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Nanostructured superhard carbon phase obtained under high pressure with shear deformation from single-wall nanotubes HiPco
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Physica B 382 (2006) 58–64 www.elsevier.com/locate/physb
Nanostructured superhard carbon phase obtained under high pressure with shear deformation from single-wall nanotubes HiPco V.D. Blanka, V.N. Denisova,b, A.N. Kirichenkoa, N.A. Lvovaa, S.Y. Martyushova, B.N. Mavrinb,, D.M. Popovaa, M.Yu. Popova,b, E.V. Tat’yanina, A.A. Zakhidovc a
Technological Institute for Superhard and Novel Carbon Materials Troitsk, Moscow reg., Russian Federation b Institute of Spectroscopy of the Russian Academy of Sciences, Troitsk, Moscow reg., Russian Federation c NanoTech Institute and Department of Physics, University of Texas at Dallas, Richardson, USA Received 22 November 2005; received in revised form 30 January 2006; accepted 30 January 2006
Abstract Raman spectra of single-wall nanotubes under high pressure combined with shear deformation are investigated in situ in a diamond cell. Shear deformation applied under 35 GPa led to pressure multiplication up to 60 GPa, and increased the intensity of Raman bands more than ten times without essential change of G-mode position while causing its essential broadening. The G-mode remained broad after pressure unloading and shifted to 1534 cm1. The hardness of the superhard material was 5876 GPa, comparable to the hardness of carbo-boro-nitride. A broad band appeared in the photoluminescence spectrum with a maximum at about 2 eV, which allowed to assume a high content of sp3-bonds in the sample. The large dispersion of the G-mode almost vanished after pressure unloading. However, a noticeable dispersion of the D-mode was found, which is a sign of certain ordering in the superhard phase. TEM study of the superhard phase detected clusters with graphene sheets with a size about 1 nm. r 2006 Elsevier B.V. All rights reserved. PACS: 61.46.+w; 62.50.+p; 78.30.Ly Keywords: Nanotube; High pressure; Raman; Hardness
1. Introduction The carbon superhard phase synthesis is based on the transformation of the carbon sp2-hybridized states (which are the stable states of a two-dimensional carbon atom phase, such as graphene sheets ) into sp3-states. Atoms in a sp3-state are bound to s-bonds, which form a threedimensional network with high elastic characteristics. As usual, sp2–sp3 transformations take place under high pressures and high temperatures (for instance, phase transformations from graphite into diamond [1], or fullerite into superhard phase [2–4]). In Ref. [5] it was reported that a graphite single crystal, when compressed along the c-axis, irreversibly transforms
Corresponding author. Tel.: +7 095 334 0238; fax: +7 095 334 0238.
E-mail address: [email protected] (B.N. Mavrin). 0921-4526/$ - see front matter r 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.physb.2006.01.519
to hexagonal diamond at a pressure around 12 GPa and upon heating above 1000 1C. This phase transition was studied at room temperature using in situ X-ray measurements [6]. Under the experimental conditions of Ref. [6], this transition was reversible, i.e. upon pressure unloading hexagonal diamond transforms back to graphite. When graphite was compressed along the c-axis, the phase transition was detected at 14 GPa and room temperature [6]. The formation of sp3-bonds between graphene sheets at room temperature is sensitive to the stress tensor in the specimen. For example, at hydrostatic conditions graphite does not transform to diamond under pressures of at least up to 80 GPa [7], while the stress tensor variation (by applying a shear deformation under load) and a large plastic deformation lead to a direct irreversible transformation of graphite to hexagonal diamond at a pressure of 17 GPa at room temperature [1].
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As for carbon nanotubes under high pressure, Raman spectra [8–16], X-ray [17–19] and neutron [20] diffraction under high pressures were investigated. In the Raman spectra it was revealed that the intensity of the radial breathing mode vanishes at about 2 GPa, which was explained by either hexagonal [9] or oval [11] distortion of a cylindrical section or its collapse [16]. The theoretical investigations of the nanotube structure under increasing pressure came to the conclusion that a certain critical pressure exists, which is inversely proportional to the nanotube diameter, over which the cross-section of a nanotube becomes elliptical with subsequent collapse [21,22]. Recently it was shown that the pressure effects are reversible (at least, partially) on release of pressures as high as 25.9 GPa [13], 55 GPa [14,15] and 40 GPa [23], which were essentially higher than the critical ones. Possible polymerization of nanotubes was predicted in some papers. However, for the polymerization, pressures [24] higher than the critical collapse pressure were required [21,22]. In the papers [3,14,15] polymerization of nanotubes under the pressure of 24 GPa under conditions of plastic deformation was assumed. The conclusion about the polymerization was made on the basis of the conservation of the spectral features of nanotubes in the Raman spectra of nanotube samples after application of a high pressure of 30 GPa, while the hardness of the nanotube sample significantly increased, becoming comparable with hardness of c-BN. In the present study we have investigated Raman spectra of single-wall nanotubes, synthesized by HiPco method, in a diamond high-pressure cell with the simultaneous application of shear deformation. The photoluminescence spectra and the hardness of the samples after the loading were measured as well. We have found that the application of shear deformation under a pressure higher than 35 GPa allowed the formation of a new nano-structured carbon superhard phase, which remains stable after pressure unloading.
2. Experimental procedure The initial single-wall nanotubes (SWNT) made by HiPco (Carbon Technologies, Inc.) were cleaned from iron atoms by the method described in Ref. [19]. Although the HiPcoSWNTs contain a wider diameter distribution (0.7–1.4 nm), the nanotubes with the narrow distribution range of 0.88–0.91 nm diameters had contributed to the most intense reflections in the X-ray diffraction patterns [19]. All high-pressure experiments were performed using a shear diamond anvil cell (SDAC). Specimens were loaded into the gasket without a pressure transfer medium. In the SDAC a controlled shear deformation is applied to a specimen under pressure by rotation of one of the anvils around the load axis. The application of shear deformation decreases the hysteresis of the phase transformation and
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makes it possible to obtain a homogeneous phase (see, e.g. Ref. [3] and references therein). A spectrometer with three-fold monochromatization, a CCD-camera cooled by nitrogen and a confocal Ramanmicroscope with spatial resolution of 1–2 mm were used for the detection of Raman spectra, which were excited by Ar–Kr laser lines. The pressure inside the diamond cell was determined by a piezospectroscopy method, thus circumventing the uncertainty of the usage of rubyscale under the conditions, when the ruby crystal is pressed in the hard matrix of the sample [3,15,25]. The first-order Raman spectrum of diamond exhibits a single peak that corresponds to the triply degenerate optical phonons in the absence of stresses; the uniaxial stress causes splittings and shifts, which are linear with the stress. In the case of compression along the [1 0 0] diamond crystallographic direction, the single diamond peak shifts to the high-frequency region and splits into singlet and doublet modes. The mode splitting is proportional to the deviator part (pure shear) of the stress tensor, while the shift of the centroid of the split bands is proportional to the spherical part of the stress tensor. This effect can be used for measurements of stresses in diamond anvils (normally the diamond anvils are loaded along the [1 0 0] axis) and pressure in a sample compressed between the anvils [15,25]. The luminescence spectra were excited by a 488 nm line of a laser and were recorded by a TRIAX spectrograph. Hardness measurements were performed on the submicron length scale using the NanoScan (NS) measurement system based on the principles of the scanning force microscopy [26–28]. Earlier this measurement method for hardness was tested on other hard and superhard materials, including the hardness measurements of c-BN, diamond and ultrahard fullerite [27,28]. Nanoscan also allows to measure the Young modulus of the samples [29,30]. 3. Results of the experiment Investigating the Raman spectra, we concentrated the main attention on the spectral region above 1000 cm1, where the tangential modes of nanotubes’ sp2-bonds and their overtones were expected. In addition to the Raman band of a diamond anvil two overlapping bands at 1592 and 1520 cm1 (seen in Fig. 1), related to the tangential modes of sp2-bonds of the nanotubes, were seen in the spectrum at the initial stages of the high-pressure application cycle ðP ¼ 1 GPaÞ. Above 10 GPa the intensity of the nanotubes’ spectrum clearly was reduced and the bands transformed into one wide band, with a maximum shifted with the increase of the pressure in correspondence with the preceding measurements [14,15]. The shear deformation was applied by turning one of the diamond anvils relative to another anvil at the pressure of 35 GPa. Due to shear deformation application multiplication of pressure takes place [3], and the effective pressure in the sample increased from 35 to
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Fig. 2. Comparison of Raman spectra of HiPco single-wall nanotubes before shear deformation application (at P ¼ 35 GPa) and directly after application of a shear deformation. Fig. 1. Raman spectra of single-wall nanotubes (made by HiPco method) for increasing pressure before shear deformation application.
60 GPa. After that, the intensity of the Raman spectrum increased more than 10 times, the width of the maximum increased, but its position almost did not change. The background of luminescence noticeably increased, over which the Raman spectrum can be seen (as shown in Fig. 2). During the pressure reducing cycle, the Raman spectrum intensity and its following luminescence background decreased, and the maximum of the Raman G-band shifted to low frequencies (Fig. 3). After the complete pressure unloading one broad band with a maximum at 1534 cm1 and a very faint shoulder near 1350 cm1 remained. The pressure dependence of the high-frequency mode of Raman scattering (tangential mode at 1592 cm1 in the initial nanotubes; transformed to 1534 cm1 mode after unloading) probably reflects some structural changes in the sample. The results of these two experiments are presented in Fig. 4. The dependencies of Raman frequencies on pressure are almost identical for these experiments. Sharp turns on the pressure dependence curve during pressure loading (at Fig. 4) at 15 and 24 GPa were noticed earlier and have been attributed to structural changes [14,15]. At the phase transformation, a step-like anomaly
of the radial pressure distribution in the sample due to jumps in volume and elastic modulus appears in the case of non-hydrostatic compression. In Ref. [14], these anomalies were detected at pressures of 15, 19 and 24 GPa. In Refs. [14,15], the sharp turn at 24 GPa was attributed to a possible polymerization of nanotubes. We investigated the dependence of the Raman spectrum on the laser excitation wavelength li. To observe this dependence for the Raman line at 1350 cm1 as well, we carried out an experiment with shear deformation application at lower pressure (the maximum pressure after the shear deformation raised only to P ¼ 50 GPa). In this case the 1350 cm1 line was more intensive after unloading, but the character of the spectrum right after application of shear deformation remained the same as in the first experiment. First we investigated the dependence of the spectrum on excitation wavelength li right after the shear deformation ðP ¼ 50 GPaÞ. At excitation by the li ¼ 488 nm laser line, the maximum of the Raman band was at the frequency of 1697 cm1 (Fig. 5), which red-shifted with the increase of li (to 1665 cm1 if excited by 514.5 nm and to 1650 cm1 if excited by 647 nm). The frequency shift is quite sizable, larger than 100 cm1/eV. After pressure unloading two bands were seen in the Raman spectrum, which had a less strong dependence of the maximum position on excitation li (as shown in Fig. 5) shifts to 1557
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Fig. 3. Raman spectra of HiPco single-wall nanotubes for decreasing pressure after shear deformation application.
Fig. 4. Dependence of the frequency at 1592 cm1 on pressure for two experiments (33 GPa (m) and 35 GPa (K) before shear deformation application) during loading (filled signs) and unloading (open signs).
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Fig. 5. Dependence of Raman spectra of HiPco single-wall nanotubes on the excitation wavelength at the pressure of 50 GPa (after shear deformation application) and after unloading.
and 1368 cm1 at 488 nm excitation, and shifts to 1548 and 1335 cm1 at 647 nm excitation. The frequency shift of the 1350 cm1 line corresponds to 15 cm1/eV. Relative intensities of the discussed pair of Raman bands are found to be dependent on the excitation li. While the bands excited by the li ¼ 647 nm laser line were of about the same intensity, the band at 1368 cm1 was essentially weaker than the second band at 1557 cm1, if excited by li ¼ 488 nm (see Fig. 5). Photoluminescence spectra were obtained for the samples after unloading in the spectral region of 480–1000 nm. A very broad photoluminescence band with a PL maximum about 650 nm (1.9 eV) was seen in the photoluminescence spectrum of the sample from the second experiment (Fig. 6). On the wing of the PL spectrum several narrow bands are clearly seen: the Lband (the excitation laser line at 488 nm), Raman D-band (shifted by 1368 cm1), G-band (with a Raman shift of 1557 cm1) and an overtone Raman band at 2930 cm1 were seen. The photoluminescence band maximum of the sample from the first experiment shifted to a short-wave region (570 nm).
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Fig. 6. Photoluminescence spectrum of the superhard phase after unloading: L—laser line of photoluminescence excitation at 488 nm, D—defect-induced D-mode, G—the G-mode band, PL—the photoluminescence maximum.
structure since they show only hk0 and 00l diffuse rings at around 0.3, 0.2 and 0.1 nm. The stripes in the HREM images demonstrate a different level of perfection in the packing of graphite planes. Usually, basal graphite planes are curved, and a frequent cross-linking between planes occurs. In Fig. 7a the HREM image of a typical structure, which was most frequently observed in the investigated samples, is seen. In this case the material is divided into nanometer scale domains with approximately parallel packing of graphite planes (graphene sheets) (Fig. 7b, HREM image) and with retention of a primary orientation of these planes in different domains at distances of the order of 1 nm (Fig. 7b, SAED). On the borders of these domains the packing of graphite planes (graphene sheets) is destroyed and strong buckling of these planes can be observed. Such a local texture can be related with the local primary orientation of the initial bundles of SWNT. Furthermore, graphite planes (graphene sheets) with ringtype and fingerprint-type shapes, which are typical for bucky onions, can also be observed in the HREM image of the structure of superhard carbon phase (Fig. 7a). The latter microstructure feature is similar to the one found in superhard carbon films produced by laser-arc method [31] and in carbon nitride films deposited by DC magnetron sputtering [32]. At the same time, a rough stripe contrast, observed in some particles at low resolution, allows to suggest that a complete destruction of nanotubes (in separate domains of the sample) did not occur even at very high pressures. The hardness and the Young modulus of the samples, synthesized at the pressure of 60 GPa after shear deformation application, were measured by NanoScan measurement equipment. The value of the hardness was found to be H ¼ 5876 GPa and the modulus of elasticity was as high as E ¼ 660760 GPa. The value of the hardness is close to the hardness of the samples of c-BN, measured by the same NanoScan measurement equipment and by the sclerometric method as well. 4. Discussion of the results
Fig. 7. HREM images and SAED patterns (insets) of some particles (a, b) of superhard phase.
The TEM investigations of the superhard phase samples after unloading were carried out using an electron microscope JEM-2010. TEM results were obtained by the investigation of the particles, prepared by mechanical cleaving of the samples and deposited on a copper grid coated with an amorphous carbon substrate. The selected area electron diffraction (SAED) picture and high-resolution electron image (HREM) of the superhard carbon phase are shown in Fig. 7. The structure of this hard phase can be described as a disordered graphite-like carbon. The obtained SAED patterns correspond to the turbostratic
Let us start the discussion from the photoluminescence spectrum (Fig. 6). Photoluminescence of carbon is investigated in great detail in amorphous carbon, a-C:H ([33] and references there). Its origin is attributed to the formation of defective states. It was recently shown [34] that defects in a-C:H are mainly due to the formation of sp2-clusters. The sp2-clusters consist of short olefinic chains or aromatic rings, because of which an optical gap can opens up from 1.5 to 3.5 eV depending on the content of sp2-states [35]. Suggesting the same nature of photoluminescence in our samples, the photoluminescence maximum found at 1.9 eV (in Fig. 6) corresponds to an optical gap of E 04 ¼ 2:122:3 eV. Furthermore, the dependence of the content of sp2-states in a sample on the E04 value for a-C:H was earlier measured, according to which in our sample the content of sp2-states can be estimated to be about 0.4. The
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rest belongs to sp3-states. Since sp3-states are coupled with strong s-bonds, forming a network frame, the hardness of the sample should be naturally increased. According to our measurements, the hardness of the sample after unloading in the first experiment corresponds to H about 60 GPa, which is comparable to the hardness of carbo-boro-nitride, c-BN. Our conclusion of the large content of sp3-states agrees with the observed low-frequency shift of the G-mode in Raman spectra after pressure unloading (a shift from 1589 cm1 (in the initial HiPco-SWCNT) to 1534 cm1 in superhard phase). It is known that the G-mode frequency reduces with the increase of the sp3-states content. In our experiments the position of the G-mode in the samples after unloading depended on the pressure at which a shear deformation was applied. The higher the applied pressure was, the lower the frequency of the G-mode was found, in agreement with an increase of the content of sp3-states. The bandwidth of the G-mode depends on the concentration of defects in the sample. In case of amorphous aC:H this bandwidth did not exceed 100 cm1 and the growth of the width with the concentration of defects was mostly connected with the increase of the bond-angle disorder [36]. However, the bandwidth in our samples after unloading was noticeably larger than 100 cm1 (Fig. 3), which can be explained by the large distribution of cluster sizes. As mentioned above, the Raman spectrum intensity under high pressure increased more than ten times after shear deformation application. The Raman intensity can rise due to the increase of the optical gap (2.1–2.3 eV after unloading) and, accordingly, due to the resonance gain of the Raman spectrum. Earlier [37], the conclusion about opening of the optical gap in carbon nanotubes under pressures higher than 42 GPa even without shear deformation was also made based on the observed increase of electrical resistance. The relative intensities of the bands at 1350 cm1 and the G-mode in Raman spectra depended both on defects concentration (sp2-clusters), and on excitation wavelength li. The 1350 cm1 band is a defect-induced one (D-mode) and therefore its intensity must depend on defects concentration. On the other hand, the size of sp2-clusters can be distributed in a large interval. It is known from the investigations of polycyclic carbons, that a cluster energy gap is inversely proportional to cluster size. Therefore a resonance gain of the Raman spectrum takes place for different cluster sizes at the excitation with different wavelengths. The larger the excitation wavelength li is, the bigger is the size of the clusters, in which the resonance gain of the Raman spectrum happens. According to calculations [38], a D-mode intensity increases with a cluster size increase. Therefore the relative intensity of a Dmode with respect to the intensity of a G-mode must rise with the excitation wavelength, which is seen in Fig. 5. Earlier [39], it was observed in the spectra of disordered and amorphous carbon.
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We observed the dependence of the frequencies of both D- and G-modes on the excitation wavelength (Fig. 5). The dispersion of the G-mode is found to be the highest right after shear deformation application under high pressure, it reduced with the decrease of the pressure in the cell and was almost absent after complete unloading. On the other hand, the dispersion of the D-mode can be investigated only after unloading (in the second experiment). As it was earlier shown in the spectra of disordered and amorphous carbons [39], the dispersion of these modes depends on the degree of structure disorder. If the dispersion of the Dmode is highest for ordered structure, the dispersion of the G-mode appears and grows only with disorder increase. For example, dispersion of the D-mode is almost absent in amorphous a-C [39]. The dependence of D-mode frequency in graphites with defects on the excitation wavelength was explained within the framework of double Raman resonance [40], according to which a dispersion of 40–50 cm1/eV can be expected, which was indeed observed in the series of experiments. If the value of the dispersion of the D-mode also correlates with the ordering in the structure, then it can be suggested that some elements of the ordered structure retain in the superhard phase, because the D-mode dispersion after unloading was small, corresponding to only 15 cm1/eV. A weak dispersion of the G-mode was earlier observed in amorphous a-C [39]. The great G-mode dispersion in the superhard phase spectra (larger than 100 cm1/eV) under high pressure after shear deformation application (Fig. 5) can be the evidence of the high degree of disorder. The gradual decrease of the G-mode dispersion with decreasing pressure and almost complete disappearance of the dispersion after complete unloading can be due to some structural order in the superhard phase after unloading. 5. Conclusions It is shown that application of shear deformation under high pressure permits to obtain a new superhard carbon phase from single-wall nanotubes (made by HiPco method), remaining after pressure unloading at room temperature. The hardness of this phase is quite high, about 60 GPa, which is close to the hardness of crystalline c-BN, the second crystal after diamond by hardness. The luminescence spectra allowed to estimate both the magnitude of the optical band gap E04 and the content of sp2states (estimated to be about 0.4) in the new superhard phase. The formation of the superhard phase was accompanied both by a step-wise increase of the Raman intensity as a result of the shift of the optical gap into the visible region, and by essential broadening of the G-mode band because of the growth of the concentration of defects. The investigations of the dependence of the Raman spectrum on the excitation wavelength showed that the Gmode dispersion under high pressure right after shear deformation application was very significant, being evidence of a high structural disorder. The structural disorder
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decreased with decreasing pressure, which was also confirmed by observation of the D-mode dispersion in the superhard phase Raman spectrum after unloading. A large bandwidth of the G-mode after unloading is evidence of a much wider distribution of the bond-angle disorder as compared to polymer- or diamond-like amorphous carbons a-C:H. According to our TEM-investigations, the structure of the superhard phase is highly disordered, containing clusters (with sizes of about 1 nm) of parallel graphene sheets with high distortion at the borders and separate domains in which the graphene sheets form, typical for fullerene onions, and in which the destruction of nanotubes is incomplete. References [1] V. Aksenenkov, V.D. Blank, A. Borovikov, V. Danilov, K. Kosorezov, Phys. Dokl. 39 (1994) 700. [2] V.D. Blank, M. Popov, S. Buga, V. Davydov, V.N. Denisov, A. Ivlev, B.N. Mavrin, V. Agafonov, R. Ceolin, H. Szwarc, A. Rassat, Phys. Lett. A 188 (1994) 281. [3] M. Popov, Y. Koga, S. Fujiwara, B.N. Mavrin, V.D. Blank, New Diamond Frontier Carbon Technol. J. 4 (2002) 229. [4] V.D. Blank, S. Buga, G. Dubitsky, N. Serebryanaya, S. Sulyanov, M. Popov, V.N. Denisov, A. Ivlev, B.N. Mavrin, Phys. Lett. A 220 (1996) 149. [5] F.P. Bundy, J.S. Kasper, J. Chem. Phys. 46 (1967) 3437. [6] T. Yagi, et al., Phys. Rev. B 46 (1992) 6031. [7] A. Goncharov, JETP 71 (1990) 1025. [8] U.D. Venkateswaran, M.E. Gosselin, B. Postek, D.L. Masica, G. Chen, R. Gupta, P.C. Eklund, Phys. Stat. Sol. (B) 235 (2003) 364. [9] U.D. Venkateswaran, A.M. Rao, E. Richter, M. Menon, A. Rinzler, R.E. Smalley, P.C. Eklund, Phys. Rev. B 59 (1999) 10928. [10] C. Thomsen, S. Reich, H. Jantoljak, I. Loa, K. Syassen, M. Burghard, G.S. Duelsberg, S. Roth, Appl. Phys. A: Mater. Sci. Process. 69 (1999) 309. [11] M.J. Peters, C.E. McNail, J.P. Lu, D. Kahn, Phys. Rev. B 61 (2000) 5939. [12] U.D. Venkateswaran, E.A. Brandsen, U. Schlecht, A.M. Rao, E. Richter, I. Loa, K. Syassen, P.C. Eklund, Phys. Stat. Sol. (B) 223 (2001) 225. [13] P.V. Teredesai, A.K. Sood, S.M. Sharma, S. Karmakar, S.K. Sikka, A. Govindaraj, C.N.R. Rao, Chem. Phys. Lett. 319 (2000) 302.
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Physica B 382 (2006) 58–64 www.elsevier.com/locate/physb
Nanostructured superhard carbon phase obtained under high pressure with shear deformation from single-wall nanotubes HiPco V.D. Blanka, V.N. Denisova,b, A.N. Kirichenkoa, N.A. Lvovaa, S.Y. Martyushova, B.N. Mavrinb,, D.M. Popovaa, M.Yu. Popova,b, E.V. Tat’yanina, A.A. Zakhidovc a
Technological Institute for Superhard and Novel Carbon Materials Troitsk, Moscow reg., Russian Federation b Institute of Spectroscopy of the Russian Academy of Sciences, Troitsk, Moscow reg., Russian Federation c NanoTech Institute and Department of Physics, University of Texas at Dallas, Richardson, USA Received 22 November 2005; received in revised form 30 January 2006; accepted 30 January 2006
Abstract Raman spectra of single-wall nanotubes under high pressure combined with shear deformation are investigated in situ in a diamond cell. Shear deformation applied under 35 GPa led to pressure multiplication up to 60 GPa, and increased the intensity of Raman bands more than ten times without essential change of G-mode position while causing its essential broadening. The G-mode remained broad after pressure unloading and shifted to 1534 cm1. The hardness of the superhard material was 5876 GPa, comparable to the hardness of carbo-boro-nitride. A broad band appeared in the photoluminescence spectrum with a maximum at about 2 eV, which allowed to assume a high content of sp3-bonds in the sample. The large dispersion of the G-mode almost vanished after pressure unloading. However, a noticeable dispersion of the D-mode was found, which is a sign of certain ordering in the superhard phase. TEM study of the superhard phase detected clusters with graphene sheets with a size about 1 nm. r 2006 Elsevier B.V. All rights reserved. PACS: 61.46.+w; 62.50.+p; 78.30.Ly Keywords: Nanotube; High pressure; Raman; Hardness
1. Introduction The carbon superhard phase synthesis is based on the transformation of the carbon sp2-hybridized states (which are the stable states of a two-dimensional carbon atom phase, such as graphene sheets ) into sp3-states. Atoms in a sp3-state are bound to s-bonds, which form a threedimensional network with high elastic characteristics. As usual, sp2–sp3 transformations take place under high pressures and high temperatures (for instance, phase transformations from graphite into diamond [1], or fullerite into superhard phase [2–4]). In Ref. [5] it was reported that a graphite single crystal, when compressed along the c-axis, irreversibly transforms
Corresponding author. Tel.: +7 095 334 0238; fax: +7 095 334 0238.
E-mail address: [email protected] (B.N. Mavrin). 0921-4526/$ - see front matter r 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.physb.2006.01.519
to hexagonal diamond at a pressure around 12 GPa and upon heating above 1000 1C. This phase transition was studied at room temperature using in situ X-ray measurements [6]. Under the experimental conditions of Ref. [6], this transition was reversible, i.e. upon pressure unloading hexagonal diamond transforms back to graphite. When graphite was compressed along the c-axis, the phase transition was detected at 14 GPa and room temperature [6]. The formation of sp3-bonds between graphene sheets at room temperature is sensitive to the stress tensor in the specimen. For example, at hydrostatic conditions graphite does not transform to diamond under pressures of at least up to 80 GPa [7], while the stress tensor variation (by applying a shear deformation under load) and a large plastic deformation lead to a direct irreversible transformation of graphite to hexagonal diamond at a pressure of 17 GPa at room temperature [1].
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As for carbon nanotubes under high pressure, Raman spectra [8–16], X-ray [17–19] and neutron [20] diffraction under high pressures were investigated. In the Raman spectra it was revealed that the intensity of the radial breathing mode vanishes at about 2 GPa, which was explained by either hexagonal [9] or oval [11] distortion of a cylindrical section or its collapse [16]. The theoretical investigations of the nanotube structure under increasing pressure came to the conclusion that a certain critical pressure exists, which is inversely proportional to the nanotube diameter, over which the cross-section of a nanotube becomes elliptical with subsequent collapse [21,22]. Recently it was shown that the pressure effects are reversible (at least, partially) on release of pressures as high as 25.9 GPa [13], 55 GPa [14,15] and 40 GPa [23], which were essentially higher than the critical ones. Possible polymerization of nanotubes was predicted in some papers. However, for the polymerization, pressures [24] higher than the critical collapse pressure were required [21,22]. In the papers [3,14,15] polymerization of nanotubes under the pressure of 24 GPa under conditions of plastic deformation was assumed. The conclusion about the polymerization was made on the basis of the conservation of the spectral features of nanotubes in the Raman spectra of nanotube samples after application of a high pressure of 30 GPa, while the hardness of the nanotube sample significantly increased, becoming comparable with hardness of c-BN. In the present study we have investigated Raman spectra of single-wall nanotubes, synthesized by HiPco method, in a diamond high-pressure cell with the simultaneous application of shear deformation. The photoluminescence spectra and the hardness of the samples after the loading were measured as well. We have found that the application of shear deformation under a pressure higher than 35 GPa allowed the formation of a new nano-structured carbon superhard phase, which remains stable after pressure unloading.
2. Experimental procedure The initial single-wall nanotubes (SWNT) made by HiPco (Carbon Technologies, Inc.) were cleaned from iron atoms by the method described in Ref. [19]. Although the HiPcoSWNTs contain a wider diameter distribution (0.7–1.4 nm), the nanotubes with the narrow distribution range of 0.88–0.91 nm diameters had contributed to the most intense reflections in the X-ray diffraction patterns [19]. All high-pressure experiments were performed using a shear diamond anvil cell (SDAC). Specimens were loaded into the gasket without a pressure transfer medium. In the SDAC a controlled shear deformation is applied to a specimen under pressure by rotation of one of the anvils around the load axis. The application of shear deformation decreases the hysteresis of the phase transformation and
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makes it possible to obtain a homogeneous phase (see, e.g. Ref. [3] and references therein). A spectrometer with three-fold monochromatization, a CCD-camera cooled by nitrogen and a confocal Ramanmicroscope with spatial resolution of 1–2 mm were used for the detection of Raman spectra, which were excited by Ar–Kr laser lines. The pressure inside the diamond cell was determined by a piezospectroscopy method, thus circumventing the uncertainty of the usage of rubyscale under the conditions, when the ruby crystal is pressed in the hard matrix of the sample [3,15,25]. The first-order Raman spectrum of diamond exhibits a single peak that corresponds to the triply degenerate optical phonons in the absence of stresses; the uniaxial stress causes splittings and shifts, which are linear with the stress. In the case of compression along the [1 0 0] diamond crystallographic direction, the single diamond peak shifts to the high-frequency region and splits into singlet and doublet modes. The mode splitting is proportional to the deviator part (pure shear) of the stress tensor, while the shift of the centroid of the split bands is proportional to the spherical part of the stress tensor. This effect can be used for measurements of stresses in diamond anvils (normally the diamond anvils are loaded along the [1 0 0] axis) and pressure in a sample compressed between the anvils [15,25]. The luminescence spectra were excited by a 488 nm line of a laser and were recorded by a TRIAX spectrograph. Hardness measurements were performed on the submicron length scale using the NanoScan (NS) measurement system based on the principles of the scanning force microscopy [26–28]. Earlier this measurement method for hardness was tested on other hard and superhard materials, including the hardness measurements of c-BN, diamond and ultrahard fullerite [27,28]. Nanoscan also allows to measure the Young modulus of the samples [29,30]. 3. Results of the experiment Investigating the Raman spectra, we concentrated the main attention on the spectral region above 1000 cm1, where the tangential modes of nanotubes’ sp2-bonds and their overtones were expected. In addition to the Raman band of a diamond anvil two overlapping bands at 1592 and 1520 cm1 (seen in Fig. 1), related to the tangential modes of sp2-bonds of the nanotubes, were seen in the spectrum at the initial stages of the high-pressure application cycle ðP ¼ 1 GPaÞ. Above 10 GPa the intensity of the nanotubes’ spectrum clearly was reduced and the bands transformed into one wide band, with a maximum shifted with the increase of the pressure in correspondence with the preceding measurements [14,15]. The shear deformation was applied by turning one of the diamond anvils relative to another anvil at the pressure of 35 GPa. Due to shear deformation application multiplication of pressure takes place [3], and the effective pressure in the sample increased from 35 to
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Fig. 2. Comparison of Raman spectra of HiPco single-wall nanotubes before shear deformation application (at P ¼ 35 GPa) and directly after application of a shear deformation. Fig. 1. Raman spectra of single-wall nanotubes (made by HiPco method) for increasing pressure before shear deformation application.
60 GPa. After that, the intensity of the Raman spectrum increased more than 10 times, the width of the maximum increased, but its position almost did not change. The background of luminescence noticeably increased, over which the Raman spectrum can be seen (as shown in Fig. 2). During the pressure reducing cycle, the Raman spectrum intensity and its following luminescence background decreased, and the maximum of the Raman G-band shifted to low frequencies (Fig. 3). After the complete pressure unloading one broad band with a maximum at 1534 cm1 and a very faint shoulder near 1350 cm1 remained. The pressure dependence of the high-frequency mode of Raman scattering (tangential mode at 1592 cm1 in the initial nanotubes; transformed to 1534 cm1 mode after unloading) probably reflects some structural changes in the sample. The results of these two experiments are presented in Fig. 4. The dependencies of Raman frequencies on pressure are almost identical for these experiments. Sharp turns on the pressure dependence curve during pressure loading (at Fig. 4) at 15 and 24 GPa were noticed earlier and have been attributed to structural changes [14,15]. At the phase transformation, a step-like anomaly
of the radial pressure distribution in the sample due to jumps in volume and elastic modulus appears in the case of non-hydrostatic compression. In Ref. [14], these anomalies were detected at pressures of 15, 19 and 24 GPa. In Refs. [14,15], the sharp turn at 24 GPa was attributed to a possible polymerization of nanotubes. We investigated the dependence of the Raman spectrum on the laser excitation wavelength li. To observe this dependence for the Raman line at 1350 cm1 as well, we carried out an experiment with shear deformation application at lower pressure (the maximum pressure after the shear deformation raised only to P ¼ 50 GPa). In this case the 1350 cm1 line was more intensive after unloading, but the character of the spectrum right after application of shear deformation remained the same as in the first experiment. First we investigated the dependence of the spectrum on excitation wavelength li right after the shear deformation ðP ¼ 50 GPaÞ. At excitation by the li ¼ 488 nm laser line, the maximum of the Raman band was at the frequency of 1697 cm1 (Fig. 5), which red-shifted with the increase of li (to 1665 cm1 if excited by 514.5 nm and to 1650 cm1 if excited by 647 nm). The frequency shift is quite sizable, larger than 100 cm1/eV. After pressure unloading two bands were seen in the Raman spectrum, which had a less strong dependence of the maximum position on excitation li (as shown in Fig. 5) shifts to 1557
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Fig. 3. Raman spectra of HiPco single-wall nanotubes for decreasing pressure after shear deformation application.
Fig. 4. Dependence of the frequency at 1592 cm1 on pressure for two experiments (33 GPa (m) and 35 GPa (K) before shear deformation application) during loading (filled signs) and unloading (open signs).
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Fig. 5. Dependence of Raman spectra of HiPco single-wall nanotubes on the excitation wavelength at the pressure of 50 GPa (after shear deformation application) and after unloading.
and 1368 cm1 at 488 nm excitation, and shifts to 1548 and 1335 cm1 at 647 nm excitation. The frequency shift of the 1350 cm1 line corresponds to 15 cm1/eV. Relative intensities of the discussed pair of Raman bands are found to be dependent on the excitation li. While the bands excited by the li ¼ 647 nm laser line were of about the same intensity, the band at 1368 cm1 was essentially weaker than the second band at 1557 cm1, if excited by li ¼ 488 nm (see Fig. 5). Photoluminescence spectra were obtained for the samples after unloading in the spectral region of 480–1000 nm. A very broad photoluminescence band with a PL maximum about 650 nm (1.9 eV) was seen in the photoluminescence spectrum of the sample from the second experiment (Fig. 6). On the wing of the PL spectrum several narrow bands are clearly seen: the Lband (the excitation laser line at 488 nm), Raman D-band (shifted by 1368 cm1), G-band (with a Raman shift of 1557 cm1) and an overtone Raman band at 2930 cm1 were seen. The photoluminescence band maximum of the sample from the first experiment shifted to a short-wave region (570 nm).
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Fig. 6. Photoluminescence spectrum of the superhard phase after unloading: L—laser line of photoluminescence excitation at 488 nm, D—defect-induced D-mode, G—the G-mode band, PL—the photoluminescence maximum.
structure since they show only hk0 and 00l diffuse rings at around 0.3, 0.2 and 0.1 nm. The stripes in the HREM images demonstrate a different level of perfection in the packing of graphite planes. Usually, basal graphite planes are curved, and a frequent cross-linking between planes occurs. In Fig. 7a the HREM image of a typical structure, which was most frequently observed in the investigated samples, is seen. In this case the material is divided into nanometer scale domains with approximately parallel packing of graphite planes (graphene sheets) (Fig. 7b, HREM image) and with retention of a primary orientation of these planes in different domains at distances of the order of 1 nm (Fig. 7b, SAED). On the borders of these domains the packing of graphite planes (graphene sheets) is destroyed and strong buckling of these planes can be observed. Such a local texture can be related with the local primary orientation of the initial bundles of SWNT. Furthermore, graphite planes (graphene sheets) with ringtype and fingerprint-type shapes, which are typical for bucky onions, can also be observed in the HREM image of the structure of superhard carbon phase (Fig. 7a). The latter microstructure feature is similar to the one found in superhard carbon films produced by laser-arc method [31] and in carbon nitride films deposited by DC magnetron sputtering [32]. At the same time, a rough stripe contrast, observed in some particles at low resolution, allows to suggest that a complete destruction of nanotubes (in separate domains of the sample) did not occur even at very high pressures. The hardness and the Young modulus of the samples, synthesized at the pressure of 60 GPa after shear deformation application, were measured by NanoScan measurement equipment. The value of the hardness was found to be H ¼ 5876 GPa and the modulus of elasticity was as high as E ¼ 660760 GPa. The value of the hardness is close to the hardness of the samples of c-BN, measured by the same NanoScan measurement equipment and by the sclerometric method as well. 4. Discussion of the results
Fig. 7. HREM images and SAED patterns (insets) of some particles (a, b) of superhard phase.
The TEM investigations of the superhard phase samples after unloading were carried out using an electron microscope JEM-2010. TEM results were obtained by the investigation of the particles, prepared by mechanical cleaving of the samples and deposited on a copper grid coated with an amorphous carbon substrate. The selected area electron diffraction (SAED) picture and high-resolution electron image (HREM) of the superhard carbon phase are shown in Fig. 7. The structure of this hard phase can be described as a disordered graphite-like carbon. The obtained SAED patterns correspond to the turbostratic
Let us start the discussion from the photoluminescence spectrum (Fig. 6). Photoluminescence of carbon is investigated in great detail in amorphous carbon, a-C:H ([33] and references there). Its origin is attributed to the formation of defective states. It was recently shown [34] that defects in a-C:H are mainly due to the formation of sp2-clusters. The sp2-clusters consist of short olefinic chains or aromatic rings, because of which an optical gap can opens up from 1.5 to 3.5 eV depending on the content of sp2-states [35]. Suggesting the same nature of photoluminescence in our samples, the photoluminescence maximum found at 1.9 eV (in Fig. 6) corresponds to an optical gap of E 04 ¼ 2:122:3 eV. Furthermore, the dependence of the content of sp2-states in a sample on the E04 value for a-C:H was earlier measured, according to which in our sample the content of sp2-states can be estimated to be about 0.4. The
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rest belongs to sp3-states. Since sp3-states are coupled with strong s-bonds, forming a network frame, the hardness of the sample should be naturally increased. According to our measurements, the hardness of the sample after unloading in the first experiment corresponds to H about 60 GPa, which is comparable to the hardness of carbo-boro-nitride, c-BN. Our conclusion of the large content of sp3-states agrees with the observed low-frequency shift of the G-mode in Raman spectra after pressure unloading (a shift from 1589 cm1 (in the initial HiPco-SWCNT) to 1534 cm1 in superhard phase). It is known that the G-mode frequency reduces with the increase of the sp3-states content. In our experiments the position of the G-mode in the samples after unloading depended on the pressure at which a shear deformation was applied. The higher the applied pressure was, the lower the frequency of the G-mode was found, in agreement with an increase of the content of sp3-states. The bandwidth of the G-mode depends on the concentration of defects in the sample. In case of amorphous aC:H this bandwidth did not exceed 100 cm1 and the growth of the width with the concentration of defects was mostly connected with the increase of the bond-angle disorder [36]. However, the bandwidth in our samples after unloading was noticeably larger than 100 cm1 (Fig. 3), which can be explained by the large distribution of cluster sizes. As mentioned above, the Raman spectrum intensity under high pressure increased more than ten times after shear deformation application. The Raman intensity can rise due to the increase of the optical gap (2.1–2.3 eV after unloading) and, accordingly, due to the resonance gain of the Raman spectrum. Earlier [37], the conclusion about opening of the optical gap in carbon nanotubes under pressures higher than 42 GPa even without shear deformation was also made based on the observed increase of electrical resistance. The relative intensities of the bands at 1350 cm1 and the G-mode in Raman spectra depended both on defects concentration (sp2-clusters), and on excitation wavelength li. The 1350 cm1 band is a defect-induced one (D-mode) and therefore its intensity must depend on defects concentration. On the other hand, the size of sp2-clusters can be distributed in a large interval. It is known from the investigations of polycyclic carbons, that a cluster energy gap is inversely proportional to cluster size. Therefore a resonance gain of the Raman spectrum takes place for different cluster sizes at the excitation with different wavelengths. The larger the excitation wavelength li is, the bigger is the size of the clusters, in which the resonance gain of the Raman spectrum happens. According to calculations [38], a D-mode intensity increases with a cluster size increase. Therefore the relative intensity of a Dmode with respect to the intensity of a G-mode must rise with the excitation wavelength, which is seen in Fig. 5. Earlier [39], it was observed in the spectra of disordered and amorphous carbon.
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We observed the dependence of the frequencies of both D- and G-modes on the excitation wavelength (Fig. 5). The dispersion of the G-mode is found to be the highest right after shear deformation application under high pressure, it reduced with the decrease of the pressure in the cell and was almost absent after complete unloading. On the other hand, the dispersion of the D-mode can be investigated only after unloading (in the second experiment). As it was earlier shown in the spectra of disordered and amorphous carbons [39], the dispersion of these modes depends on the degree of structure disorder. If the dispersion of the Dmode is highest for ordered structure, the dispersion of the G-mode appears and grows only with disorder increase. For example, dispersion of the D-mode is almost absent in amorphous a-C [39]. The dependence of D-mode frequency in graphites with defects on the excitation wavelength was explained within the framework of double Raman resonance [40], according to which a dispersion of 40–50 cm1/eV can be expected, which was indeed observed in the series of experiments. If the value of the dispersion of the D-mode also correlates with the ordering in the structure, then it can be suggested that some elements of the ordered structure retain in the superhard phase, because the D-mode dispersion after unloading was small, corresponding to only 15 cm1/eV. A weak dispersion of the G-mode was earlier observed in amorphous a-C [39]. The great G-mode dispersion in the superhard phase spectra (larger than 100 cm1/eV) under high pressure after shear deformation application (Fig. 5) can be the evidence of the high degree of disorder. The gradual decrease of the G-mode dispersion with decreasing pressure and almost complete disappearance of the dispersion after complete unloading can be due to some structural order in the superhard phase after unloading. 5. Conclusions It is shown that application of shear deformation under high pressure permits to obtain a new superhard carbon phase from single-wall nanotubes (made by HiPco method), remaining after pressure unloading at room temperature. The hardness of this phase is quite high, about 60 GPa, which is close to the hardness of crystalline c-BN, the second crystal after diamond by hardness. The luminescence spectra allowed to estimate both the magnitude of the optical band gap E04 and the content of sp2states (estimated to be about 0.4) in the new superhard phase. The formation of the superhard phase was accompanied both by a step-wise increase of the Raman intensity as a result of the shift of the optical gap into the visible region, and by essential broadening of the G-mode band because of the growth of the concentration of defects. The investigations of the dependence of the Raman spectrum on the excitation wavelength showed that the Gmode dispersion under high pressure right after shear deformation application was very significant, being evidence of a high structural disorder. The structural disorder
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decreased with decreasing pressure, which was also confirmed by observation of the D-mode dispersion in the superhard phase Raman spectrum after unloading. A large bandwidth of the G-mode after unloading is evidence of a much wider distribution of the bond-angle disorder as compared to polymer- or diamond-like amorphous carbons a-C:H. According to our TEM-investigations, the structure of the superhard phase is highly disordered, containing clusters (with sizes of about 1 nm) of parallel graphene sheets with high distortion at the borders and separate domains in which the graphene sheets form, typical for fullerene onions, and in which the destruction of nanotubes is incomplete. References [1] V. Aksenenkov, V.D. Blank, A. Borovikov, V. Danilov, K. Kosorezov, Phys. Dokl. 39 (1994) 700. [2] V.D. Blank, M. Popov, S. Buga, V. Davydov, V.N. Denisov, A. Ivlev, B.N. Mavrin, V. Agafonov, R. Ceolin, H. Szwarc, A. Rassat, Phys. Lett. A 188 (1994) 281. [3] M. Popov, Y. Koga, S. Fujiwara, B.N. Mavrin, V.D. Blank, New Diamond Frontier Carbon Technol. J. 4 (2002) 229. [4] V.D. Blank, S. Buga, G. Dubitsky, N. Serebryanaya, S. Sulyanov, M. Popov, V.N. Denisov, A. Ivlev, B.N. Mavrin, Phys. Lett. A 220 (1996) 149. [5] F.P. Bundy, J.S. Kasper, J. Chem. Phys. 46 (1967) 3437. [6] T. Yagi, et al., Phys. Rev. B 46 (1992) 6031. [7] A. Goncharov, JETP 71 (1990) 1025. [8] U.D. Venkateswaran, M.E. Gosselin, B. Postek, D.L. Masica, G. Chen, R. Gupta, P.C. Eklund, Phys. Stat. Sol. (B) 235 (2003) 364. [9] U.D. Venkateswaran, A.M. Rao, E. Richter, M. Menon, A. Rinzler, R.E. Smalley, P.C. Eklund, Phys. Rev. B 59 (1999) 10928. [10] C. Thomsen, S. Reich, H. Jantoljak, I. Loa, K. Syassen, M. Burghard, G.S. Duelsberg, S. Roth, Appl. Phys. A: Mater. Sci. Process. 69 (1999) 309. [11] M.J. Peters, C.E. McNail, J.P. Lu, D. Kahn, Phys. Rev. B 61 (2000) 5939. [12] U.D. Venkateswaran, E.A. Brandsen, U. Schlecht, A.M. Rao, E. Richter, I. Loa, K. Syassen, P.C. Eklund, Phys. Stat. Sol. (B) 223 (2001) 225. [13] P.V. Teredesai, A.K. Sood, S.M. Sharma, S. Karmakar, S.K. Sikka, A. Govindaraj, C.N.R. Rao, Chem. Phys. Lett. 319 (2000) 302.
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