa) was calculated by SRIM 2008 [26] to be 6.8 dpa. In order to synthesize heterostructured SiC nanotubes with different crystal states, the nickel grid holder with both SiC and CeSiC nanotubes was irradiated with a nickel mask in front of the nickel grid holder. Microstructural observations were performed through transmission electron microscopy (TEM, Model 2100F, JEOL Ltd., Akishima, Japan) operating at 200 kV. Electron energy-loss spectroscopy (EELS, Enfinium spectrometer, Nippon Gatan, Nishi-Tokyo, Japan) was performed to evaluate the plasmon energies of each crystallinity phase present in the SiC nanotubes, such as polycrystalline and amorphous, by both TEM and scanning TEM (STEM) modes. By using STEM, the plasmon energies of an area of one SiC nanotube less than 1 nm in diameter can be evaluated. The volume swelling of SiC nanotube by amorphization was evaluated from the shift of plasmon energies in the polycrystalline and amorphous regions [27]. The experimental error in measurements of volume swelling in SiC as evaluated by this method is reported to be less than ±8% of the average measured value [27]. 3. Results and discussion Typical TEM images and selected area electron diffraction (SAED) patterns of SiC and CeSiC nanotubes irradiated with 340 keV Siþ ions are shown in Fig. 1. The SAED pattern of the SiC nanotube, depicted in the inset in Fig. 1(a), shows a hazy halo ring, indicating the amorphization of the polycrystalline SiC in the nanotube. This result demonstrates for the first time that amorphous SiC nanotubes can be synthesized by ion irradiation. According to high-resolution TEM (HR-TEM) observation, the amorphous structure is observed in the SiC nanotube irradiated with Siþ ions. In addition to the amorphization of single-phase SiC
nanotubes, polycrystalline CeSiC coaxial nanotubes are also transformed to amorphous CeSiC coaxial nanotubes by ion irradiation at a temperature less than 100 C, as shown in Fig. 1(c) and (d). The crystal states of the SiC layers in both the SiC and CeSiC nanotubes transform from polycrystalline to amorphous, while the tubular shapes of the SiC and CeSiC nanotubes are retained even after irradiation with Siþ. The orientation of the graphitic shells in the CeSiC nanotubes before ion irradiation was almost parallel to the length directions of the nanotubes [11,12]. However, HR-TEM observation indicates that the orientation of graphitic shells after ion irradiation changes such that they are completely parallel to the radial directions of the nanotubes. These results are consistent with that of the SAED pattern from the CeSiC nanotube, as depicted in the inset in Fig. 1(c). The graphite (002) spots in the SAED image are clearly observed even after ion irradiation up to 6.8 dpa, indicating the certain crystallinity of the carbon layer in the CeSiC nanotube. In ordinary microstructural changes in a multi-walled carbon nanotube (MWNT) by electron irradiation, the graphitic shells were deformed to induce interstitial loops [28]. The basal planes were broken into small segments, and the graphitic shells bent toward the interior hole [28]. Only after these deformations did the microstructures of MWNTs amorphize [28]. Krasheninnikov et al. reported a theoretical study of the microstructural changes of MWNT irradiated with 300 eV Ar ions by molecular dynamics simulations [29]. Defect-mediated covalent bonds between the graphitic shells were simulated to appear in the MWNT; these led to the amorphization of the graphitic shell [29]. Furthermore, it has been experimentally reported that amorphous carbon nanorods can be produced from MWNTs by 5 or 40 keV Arþ irradiation [30,31] with irradiation damage as calculated by SRIM 2008 of approximately 7 dpa [26], and 3.6 dpa [31], respectively. However, B. Q. Wei et al. reported that both the bending of the inner shells of the MWNT irradiated with 50 keV Gaþ ions to align with the radial direction of the nanotube completely and the formation of pillbox-like compartments occur at a dose of less than 1 1017 ions/m2 and an estimated irradiation damage calculated by SRIM 2008 [26] of 1.5 102 dpa [32]. Increasing the dose of Gaþ ions destroyed the nanocompartments and finally induced the formation of a homogenous amorphous carbon nanorod at a dose exceeding 1 1019 ions/m2 with approximately 1.5 dpa damage calculated by SRIM 2008 [26,32]. In this study, the CeSiC nanotubes are irradiated up to 6.8 dpa. However, the carbon layers in the CeSiC nanotubes are not transformed to a completely amorphous structure. The graphitic structure is still observed in HR-TEM images of the shells, and graphite (002) spots in SAED are clearly observed. The remaining graphitic portions of the shells have bent to align completely with the radial directions of the nanotubes, as shown in Fig. 1(d). The synthesis of MWNT with graphitic shells completely oriented parallel to the radial direction of the nanotube has not been reported yet. Such an MWNT with this new structure may be expected to find new applications. The SiC layer of the CeSiC nanotube can be used to investigate the deformation and bending of graphitic shells; the ejection of carbon atoms from both the inner and outer surfaces of the carbon layer is restrained by the SiC. Thus, the morphological stability of the carbon layer against irradiation with electrons or ions can be improved by coating the nanotube with a SiC layer. This technique may be an obstacle to the transformation of the carbon layer to a completely amorphous structure in a CeSiC nanotube. The volume of both the carbon [28] and the SiC layers [18] swells by ion irradiation. The carbon layer in the CeSiC nanotube is subjected to an applied stress because of the different rates of volume swelling between the carbon and SiC layers. This stress may be one of the driving forces of the bending of the graphitic shells.
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Fig. 1. Typical TEM images and SAED patterns of (a), (b) SiC nanotube, and (c), (d) CeSiC nanotube irradiated with 340 keV Siþ ions.
Fig. 2 shows typical TEM images and SAED patterns of a SiC nanotube irradiated with Siþ ions with a nickel mask. In Fig. 2(a), the SAED of the upper region of the SiC nanotube irradiated with a
mask clearly shows the existence of diffraction spots, while that of the lower region shows a hazy halo ring. These results indicate that a heterostructured SiC nanotube, with two different crystallinities
Fig. 2. Typical TEM images and SAED patterns of (a) and (b) SiC nanotube irradiated with Siþ ions with a mask. HR-TEM images of (c) crystalline, (d) interfacial, and (e) amorphous regions.
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and microstructures such as polycrystalline and amorphous in a single nanotube, was synthesized for the first time by ion irradiation with a mask. The inner diameter of the amorphous region is decreased compared to that of the polycrystalline region in all heterostructured SiC nanotubes, as shown in Fig. 2(b) and Fig. S1. According to the HR-TEM image of the interface between the amorphous and polycrystalline regions (Fig. 2(d)), the two regions are clearly separated by an interface. EELS spectra across the SiC nanotube heterostructure were obtained by TEM or STEM. The Si-L and CeK energy-loss near-edge structures (ELNES) of the amorphous, interfacial, and polycrystalline regions of the heterostructured SiC nanotube are shown in Fig. 3. Both fine structures of Si-L and CeK ELNES are lost by the amorphization, which is more conspicuously emphasized in the CeK ELNES. The loss of long-range order and chemical disordering causes these losses of fine structures. Fig. 4(a) shows the lowenergy EELS spectra taken from the amorphous, interfacial, and polycrystalline regions of the heterostructured SiC nanotube by TEM. The prominent peaks in all regions are located at 21e22 eV, corresponding to bulk plasmons. Fig. 4(b) shows the dependence of the plasmon energy of heterostructured SiC nanotubes on the distance from the interface between the amorphous and polycrystalline regions, as observed by STEM. The plasmon energies in the heterostructured SiC nanotube are decreased by the amorphization of the SiC. The plasmon energy Ep is given by.
sffiffiffiffiffiffiffiffiffiffiffi ne2 Ep ¼ Z ε0 m0
(1)
based on the Drude model for metals, where ħ is the Plank constant, n is the number density of valence electrons, e is the electron charge, m0 is the electron mass, and ε0 is the dielectric constant of vacuum [27]. In this study, the plasmon energy shift is considered to reflect the change in n, which is directly related to the material density [33]. Assuming the plasmon energy shift to be caused by the reduction of the atomic density of the amorphous region, the corresponding volume swelling DV by amorphization is evaluated by the following equation [33]:
DV ¼
Vamorphous Vcrystal ¼ Vcrystal
Ep; crystal Ep; amorphous
!2 1
(2)
The evaluated DV in this study is approximately 5.0%. The volume swelling of bulk SiC irradiated by neutrons or ions at room temperature has been reported to be between 10 and 15%, although DV also depends on the irradiation dose [34e36]. The irradiation damage of 6.8 dpa withstood by the SiC nanotubes in this study is sufficiently high to saturate the volume swelling [36]. The volume
swelling measured in this study is much smaller than those of bulk SiC from previous studies. The SiC nanotubes have very high surface areas and many grain boundaries as a result of their tubular shape, nanometer size, and small grain size. Therefore, point defects and cavities induced by the ion irradiation can be assumed to diffuse easily from the surface and grain boundaries in the SiC nanotubes in this study. It is more difficult to accumulate point defects and cavities in the SiC nanotubes than in bulk SiC. From these results, the volume swelling of SiC nanotubes in this study is much smaller than that of bulk SiC in the previous studies. Heera et al. reported the existence of a relaxed amorphous state of SiC, which has a higher density than that of as-amorphized SiC, by annealing SiC at 500 C [37]. The density change of the SiC nanotubes in this study (approximately 5%) is smaller than that of the SiC transformed from the amorphous to the relaxed amorphous state (approximately 7%) [37]. The downsizing of the SiC material to the nanoscale causes the formation of an amorphous SiC phase with higher density than that annealed at 500 C after ion irradiation. These results lead to the conclusion that relaxed amorphous SiC nanotubes with higher density, and thus lower amounts of vacancies and point defects, can be produced by the ion irradiation of polycrystalline SiC nanotubes. The measurements of the physical properties of such relaxed amorphous SiC nanotubes are expected in future work. Fig. 5 shows typical TEM images and SAED patterns of CeSiC nanotubes irradiated by 340 keV Siþ ions with a nickel mask. As in the SiC nanotube irradiated with a mask, the SAED of the right region of the CeSiC nanotube irradiated with a mask clearly shows the existence of diffraction spots, while that of the left region shows a hazy halo ring. These results demonstrate the synthesis of a polycrystalline/amorphous heterostructured CeSiC nanotube. According to TEM observations, the inner diameter of the amorphous region is smaller than that of the polycrystalline region in the heterostructured CeSiC nanotube. The graphitic structure of the shells of the irradiated region in the heterostructured CeSiC nanotube are still clearly observed; they are seen to be aligned with the radial direction of the nanotube, as shown in Fig. 5(e). Fig. 6 also shows typical TEM images and SAED patterns of a second type of CeSiC nanotube irradiated with a mask. As in Fig. 5, amorphous and polycrystalline regions of the SiC layer are seen to coexist in the CeSiC nanotube irradiated with a mask. In contrast to Fig. 5, the inner diameter in the amorphous region has decreased to zero. These results demonstrate the synthesis of a polycrystalline/ amorphous heterostructured CeSiC nanotube/nanowire. In both Figs. 5 and 6, the inner diameters do not change suddenly at the interface between polycrystalline and amorphous regions, but instead decrease continuously over more than 100 nm in the transition from polycrystalline to amorphous structures. It is possible that the nickel mask moves slightly during the ion
Fig. 3. (a) Si-L and (b) CeK energy-loss near-edge structures (ELNES) of the amorphous, interfacial, and polycrystalline regions of the heterostructured SiC nanotube.
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Fig. 4. (a) Low-energy EELS spectra taken from amorphous, interfacial, and polycrystalline regions of the heterostructured SiC nanotube by TEM. (b) Relationship between energies of plasmon peaks obtained by STEM and distance from interface.
Fig. 5. Typical TEM images and SAED patterns of (a) and (b) polycrystalline/amorphous CeSiC nanotube. HR-TEM images of (c) interface between polycrystalline and amorphous regions of SiC layer, (d) polycrystalline and (e) amorphous regions of carbon layer.
irradiation. However, the sharp interface between the polycrystalline and amorphous regions in the SiC layers, as shown in Figs. 5(c) and 6(c), can be clearly observed in both cases. Therefore, the possibility that the mask moves more than 100 nm is unlikely. The inner diameter does not change suddenly at the interface; it decreases gradually in order to relax the residual stress arising from the mismatch of the inner diameter size in the polycrystalline and amorphous regions. The schematic of shrinkage and the mechanism of the disappearance of the inner diameter in the CeSiC nanotube are shown in Fig. 7. The direction of the alignment of the graphitic shells in the CeSiC nanotube prior to ion irradiation is almost parallel to the length direction of the nanotube. Because the MWNTs used in this study are of the cup-stacked type, originally the graphitic shells are aligned with a slight bend relative to the radial direction, and the crystallinity of the MWNT is not good. According to TEM observations, the ion irradiation causes the graphitic shells to bend toward the radial direction of the nanotube. The bending angle of the graphitic shells in a CeSiC nanotube may increase with increasing irradiation damage. The inner diameter of the nanotube decreases gradually as the bending angle increases. In the case of a large initial inner diameter, even if the graphitic shells become completely aligned to the radial direction of the nanotube, the inner diameter is decreased but remains non-zero, as shown in Fig. 7(a).
Furthermore, the new MWNT with the graphitic shells aligned completely parallel to the radial direction of the nanotube is produced inside the amorphous SiC tubular layer. However, in the case of a small initial inner diameter, the inner diameter decreases and finally disappears because of the curvature of the graphitic shells, as shown in Fig. 7(b). In other words, the decrease and disappearance of the inner diameter depend on the initial size of the inner diameter and the thickness of the carbon layer. Table 1 provides a summary of the ion irradiation processing and the samples synthesized in this study. As shown in Table 1, amorphous SiC nanotubes, polycrystalline/amorphous heterostructured SiC nanotubes, and MWNTs with graphitic shells aligned parallel to the radial direction of the nanotube inside amorphous SiC nanotubes are synthesized for the first time in this study. These new SiC nanotubes can be expected to possess potential applications in nanoscale electronic and optical devices. The evaluation of the electronic properties, such as currentevoltage curves, of these new SiC nanotubes is required in further research. 4. Conclusions Amorphous SiC nanotubes were synthesized by 340 keV Siþ ion irradiation of polycrystalline SiC nanotubes for the first time. The crystal phases of the SiC layers in both the SiC and CeSiC nanotubes
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Fig. 6. Typical TEM images and SAED patterns of (a) and (b) polycrystalline/amorphous CeSiC nanotube/nanowire. HR-TEM images of (c) interface between polycrystalline and amorphous regions of SiC layer, (d) polycrystalline and (e) amorphous regions of carbon layer.
Fig. 7. Schematic of shrinkage and mechanism of disappearance of inner diameter in CeSiC nanotubes of (a) large and (b) small caliber.
Table 1 Summary of ion irradiation processing and samples. No.
Raw materials
State before irradiation
Irradiation process
State after irradiation
1 2 3
SiC nanotube SiC nanotube CeSiC nanotube
polycrystalline polycrystalline polycrystalline
without mask with mask without mask
4 5
CeSiC nanotube of large caliber CeSiC nanotube of small caliber
polycrystalline polycrystalline
with mask with mask
Amorphous SiC nanotube Polycrystalline/amorphous SiC nanotube MWNT with graphitic shells parallel to the radial direction inside amorphous SiC nanotube Polycrystalline/amorphous heterostructured CeSiC nanotube Polycrystalline/amorphous heterostructured CeSiC nanotube/nanowire
T. Taguchi et al. / Carbon 95 (2015) 279e285
transformed from polycrystalline to amorphous, while the tubular structures of the SiC and CeSiC nanotubes remained even after ion irradiation. The graphitic shells in the carbon layer of the CeSiC nanotubes gradually bent to become aligned with the radial directions of the nanotubes by ion irradiation. Since the graphite (002) spots in the SAED were clearly observed even after ion irradiation, the carbon layer in CeSiC nanotube was confirmed to remain crystalline. A new MWNT with graphitic shells aligned completely parallel to the radial direction of the nanotube was produced inside the amorphous SiC tubular layer in the case of CeSiC nanotubes with large inner diameters. A polycrystalline/ amorphous heterostructured SiC nanotube, in which polycrystalline and amorphous SiC coexisted in the same nanotube, was also synthesized by ion irradiation with a mask in front of a polycrystalline SiC nanotube. TEM observation revealed that the inner diameter of the SiC nanotubes shrank depending on the crystalline state of the SiC. According to EELS evaluation, the plasmon energies of the SiC nanotubes changed rapidly at the interfaces between polycrystalline and amorphous regions. The degree of volume swelling, evaluated from the differences of plasmon energies, was found to be approximately 5.0%. This is much smaller than those of bulk SiC materials measured in previous studies. The decreased measured volume swelling was considered to result from the easy diffusion of point defects and cavities induced by the ion irradiation from the surface and grain boundaries in the SiC nanotubes. This result revealed that relaxed amorphous SiC nanotubes with higher density could be produced by ion irradiation. Polycrystalline/ amorphous heterostructured CeSiC nanotube and polycrystalline/ amorphous heterostructured CeSiC nanotube/nanowire were also synthesized by ion irradiation with a mask of polycrystalline CeSiC nanotubes. The decrease and disappearance of the inner diameter of the CeSiC nanotube depended on the initial size of the inner diameter and the thickness of the carbon layer. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.carbon.2015.08.056. References [1] S. Iijima, Helical microtubules of graphitic carbon, Nature 354 (1991) 56e58. [2] Y. Xia, P. Yang, Y. Sun, Y. Wu, B. Mayers, B. Gates, et al., One-dimensional nanostructures: synthesis, characterization, and applications, Adv. Mater 15 (2003) 353e389. [3] H. Dai, Carbon nanotubes: opportunities and challenges, Surf. Sci. 500 (2002) 218e241. [4] H. Dai, E.W. Wong, Y.Z. Lu, S. Fan, C.M. Lieber, Synthesis and characterization of carbide nanorods, Nature 375 (1995) 769e772. [5] N.G. Chopra, R.J. Luyken, K. Cherrey, V.H. Crespi, M.L. Cohen, S.G. Louie, et al., Boron nitride nanotubes, Science 269 (1995) 966e967. [6] M. Zhang, Y. Bando, K. Wada, K. Kurashima, Synthesis of nanotubes and nanowires of silicon oxide, J. Mater. Sci. Lett. 18 (1999) 1911e1913. [7] J. Zhu, S. Fan, Nanostructure of GaN and SiC nanowires base on carbon nanotubes, J. Mater. Res. 14 (1999) 1175e1177. [8] T. Taguchi, H. Yamamoto, S. Shamoto, Synthesis and characterization of singlephase TiC nanotubes, TiC nanowires, and carbon nanotubes equipped with TiC nanoparticles, J. Phys. Chem. C 111 (2007) 18888e18891. [9] A. Fissel, B. Schroter, W. Richter, Low-temperature growth of SiC thin films on Si and 6H-SiC by solid-source molecular beam epitaxy, Appl. Phys. Lett. 66 (1995) 3182e3184. [10] Y. Katoh, L.L. Snead, C.H. Henager Jr., T. Nozawa, T. Hinoki, A. Ivekovic, et al., Current status and recent research achievements in SiC/SiC composites,
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Synthesis of heterostructured SiC and CeSiC nanotubes by ion irradiation-induced changes in crystallinity Tomitsugu Taguchi*, Shunya Yamamoto, Katsuaki Kodama, Hidehito Asaoka Quantum Beam Science Center, Japan Atomic Energy Agency, Tokai-mura, Ibaraki-ken 319-1195 Japan
a r t i c l e i n f o
a b s t r a c t
Article history: Received 1 May 2015 Received in revised form 17 August 2015 Accepted 18 August 2015 Available online 20 August 2015
Amorphous SiC nanotubes are successfully synthesized for the first time by the irradiation of polycrystalline SiC nanotubes with 340 keV Siþ ions. A polycrystalline/amorphous heterostructured SiC nanotube, in which polycrystalline SiC and amorphous SiC coexist in the same nanotube, is also synthesized by ion irradiation with a mask in front of the polycrystalline SiC nanotube. According to evaluation by electron energy loss spectroscopy, the plasmon energies of the SiC nanotube change rapidly at the interface between the polycrystalline and amorphous regions. The volume swelling induced by amorphization, as evaluated from the differences in plasmon energies, is approximately 5.0%. This result reveals that more relaxed amorphous SiC nanotubes with higher densities can be produced. The graphitic shells in the carbon layer of CeSiC nanotubes are found to gradually bend to align with the radial directions of the nanotubes by ion irradiation. Since graphite (002) spots in the selected-area electron diffraction pattern are clearly observed even after ion irradiation, the carbon layer in the CeSiC nanotube maintains crystallinity. Moreover, a new multi-walled carbon nanotube with graphitic shells completely parallel to the radial direction of the nanotube is also produced inside the amorphous SiC tubular layer, in the case of large-caliber CeSiC nanotubes. © 2015 Elsevier Ltd. All rights reserved.
Keywords: Silicon carbide Nanotubes Transmission electron microscopy Microstructure Interfaces
1. Introduction Since the discovery of carbon nanotubes (CNTs) in 1991 [1], many researchers have synthesized new one-dimensional (1-D) nanostructured materials such as nanotubes, nanorods, and nanowires for potential applications because of their fascinating geometries and physical properties [2,3]. Some have reported the fabrication of many nanomaterials such as oxide, carbide, and nitride ceramics nanostructures, using CNTs as a template [4e8]. SiC is one of the most important wide-band-gap semiconducting materials for high-temperature and high-power applications [9]. Furthermore, SiC is also an important structural material at high temperatures [10]. Therefore, SiC offers exciting opportunities in both electronic devices and structural materials at high temperature [9,10]. We have previously reported on the formation of CeSiC coaxial nanotubes, in which CNTs were sheathed with SiC. Moreover, single-phase SiC nanotubes were successfully synthesized by heating the CeSiC coaxial nanotubes in air [11,12]. However, many
* Corresponding author. E-mail address: [email protected] (T. Taguchi). http://dx.doi.org/10.1016/j.carbon.2015.08.056 0008-6223/© 2015 Elsevier Ltd. All rights reserved.
grain boundaries were observed in the SiC nanotubes, because of the polycrystalline state of the SiC nanotubes fabricated in the previous studies [11,12]. SiC materials have four possible crystal states: single-crystalline, poly-crystalline, nano-crystalline, and amorphous. It is well-known that SiC materials of different crystal states have different physical properties, such as elastic constants and band gaps [13,14]. Amorphous SiC has attracted considerable interest as a material for constructing high-efficiency silicon solar cells [15]. Amorphous SiC is also applied as a passivation layer for crystalline silicon [16] and as a metal diffusion barrier for copper [17]. In general, the amorphization of a crystal material causes the uniformization of physical properties in the material, since neither grain boundaries nor crystal anisotropy exists. Therefore, the synthesis of amorphous SiC nanotubes without grain boundaries is required to attain uniformity. Bulk SiC materials have been reported to be amorphized by both ion and neutron irradiation [18]. Here we report, for the first time, the synthesis of amorphous SiC nanotubes by ion irradiation. Furthermore, it has been reported that the synthesis of 1-D heterostructures with different materials can produce materials with emphasized and novel electronic and optical properties, compared to the synthesis of 1-D homogeneous nanostructures [19]. Therefore, many 1-D heterostructures, such as axial
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and coaxial heterostructure nanowires and nanotubes, have been synthesized, including SieSiGe, GaPeGaAS, and InPeInAs axial nanowires and TiO2eSiO2, AlNeBN, ans CeSiCeSiO2 coaxial nanotubes [20e25]. However, the synthesis of heterostructured nanotubes with different crystal states of the same material has not yet been reported. Here, we report the successful synthesis of heterostructured SiC nanotubes, in which the polycrystalline and amorphous states coexist in one nanotube. 2. Experimental procedure Commercial CNTs (GSI Creos Corporation, Tokyo, Japan) were used as template materials. CeSiC coaxial nanotubes were synthesized by heating CNTs with Si powder (The Nilaco Corporation, Tokyo, Japan) at 1200 C for 100 h in a vacuum of around 5 104 Pa. The Si powder did not directly contact the CNTs. The amount of Si powder was much larger than that of the CNTs. Singlephase SiC nanotubes were produced by the heat treatment of CeSiC coaxial nanotubes at 600 C for 2 h in air. Both the single-phase SiC nanotubes and CeSiC coaxial nanotubes were ultrasonically dispersed in ethanol for 20 min. The dispersions of nanotubes in ethanol were each deposited on a nickel grid sample holder and dried at room temperature. The nickel grid holder, which held depositions of both single-phase SiC nanotubes and CeSiC coaxial nanotubes, was irradiated with 340 keV Siþ ions at a temperature below 100 C in the Takasaki Ion Accelerators for Advanced Radiation Application (TIARA) facility of the Japan Atomic Energy Agency. Silicon ions were chosen as the irradiation species because they were self-ions of the SiC nanotubes. The projected range of the 340 keV Siþ ions into the SiC target was calculated to be approximately 400 nm by Stopping and Range of Ions in Matter (SRIM) 2008 software [26], which indicated that the irradiated Siþ would completely penetrate the SiC nanotubes. The ion fluence was 2.4 1020 ions/m2, and the corresponding irradiation damage (displacement per atom, dpa) was calculated by SRIM 2008 [26] to be 6.8 dpa. In order to synthesize heterostructured SiC nanotubes with different crystal states, the nickel grid holder with both SiC and CeSiC nanotubes was irradiated with a nickel mask in front of the nickel grid holder. Microstructural observations were performed through transmission electron microscopy (TEM, Model 2100F, JEOL Ltd., Akishima, Japan) operating at 200 kV. Electron energy-loss spectroscopy (EELS, Enfinium spectrometer, Nippon Gatan, Nishi-Tokyo, Japan) was performed to evaluate the plasmon energies of each crystallinity phase present in the SiC nanotubes, such as polycrystalline and amorphous, by both TEM and scanning TEM (STEM) modes. By using STEM, the plasmon energies of an area of one SiC nanotube less than 1 nm in diameter can be evaluated. The volume swelling of SiC nanotube by amorphization was evaluated from the shift of plasmon energies in the polycrystalline and amorphous regions [27]. The experimental error in measurements of volume swelling in SiC as evaluated by this method is reported to be less than ±8% of the average measured value [27]. 3. Results and discussion Typical TEM images and selected area electron diffraction (SAED) patterns of SiC and CeSiC nanotubes irradiated with 340 keV Siþ ions are shown in Fig. 1. The SAED pattern of the SiC nanotube, depicted in the inset in Fig. 1(a), shows a hazy halo ring, indicating the amorphization of the polycrystalline SiC in the nanotube. This result demonstrates for the first time that amorphous SiC nanotubes can be synthesized by ion irradiation. According to high-resolution TEM (HR-TEM) observation, the amorphous structure is observed in the SiC nanotube irradiated with Siþ ions. In addition to the amorphization of single-phase SiC
nanotubes, polycrystalline CeSiC coaxial nanotubes are also transformed to amorphous CeSiC coaxial nanotubes by ion irradiation at a temperature less than 100 C, as shown in Fig. 1(c) and (d). The crystal states of the SiC layers in both the SiC and CeSiC nanotubes transform from polycrystalline to amorphous, while the tubular shapes of the SiC and CeSiC nanotubes are retained even after irradiation with Siþ. The orientation of the graphitic shells in the CeSiC nanotubes before ion irradiation was almost parallel to the length directions of the nanotubes [11,12]. However, HR-TEM observation indicates that the orientation of graphitic shells after ion irradiation changes such that they are completely parallel to the radial directions of the nanotubes. These results are consistent with that of the SAED pattern from the CeSiC nanotube, as depicted in the inset in Fig. 1(c). The graphite (002) spots in the SAED image are clearly observed even after ion irradiation up to 6.8 dpa, indicating the certain crystallinity of the carbon layer in the CeSiC nanotube. In ordinary microstructural changes in a multi-walled carbon nanotube (MWNT) by electron irradiation, the graphitic shells were deformed to induce interstitial loops [28]. The basal planes were broken into small segments, and the graphitic shells bent toward the interior hole [28]. Only after these deformations did the microstructures of MWNTs amorphize [28]. Krasheninnikov et al. reported a theoretical study of the microstructural changes of MWNT irradiated with 300 eV Ar ions by molecular dynamics simulations [29]. Defect-mediated covalent bonds between the graphitic shells were simulated to appear in the MWNT; these led to the amorphization of the graphitic shell [29]. Furthermore, it has been experimentally reported that amorphous carbon nanorods can be produced from MWNTs by 5 or 40 keV Arþ irradiation [30,31] with irradiation damage as calculated by SRIM 2008 of approximately 7 dpa [26], and 3.6 dpa [31], respectively. However, B. Q. Wei et al. reported that both the bending of the inner shells of the MWNT irradiated with 50 keV Gaþ ions to align with the radial direction of the nanotube completely and the formation of pillbox-like compartments occur at a dose of less than 1 1017 ions/m2 and an estimated irradiation damage calculated by SRIM 2008 [26] of 1.5 102 dpa [32]. Increasing the dose of Gaþ ions destroyed the nanocompartments and finally induced the formation of a homogenous amorphous carbon nanorod at a dose exceeding 1 1019 ions/m2 with approximately 1.5 dpa damage calculated by SRIM 2008 [26,32]. In this study, the CeSiC nanotubes are irradiated up to 6.8 dpa. However, the carbon layers in the CeSiC nanotubes are not transformed to a completely amorphous structure. The graphitic structure is still observed in HR-TEM images of the shells, and graphite (002) spots in SAED are clearly observed. The remaining graphitic portions of the shells have bent to align completely with the radial directions of the nanotubes, as shown in Fig. 1(d). The synthesis of MWNT with graphitic shells completely oriented parallel to the radial direction of the nanotube has not been reported yet. Such an MWNT with this new structure may be expected to find new applications. The SiC layer of the CeSiC nanotube can be used to investigate the deformation and bending of graphitic shells; the ejection of carbon atoms from both the inner and outer surfaces of the carbon layer is restrained by the SiC. Thus, the morphological stability of the carbon layer against irradiation with electrons or ions can be improved by coating the nanotube with a SiC layer. This technique may be an obstacle to the transformation of the carbon layer to a completely amorphous structure in a CeSiC nanotube. The volume of both the carbon [28] and the SiC layers [18] swells by ion irradiation. The carbon layer in the CeSiC nanotube is subjected to an applied stress because of the different rates of volume swelling between the carbon and SiC layers. This stress may be one of the driving forces of the bending of the graphitic shells.
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Fig. 1. Typical TEM images and SAED patterns of (a), (b) SiC nanotube, and (c), (d) CeSiC nanotube irradiated with 340 keV Siþ ions.
Fig. 2 shows typical TEM images and SAED patterns of a SiC nanotube irradiated with Siþ ions with a nickel mask. In Fig. 2(a), the SAED of the upper region of the SiC nanotube irradiated with a
mask clearly shows the existence of diffraction spots, while that of the lower region shows a hazy halo ring. These results indicate that a heterostructured SiC nanotube, with two different crystallinities
Fig. 2. Typical TEM images and SAED patterns of (a) and (b) SiC nanotube irradiated with Siþ ions with a mask. HR-TEM images of (c) crystalline, (d) interfacial, and (e) amorphous regions.
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and microstructures such as polycrystalline and amorphous in a single nanotube, was synthesized for the first time by ion irradiation with a mask. The inner diameter of the amorphous region is decreased compared to that of the polycrystalline region in all heterostructured SiC nanotubes, as shown in Fig. 2(b) and Fig. S1. According to the HR-TEM image of the interface between the amorphous and polycrystalline regions (Fig. 2(d)), the two regions are clearly separated by an interface. EELS spectra across the SiC nanotube heterostructure were obtained by TEM or STEM. The Si-L and CeK energy-loss near-edge structures (ELNES) of the amorphous, interfacial, and polycrystalline regions of the heterostructured SiC nanotube are shown in Fig. 3. Both fine structures of Si-L and CeK ELNES are lost by the amorphization, which is more conspicuously emphasized in the CeK ELNES. The loss of long-range order and chemical disordering causes these losses of fine structures. Fig. 4(a) shows the lowenergy EELS spectra taken from the amorphous, interfacial, and polycrystalline regions of the heterostructured SiC nanotube by TEM. The prominent peaks in all regions are located at 21e22 eV, corresponding to bulk plasmons. Fig. 4(b) shows the dependence of the plasmon energy of heterostructured SiC nanotubes on the distance from the interface between the amorphous and polycrystalline regions, as observed by STEM. The plasmon energies in the heterostructured SiC nanotube are decreased by the amorphization of the SiC. The plasmon energy Ep is given by.
sffiffiffiffiffiffiffiffiffiffiffi ne2 Ep ¼ Z ε0 m0
(1)
based on the Drude model for metals, where ħ is the Plank constant, n is the number density of valence electrons, e is the electron charge, m0 is the electron mass, and ε0 is the dielectric constant of vacuum [27]. In this study, the plasmon energy shift is considered to reflect the change in n, which is directly related to the material density [33]. Assuming the plasmon energy shift to be caused by the reduction of the atomic density of the amorphous region, the corresponding volume swelling DV by amorphization is evaluated by the following equation [33]:
DV ¼
Vamorphous Vcrystal ¼ Vcrystal
Ep; crystal Ep; amorphous
!2 1
(2)
The evaluated DV in this study is approximately 5.0%. The volume swelling of bulk SiC irradiated by neutrons or ions at room temperature has been reported to be between 10 and 15%, although DV also depends on the irradiation dose [34e36]. The irradiation damage of 6.8 dpa withstood by the SiC nanotubes in this study is sufficiently high to saturate the volume swelling [36]. The volume
swelling measured in this study is much smaller than those of bulk SiC from previous studies. The SiC nanotubes have very high surface areas and many grain boundaries as a result of their tubular shape, nanometer size, and small grain size. Therefore, point defects and cavities induced by the ion irradiation can be assumed to diffuse easily from the surface and grain boundaries in the SiC nanotubes in this study. It is more difficult to accumulate point defects and cavities in the SiC nanotubes than in bulk SiC. From these results, the volume swelling of SiC nanotubes in this study is much smaller than that of bulk SiC in the previous studies. Heera et al. reported the existence of a relaxed amorphous state of SiC, which has a higher density than that of as-amorphized SiC, by annealing SiC at 500 C [37]. The density change of the SiC nanotubes in this study (approximately 5%) is smaller than that of the SiC transformed from the amorphous to the relaxed amorphous state (approximately 7%) [37]. The downsizing of the SiC material to the nanoscale causes the formation of an amorphous SiC phase with higher density than that annealed at 500 C after ion irradiation. These results lead to the conclusion that relaxed amorphous SiC nanotubes with higher density, and thus lower amounts of vacancies and point defects, can be produced by the ion irradiation of polycrystalline SiC nanotubes. The measurements of the physical properties of such relaxed amorphous SiC nanotubes are expected in future work. Fig. 5 shows typical TEM images and SAED patterns of CeSiC nanotubes irradiated by 340 keV Siþ ions with a nickel mask. As in the SiC nanotube irradiated with a mask, the SAED of the right region of the CeSiC nanotube irradiated with a mask clearly shows the existence of diffraction spots, while that of the left region shows a hazy halo ring. These results demonstrate the synthesis of a polycrystalline/amorphous heterostructured CeSiC nanotube. According to TEM observations, the inner diameter of the amorphous region is smaller than that of the polycrystalline region in the heterostructured CeSiC nanotube. The graphitic structure of the shells of the irradiated region in the heterostructured CeSiC nanotube are still clearly observed; they are seen to be aligned with the radial direction of the nanotube, as shown in Fig. 5(e). Fig. 6 also shows typical TEM images and SAED patterns of a second type of CeSiC nanotube irradiated with a mask. As in Fig. 5, amorphous and polycrystalline regions of the SiC layer are seen to coexist in the CeSiC nanotube irradiated with a mask. In contrast to Fig. 5, the inner diameter in the amorphous region has decreased to zero. These results demonstrate the synthesis of a polycrystalline/ amorphous heterostructured CeSiC nanotube/nanowire. In both Figs. 5 and 6, the inner diameters do not change suddenly at the interface between polycrystalline and amorphous regions, but instead decrease continuously over more than 100 nm in the transition from polycrystalline to amorphous structures. It is possible that the nickel mask moves slightly during the ion
Fig. 3. (a) Si-L and (b) CeK energy-loss near-edge structures (ELNES) of the amorphous, interfacial, and polycrystalline regions of the heterostructured SiC nanotube.
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Fig. 4. (a) Low-energy EELS spectra taken from amorphous, interfacial, and polycrystalline regions of the heterostructured SiC nanotube by TEM. (b) Relationship between energies of plasmon peaks obtained by STEM and distance from interface.
Fig. 5. Typical TEM images and SAED patterns of (a) and (b) polycrystalline/amorphous CeSiC nanotube. HR-TEM images of (c) interface between polycrystalline and amorphous regions of SiC layer, (d) polycrystalline and (e) amorphous regions of carbon layer.
irradiation. However, the sharp interface between the polycrystalline and amorphous regions in the SiC layers, as shown in Figs. 5(c) and 6(c), can be clearly observed in both cases. Therefore, the possibility that the mask moves more than 100 nm is unlikely. The inner diameter does not change suddenly at the interface; it decreases gradually in order to relax the residual stress arising from the mismatch of the inner diameter size in the polycrystalline and amorphous regions. The schematic of shrinkage and the mechanism of the disappearance of the inner diameter in the CeSiC nanotube are shown in Fig. 7. The direction of the alignment of the graphitic shells in the CeSiC nanotube prior to ion irradiation is almost parallel to the length direction of the nanotube. Because the MWNTs used in this study are of the cup-stacked type, originally the graphitic shells are aligned with a slight bend relative to the radial direction, and the crystallinity of the MWNT is not good. According to TEM observations, the ion irradiation causes the graphitic shells to bend toward the radial direction of the nanotube. The bending angle of the graphitic shells in a CeSiC nanotube may increase with increasing irradiation damage. The inner diameter of the nanotube decreases gradually as the bending angle increases. In the case of a large initial inner diameter, even if the graphitic shells become completely aligned to the radial direction of the nanotube, the inner diameter is decreased but remains non-zero, as shown in Fig. 7(a).
Furthermore, the new MWNT with the graphitic shells aligned completely parallel to the radial direction of the nanotube is produced inside the amorphous SiC tubular layer. However, in the case of a small initial inner diameter, the inner diameter decreases and finally disappears because of the curvature of the graphitic shells, as shown in Fig. 7(b). In other words, the decrease and disappearance of the inner diameter depend on the initial size of the inner diameter and the thickness of the carbon layer. Table 1 provides a summary of the ion irradiation processing and the samples synthesized in this study. As shown in Table 1, amorphous SiC nanotubes, polycrystalline/amorphous heterostructured SiC nanotubes, and MWNTs with graphitic shells aligned parallel to the radial direction of the nanotube inside amorphous SiC nanotubes are synthesized for the first time in this study. These new SiC nanotubes can be expected to possess potential applications in nanoscale electronic and optical devices. The evaluation of the electronic properties, such as currentevoltage curves, of these new SiC nanotubes is required in further research. 4. Conclusions Amorphous SiC nanotubes were synthesized by 340 keV Siþ ion irradiation of polycrystalline SiC nanotubes for the first time. The crystal phases of the SiC layers in both the SiC and CeSiC nanotubes
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Fig. 6. Typical TEM images and SAED patterns of (a) and (b) polycrystalline/amorphous CeSiC nanotube/nanowire. HR-TEM images of (c) interface between polycrystalline and amorphous regions of SiC layer, (d) polycrystalline and (e) amorphous regions of carbon layer.
Fig. 7. Schematic of shrinkage and mechanism of disappearance of inner diameter in CeSiC nanotubes of (a) large and (b) small caliber.
Table 1 Summary of ion irradiation processing and samples. No.
Raw materials
State before irradiation
Irradiation process
State after irradiation
1 2 3
SiC nanotube SiC nanotube CeSiC nanotube
polycrystalline polycrystalline polycrystalline
without mask with mask without mask
4 5
CeSiC nanotube of large caliber CeSiC nanotube of small caliber
polycrystalline polycrystalline
with mask with mask
Amorphous SiC nanotube Polycrystalline/amorphous SiC nanotube MWNT with graphitic shells parallel to the radial direction inside amorphous SiC nanotube Polycrystalline/amorphous heterostructured CeSiC nanotube Polycrystalline/amorphous heterostructured CeSiC nanotube/nanowire
T. Taguchi et al. / Carbon 95 (2015) 279e285
transformed from polycrystalline to amorphous, while the tubular structures of the SiC and CeSiC nanotubes remained even after ion irradiation. The graphitic shells in the carbon layer of the CeSiC nanotubes gradually bent to become aligned with the radial directions of the nanotubes by ion irradiation. Since the graphite (002) spots in the SAED were clearly observed even after ion irradiation, the carbon layer in CeSiC nanotube was confirmed to remain crystalline. A new MWNT with graphitic shells aligned completely parallel to the radial direction of the nanotube was produced inside the amorphous SiC tubular layer in the case of CeSiC nanotubes with large inner diameters. A polycrystalline/ amorphous heterostructured SiC nanotube, in which polycrystalline and amorphous SiC coexisted in the same nanotube, was also synthesized by ion irradiation with a mask in front of a polycrystalline SiC nanotube. TEM observation revealed that the inner diameter of the SiC nanotubes shrank depending on the crystalline state of the SiC. According to EELS evaluation, the plasmon energies of the SiC nanotubes changed rapidly at the interfaces between polycrystalline and amorphous regions. The degree of volume swelling, evaluated from the differences of plasmon energies, was found to be approximately 5.0%. This is much smaller than those of bulk SiC materials measured in previous studies. The decreased measured volume swelling was considered to result from the easy diffusion of point defects and cavities induced by the ion irradiation from the surface and grain boundaries in the SiC nanotubes. This result revealed that relaxed amorphous SiC nanotubes with higher density could be produced by ion irradiation. Polycrystalline/ amorphous heterostructured CeSiC nanotube and polycrystalline/ amorphous heterostructured CeSiC nanotube/nanowire were also synthesized by ion irradiation with a mask of polycrystalline CeSiC nanotubes. The decrease and disappearance of the inner diameter of the CeSiC nanotube depended on the initial size of the inner diameter and the thickness of the carbon layer. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.carbon.2015.08.056. References [1] S. Iijima, Helical microtubules of graphitic carbon, Nature 354 (1991) 56e58. [2] Y. Xia, P. Yang, Y. Sun, Y. Wu, B. Mayers, B. Gates, et al., One-dimensional nanostructures: synthesis, characterization, and applications, Adv. Mater 15 (2003) 353e389. [3] H. Dai, Carbon nanotubes: opportunities and challenges, Surf. Sci. 500 (2002) 218e241. [4] H. Dai, E.W. Wong, Y.Z. Lu, S. Fan, C.M. Lieber, Synthesis and characterization of carbide nanorods, Nature 375 (1995) 769e772. [5] N.G. Chopra, R.J. Luyken, K. Cherrey, V.H. Crespi, M.L. Cohen, S.G. Louie, et al., Boron nitride nanotubes, Science 269 (1995) 966e967. [6] M. Zhang, Y. Bando, K. Wada, K. Kurashima, Synthesis of nanotubes and nanowires of silicon oxide, J. Mater. Sci. Lett. 18 (1999) 1911e1913. [7] J. Zhu, S. Fan, Nanostructure of GaN and SiC nanowires base on carbon nanotubes, J. Mater. Res. 14 (1999) 1175e1177. [8] T. Taguchi, H. Yamamoto, S. Shamoto, Synthesis and characterization of singlephase TiC nanotubes, TiC nanowires, and carbon nanotubes equipped with TiC nanoparticles, J. Phys. Chem. C 111 (2007) 18888e18891. [9] A. Fissel, B. Schroter, W. Richter, Low-temperature growth of SiC thin films on Si and 6H-SiC by solid-source molecular beam epitaxy, Appl. Phys. Lett. 66 (1995) 3182e3184. [10] Y. Katoh, L.L. Snead, C.H. Henager Jr., T. Nozawa, T. Hinoki, A. Ivekovic, et al., Current status and recent research achievements in SiC/SiC composites,
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