Growth regimes and metal enhanced 6-fold ring clustering of carbon in carbon–nickel composite thin films

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Carbon 45 (2007) 2995–3006 www.elsevier.com/locate/carbon

Growth regimes and metal enhanced 6-fold ring clustering of carbon in carbon–nickel composite thin films G. Abrasonis a

a,*

, M. Krause a, A. Mu¨cklich a, K. Sedlac˘kova´ b, G. Radno´czi b, U. Kreissig a, A. Kolitsch a, W. Mo¨ller a

Institute of Ion Beam Physics and Materials Research, Forschungszentrum Dresden-Rossendorf, PF-510119, 01314 Dresden, Germany b Research Institute for Technical Physics and Materials Sciences HAS, Konkoly-Thege Mu´t 29-33, 1121 Budapest, Hungary Received 12 August 2007; accepted 17 September 2007 Available online 1 October 2007

Abstract Growth regimes of C:Ni (30 at.%) composite thin films grown by ion beam co-sputtering in the temperature range of RT-500 C are investigated. The combination of elastic recoil detection analysis, X-ray diffraction, transmission electron microscopy and Raman spectroscopy employing two excitation wavelengths was used to characterize the coexisting carbon and nickel constituents of the composite structure. Three growth regimes are identified characterized by different Ni nanoparticle shape (granular, columnar) and crystal structure (Ni3C or fcc Ni). The comparison of the Raman spectroscopy results from carbon reference and C:Ni (30 at.%) thin films shows that the presence of Ni enhances significantly the 6-fold ring clustering process at temperatures as low as RT, while at higher temperatures it favors ordering within the 6-fold ring clusters. The enhancement occurs independently on Ni nanoparticle size, shape or phase and is related to processes taking place on the surface of the growing film growth rather than in the bulk.  2007 Elsevier Ltd. All rights reserved.

1. Introduction The discovery of carbon encapsulated metal nanoparticles paved a new research pathway due to their unique structure [1–6] which results in distinct physical, chemical and mechanical properties [7–12]. Numerous metallic elements have been reported to be encapsulated in such cage structures [13], while the encapsulating property has been shown for pure carbon and other layered materials such as CNx [11] or h-BN as well [14–19]. Different methods have been reported to be able to produce such encapsulated structures based on arc evaporation [2,6,13,15,20–23], plasma enhanced chemical vapor deposition [24,25] or sputtering [9–12,17,26–29]. Of particular interest are continuous films of such encapsulated structures. Such films are essentially composites consisting of metal nanoparticles embedded in carbon or BN. Similar structures obtained after thin film growth *

Corresponding author. Fax: +49 351 260 3285. E-mail address: [email protected] (G. Abrasonis).

0008-6223/$ - see front matter  2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.carbon.2007.09.044

using different transitions metals as well as different encapsulating media (carbon, CNx or h-BN) suggest that there are common mechanisms which govern the formation of these structures. To get a comprehensive view on structure, properties and growth mechanisms, one needs to combine the techniques being appropriate to characterize both constituents of such composite structures, i.e., metal nanoparticles and the embedding media. The metal nanoparticles have been characterized usually employing X-ray diffraction (XRD) and transmission electron microscopy (TEM) [8–12,15,16,19,24,26,29–38]. The use of grazing incidence small angle X-ray scattering has been shown to be a successful method which allows probing simultaneously a large number of particles and gives a quantitative measure of particle shape and interparticle distance if an appropriate model is used to fit the experimental data [10,19,30– 33,39]. The presence or absence of foreign atoms within the nanoparticles has been revealed by extended X-ray absorption fine structure [30,32,33]. The encapsulating phase has been studied mostly by TEM to identify the presence of curved encapsulating

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sheets of layered material. However, the embedding medium does not exhibit regular layered structures and consists of a mixture of clusters of various extensions and curvature, which makes them difficult to be quantified by employing microscopic means. The complementary use of spectroscopic tools allows highlighting the structural features of embedding medium even if it appears microscopically featureless or amorphous. Concerning carbon, Raman spectroscopy is a powerful tool to characterize its different structural modifications [40–42], because of its sensitivity to chemical bonding, short and long range order, its high energy resolution, and its large penetration depth. In general, depending on the sp2 and the sp3 concentration, three classes of disordered and amorphous carbon can be distinguished, namely nanocrystalline graphite (ncgraphite), amorphous carbon (a-C) and tetrahedrically ordered amorphous carbon (ta-C) [43]. All three classes can be characterized by their Raman spectra. With decreasing sp2 content, the chemical bonding structure of carbon phases changes. In nanocrystalline graphite, a laterally confined graphite structure exists. The limited size of the graphite crystallites lifts the momentum selection rules for the Raman scattering, and phonons with a wave vector k > 0 can be excited. The Raman spectrum of nc-graphite consists therefore of the zone-center E2g phonon at 1582 cm1, which is called G-line, and a selectively enhanced disorder induced line, the so-called D-line. The intensity of the D-line ID is determined by the concentration of defects and the ring structures [43–45]. Due to the energy selective enhancement mechanism, the D-line frequency is dependent from the laser energy, and varies from 1325 cm1 for 785 nm excitation to 1370 cm1 for 488 nm excitation. The D-line is the fingerprint of graphitic carbon structures. The origin of the D-line and its dispersion have been successfully explained by a so-called ‘molecular’ approach [41,46] and solid-state double resonance approach [42,47]. In a-C, graphitic and olefin structures are present simultaneously. The G-line is therefore a superposition of the graphitic E2g line and C@C stretching vibrations of the olefin chains. With increasing chain length of the olefins, the p conjugation length increases and the absorption maximum gets red-shifted. Since the absorption maximum is correlated with the maximum of resonance Raman enhancement, the C@C stretching vibrations of longer olefin chains are observed with lower laser energy and vice versa. This is the origin for G-line dispersion in a-C. Moreover, the decreasing concentration of graphitic structures gives a smaller D-line intensity and a smaller ID/IG ratio. In ta-C, the olefin chains are the only sp2 carbon structure. Therefore, the D-line is not observed, and the G-line frequency depends on the excitation energy. Whereas, some studies of carbon–metal composites employed single wavelength Raman spectroscopy [28,29,36,48], to our knowledge, dispersive Raman spectroscopy has not been applied so far to these systems. A systematic comparison of carbon thin films grown with and without metal under otherwise identical conditions

would reveal the influence of the metal presence on the formation of carbon nanostructures which is of more general relevance. This may assist the understanding of the catalytic action of different metals in the formation of graphitic structures – a topic which is still controversial [49–52]. In this work, we present a study of C:Ni (30 at.%) composite thin films investigated by combined elastic recoil detection analysis (ERDA), XRD, TEM and Raman spectroscopy. The C:Ni thin films were grown by ion beam cosputtering. Despite of its simplicity, this synthesis method allows a high and independent control over synthesis parameters such as sputtering ion type, flux, energy, substrate temperature. Ni was chosen to be co-sputtered with carbon because it has low but nevertheless sufficient affinity to carbon which results in the formation of metastable nickel carbide Ni3C [53]. Thus the observed dependencies may bridge those which are or may be observed in the case of metals which have strong affinities to carbon such as Ti or V, and of metals which are completely immicible with carbon such as Cu or Ag. In this study, three growth regimes of Ni nanoparticles in the temperature range of RT–500 C are reported. The comparison of Raman spectra obtained by employing two different excitation wavelenghts of carbon and C:Ni will demonstrate the promotion of the 6-fold ring clustering of carbon due to the presence of Ni, at temperatures as low as RT. Moreover, at the highest temperature used in this study (500 C), Ni enhances the graphitic ordering.

2. Experimental Carbon–nickel (C:Ni) composite thin films were grown by ion beam co-sputtering employing a 3 cm Kaufman ion source which generates an Ar ion beam directed towards a 4 in. pyrolitic graphite target at a distance of 18 cm. The ions impinged at an angle of 35 with respect to the target normal. A Ni stripe is attached on the graphite target with a length slightly exceeding its diameter. The Ni content in the grown films can be controlled by the width of the Ni stripe. A width of 3 mm resulting in Ni atomic ratio of 30 at.% was used for the following experiments. The base pressure was 1 · 104 Pa rising to 7.5 · 103 Pa during sputter deposition. The total ion beam sputtering current and energy were kept constant at 35 ± 1 mA and 900 eV, respectively. The substrates were located on a substrate holder facing the graphite target at a distance of 14 cm and were heated by a boron nitride heater placed on the back side of the sample holder. The temperature was monitored by a thermocouple in contact with the back side of the substrates. Before each deposition, the target was presputtered for 15 min, then a shutter placed in front of the sample was removed without interruption of the sputtering process, and the depositions were performed for 60 min. The films were grown on the substrate in the temperature range of RT-500 C. It should be noted that during the growth of the samples without additional heating the sample temperature increased slowly during the experiment and reached a value of 40 C at the end of deposition. For simplicity reasons, in the following these samples will be denoted as grown at RT. For lower temperatures (6200 C), the C:Ni films were grown on the natural silicon oxide layer being present on square (2 · 2 cm) Si(0 0 1) substrates, while for higher temperatures the films were deposited on thermally grown SiO2 (500 nm) on Si substrates of the same dimensions, to avoid Ni loss due to inward diffusion [54]. For TEM observations, thinner layers were

G. Abrasonis et al. / Carbon 45 (2007) 2995–3006 grown for 10 or 20 min. For comparative spectroscopic investigations, a series of samples were deposited at identical conditions but without the Ni stripe, being termed carbon reference films in the following text. The film areal density and the atomic ratio of the film constituents were obtained by elastic recoil detection analysis (ERDA). The measurements were performed with 35 MeV Cl7+ ions impinging at an angle of 15 relative to the film surface. The backscattered ions and the recoils are detected with a Bragg ionization chamber placed at a scattering angle of 30. Additionally, a standard Si-detector was located at a scattering angle of 38 for hydrogen detection. In this case, an aluminum foil was employed in front of the detector to stop heavier recoils and backscattered Cl7+ ions. Near-infrared (NIR) Raman spectra were measured on a Kaiser Microprobe Raman system (Kaiser Optical Systems Inc., USA), which was equipped with a long working distance Leica objective of 100-fold magnification, an 1800 line/mm transmission grating, and a Peltier cooled CCD detector. For excitation a fiber coupled 785 nm diode laser was used. The laser power on the samples was 10–15 mW. Visible Raman spectra were recorded on a T 64000 triple spectrometer (Jobin Yvon, France) coupled to a BH2 microscope with a 100-fold magnifying objective (Olympus, Germany) using the 488 nm radiation of an Ar+ ion laser Innova 305 (Coherent Inc., USA) for excitation. Here, the laser power on the sample was 7–10 mW. The spectral resolution was 4 cm1 for NIR Raman and 2 cm1 for visible Raman experiments. The scattered light was collected in a 180 back scattering geometry on both the instruments. No sample degradation was detected under these conditions. The Raman spectra were recorded from two to three different sample areas and averaged for data processing and analysis. Thin film morphology was studied by TEM using Philips CM-20, JEOL 3010 and Philips CM300 microscopes. The phase structure of nickel was derived by XRD by employing a D-5000 (Bruker-AXS) diffractometer with Cu Ka radiation (8048 eV) used in grazing incidence mode. The incident angle was 1, and sample rotation speed was 15 rotations per minute.

3. Experimental results 3.1. Thin film composition ERDA measurements (for the details, see Supplementary material, Section 1) show carbon film areal densities of 8 · 1017 and 7.5 · 1017 cm2 for pure carbon and carbon–nickel composite films, respectively. The lower amount for the latter is related to the fact that part of the carbon target was covered with the Ni stripe, thus exposing a smaller carbon area to the sputtering ion beam. The C:Ni film deposited at 500 C shows a little lower carbon areal density of 6.3 · 1017 cm2. The Ni atomic ratio in the C:Ni films is 30 at.%. In addition, some metal impurities (<1 at.%) were found in the pure carbon films, which probably originate from the carbon target holder made from stainless steel which is sputtered by the low-intensity halo of the Ar ion beam. All samples showed the presence of oxygen at 1 at.% and hydrogen at 1–3 at.% homogeneously distributed across the thickness of the films. An exception is the carbon thin film deposited at RT, where the measured H content is 4.3 at.%. These impurities are probably due to residual gas, which adsorb on the growing thin film and are trapped by subsequently adsorbed atoms originated from the sputtered target. 3.2. XRD Fig. 1 presents the X-ray diagrams of C:Ni (30 at.%) composites grown at different substrate temperatures. It

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Fig. 1. XRD patterns of C:Ni (30 at.%) composite thin films grown at different substrate temperatures. The dotted lines indicate the fcc Ni peak positions.

should be noted that the peak structure at 54 is associated with the underlying Si. At the highest temperature of 500 C, the crystalline fcc metallic nickel phase can be clearly identified. With the temperature decreasing down to 400 C, the fcc Ni peaks decrease in intensity and broaden, and the features of another phase are observed, which become dominant at 300 C. This new phase has signatures of nickel carbide Ni3C (indicated on the figure). However, another nickel phase, the so-called hexagonal nickel, exhibits very similar X-ray diffraction pattern (see also [55,12]). In both phases, the structure of the Ni sublattice is identical. In Ni3C, carbon atoms are added into onethird of the octahedral interstices, which does not change considerably the atomic positions of Ni atoms but can significantly affect the electronic structure due to charge transfer and ligand-field splitting effects. Due to the relatively large widths of the peaks it is difficult to determine the nature of this Ni phase on the base of XRD alone. Additional spectroscopic diagnostics of the electronic structure would be necessary to identify this phase, which exceeds the scope of the present study. As it is often assigned in the literature as Ni3C phase, and as one can expect that in the presence of carbon this phase is more probable, it is labeled as Ni3C (see Fig. 1). When the deposition temperature further decreases, the sharp features in the XRD diagrams disappear. The broadness is a hint for a low degree of crystallinity of this Ni based phase. Besides, the maximum of this broad peak is shifted towards lower diffraction angles as that of fcc Ni. As Ni3C (or hexagonal nickel) phase has several peaks in this angle range, the broadening and

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overlapping of the peaks can result in an appearant shift of the overall feature towards lower diffraction angles. Thus, this broad peak is attributed to the Ni3C phase. It should be noted that the presence of Ni3C (fcc Ni) identified by XRD at lower(higher) temperatures in nickel containing composites has been also reported in Refs. [28,56]. In summary, XRD reveals three growth regimes of the Ni phase within the C:Ni composite thin films, with low degree of crystallinity at low temperatures (<300 C), crystalline Ni3C at 300 C and a predominantly fcc metallic Ni phase at >300 C. 3.3. TEM Fig. 2 shows the plan-view and cross-sectional high resolution TEM images of the thin film deposited at 300 C. The cross-sectional high resolution TEM image (Fig. 2(a)) reveals that Ni nanoparticles have formed, which exhibit columnar structure with the columns mostly extending over the thickness of the film. Similar columnar structures are observed for the samples grown at 400 and 500 C (not shown). In addition, the selected area diffraction patterns (for the details, see Supplementary material, Section 2) indicate that Ni exhibits Ni3C structure in the sample grown at 300 C. It is metallic (fcc) at 500 C. High-resolution plan-view TEM imaging (Fig. 2(b)) shows Ni3C nanoparticles of sizes in the range of 2–5 nm which are encapsulated in the graphitic planes. The latter appear as curved fringes which follow the boundaries of the nickel columns. At some locations only few graphitic planes between Ni columns can be observed, while at other locations the stacked graphitic planes of 3 nm thickness can be identified, depending on the interparticle distance. The latter indicates that most of the carbon is in graphitic state and fills the space between Ni3C columns. Cross-sectional high-resolution TEM (see Fig. 3) reveals a standing orientation of the graphitic planes perpendicular to the surface for the C:Ni thin film grown at 200 C. Extended graphitic planes are not observed for growth at <200 C while they

Fig. 3. Cross-sectional high-resolution TEM images of the C:Ni composite thin film grown at 200 C. Note that the fringes perpendicular to the substrate represent the standing graphitic planes.

are present in all the samples grown at higher temperatures (P200 C). Similar structures produced by magnetron sputtering have also been reported in the literature [11,12,29]. Closer inspection of Fig. 3 shows that the columnar structure of Ni3C nanoparticles is not continuous, but consists of small crystalline fragments which build up an irregular columnar structure different from that observed in the film grown at 300 C (see Fig. 2) or higher temperatures. This corroborates the XRD findings of different growth regimes in different temperature ranges. These regimes are also identified on low-magnification plan-view TEM images of samples grown at different temperatures, as shown in Fig. 4. Darker contrasts in the form of dots are observed for the films grown at 300 and 500 C (Fig. 4c and d). The presence of dots in the plan view image which, is a projection of 3D structure onto 2D, confirms the growth of the nanoparticles perpendicular to the substrate surface. However, blurred, non-regular contrasts are found for the film grown at 200 C (Fig. 4b). This indicates that the columnar growth is not preserved along larger heights.

Fig. 2. Cross-sectional (a) and plan-view (b) high-resolution TEM images of the C:Ni composite thin film grown at 300 C.

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Fig. 4. Plan-view TEM images of the C:Ni composite thin films grown at RT (a), 200 C (b), 300 C (c) and 500 C (d).

However, as the darker contrast due to the locations of Ni nanoparticles is relatively well separated from the carbon phase, one can conclude that also the interrupted growth continues preferentially near the sites of previously grown nickel. The contrast is less pronounced for growth at RT (Fig. 4(a)). Note also the coarsening of the grain structure as the temperature increases from RT to 500 C. In summary, TEM observations combined with selected area electron diffraction confirm the presence of Ni3C and fcc Ni observed by XRD in the samples grown at 300 and 500 C, respectively. In addition, Ni nanoparticles have fine-grained structure in the samples grown at <300 C, while at higher temperatures they have nanocolumnar shape with the particle height as large as the thickness of the thin films. Besides, graphitic carbon, consisting of curved standing planes encapsulating Ni nanoparticles, is formed at temperatures >100 C, while at lower temperatures carbon phase has a featureless appearance. 3.4. Raman spectroscopy Fig. 5 shows the Raman spectra of the carbon reference and C:Ni thin films in the wavenumber range 700– 1900 cm1 measured with 785 and 488 nm excitation light.

The D and G peaks can be observed in both types of films at all deposition temperatures, except for carbon reference films grown at RT and 100 C. For some spectra, the peak at 950 cm1 from underlying Si can be observed. For all the samples the Si peak at 520 cm1 (not shown) is observed for both excitation wavelengths, indicating that the full thickness of the films is probed by the present excitation wavelengths. To emphasize the effect of Ni on carbon, we discuss the features of Raman spectra of the carbon reference samples first. For the 785 nm excitation light (Fig. 5(a)), the Raman spectra of the films grown at RT and 100 C show very low intensity in the region of D and G peaks, and the spectra are dominated by Si Raman peaks and strong Si luminescence background starting at 1100 cm1. This indicates that the formation of NIR Raman active carbon structures is suppressed at low temperatures, while the almost constant Raman intensities indicate a similar film structure formation at deposition temperatures from 200 to 500 C. Besides, at P300 C, one can identify two distinct peaks in the whole D and G band which become clearly separated at 500 C. Using 488 nm excitation light (Fig. 5(c)), the intensity in the D peak region is significantly lower than that of the G peak region, while it increases concomitantly

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To obtain some more quantitative understanding of the temperature-dependent structural changes of the films, the Raman spectra were fitted by a combination of an exponential background, a symmetric Lorentzian for the D peak, and an asymmetric Breit–Wigner–Fano (BWF) shape for the G peak. The procedure was found to provide satisfactory fit of all spectra without restriction of parameters. The intensity I as a function of wavenumber k for the BWF line shape is described by [43,57] IðkÞ ¼

I 0 ½1 þ 2ðk  k 0 Þ=qC2 1 þ ½2ðk  k 0 Þ=C

2

;

ð1Þ

with the intensity maximum Imax I max ¼ I 0  ½1 þ 1=q2 

ð2Þ

positioned at [43] k max ¼ k 0 þ C=2q;

Fig. 5. Raman spectra of C ((a) and (c)) and C:Ni composite ((b) and (d)) thin films grown at different substrate temperatures measured with 785 and 488 nm excitation light. The spectra were vertically shifted, but the original intensity ratios for each measurement series are preserved. For the clarity reasons, the Raman spectra of C films grown at RT and 100 C measured with 785 nm are represented by a dashed and a dotted line, respectively.

with temperature. Two individual contributions from D and G peaks could be clearly separated except for the film grown at RT. The blue shift of the G peak with temperature at both excitation wavelengths should be noted. On the other hand, Raman spectra of C:Ni composites grown at low temperatures show an intensive contribution of the D band for both excitation energies (see Fig. 5(b) and (d)) in contrast to the carbon reference films. The total intensity of the D–G band increases concomitantly with temperature. Besides, the splitting of the D and G peaks into individual features become slightly observable at 400 C and can be clearly defined at 500 C. Here, at 500 C, the D peak is significantly more pronounced than in corresponding carbon reference film, with its intensity being significantly higher than that of the G peak. As in the case of the carbon reference films, the G peak shifts towards higher wavenumbers for both excitation wavelengths, concomitantly with temperature.

ð3Þ

where q is the BWF coupling coefficient. In the limit q ! 1, the Lorentzian line shape is reproduced, I0, k0 and C being Lorentzian peak intensity, position and full width at half maximum, respectively. In the following, the intensity maximum defined by Eq. (2) will be used to denote the intensity of the G and D peaks, the position defined by Eq. (3) as the G and D peak positions, the width of the G peak as the parameter C in Eq. (1) and the width of the D peak as the full width at half maximum. In the following, we will focus the discussion on the D and G peak positions and their intensity ratio ID/IG. However, for the sake of completeness, all the fitting parameters (the D and G peak positions, widths, coupling coefficients q and ID/IG ratios) are plotted in Fig. 6 versus deposition temperature for both C reference and C:Ni films. For the former, the D peak positions almost coincide at RT for both wavelengths and are equal to 1380 cm1 (the slightly higher value at 785 nm can be a fitting artifact due to low overall intensity of the D–G band, see Fig. 5(a)). At 785 nm, the D peak redshifts down to 1328 cm1 when the temperature increases from RT to 500 C, while at 488 nm the D peak position does not vary considerably with temperature. This means that the D peak dispersion increases from 0 cm1 eV1 concomitantly with temperature until it reaches a saturation value of 52 cm1 eV1 at 500 C. The above D position values should be compared to those of the graphite at given excitation wavelengths, which can be estimated by taking the D mode wavenumber at the K point equal to 1245 cm1 [58], which corresponds to the D mode frequency for zero excitation energy and a measured D mode dispersion value of 52 cm1eV1 [59]. This yields 1324 and 1372 cm1 for 785 and 488 nm, respectively. The comparison with the above fitting values shows that the D position and dispersion values of the C reference films resemble those of graphite only above 300 C. For the C:Ni films, the D peak position increases slightly from 1315 to 1328 cm1 and from 1364 to 1372 cm1 at 785

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3001

films grown at corresponding temperatures. Again, the G peak is almost dispersion-free at 500 C, and its position (1583 and 1588 cm1 for 785 and 488 nm, respectively) is very close to that of graphite. For the carbon reference films, the ID/IG ratio increases with temperature (from 0.25 and 0.7 at 785 and 488 nm, respectively) and reaches a plateau (0.75 and 1.4 at 300 C for 488 and 785 nm, respectively). For C:Ni, ID/ IG does not vary with temperature at 488 nm and is equal to 0.75, while it continuously increases for 785 nm from 1.1 at RT to 1.6 at 400 C. When the temperature increases further up to 500 C, there is no significant change. The ID/IG ratio dispersion increases with temperature in both types of films and reaches 0.7 and 0.8 eV1 at 500 C for C and C:Ni films, respectively. 4. Discussion 4.1. Structural implications from Raman spectra

Fig. 6. Position and width of the D and G peaks, BWF coupling parameter q and ID/IG ratio of C and C:Ni thin films grown at different substrate temperatures for 785 and 488 nm excitation light. The lines are the guides-for-the-eye only.

and 488 nm, respectively, when the temperature increases from RT to 400 C and drops at higher temperature by 8 cm1 for both excitation wavelengths. The latter is not observed for the carbon reference films. The resulting D peak dispersion is almost independent of temperature (it varies slightly in the range of 40–50 cm1 eV1), being close to that of graphite. For the carbon reference films, the G peak position increases with temperature up to 1580 cm1, from 1520 and 1560 cm1 at 785 and 488 nm, respectively. The small irregularity for the RT sample measured at 785 nm can be due to some fitting artifact because of the low intensity of the D–G band as mentioned above. The G peak is dispersive at all deposition temperatures except 500 C where its position (1581 and 1586 cm1 for 785 and 488 nm, respectively) is very close to that of graphite which is 1581 cm1 [40,58]. The variation of the G position of the C:Ni films with temperature is very close to that of pure carbon films at 785 nm, while the G peak dispersion is considerably lower in comparison to the carbon reference

It has been demonstrated in the previous section that the spectroscopic features of the Raman spectra from the carbon reference and the C:Ni films change differently when varying the deposition temperature. These different spectroscopic features imply that the carbon phase undergoes different graphitization paths. The total intensity of the D–G band of the 785 nm Raman spectra of reference films deposited at RT and 100 C is very low, and the spectra are dominated by Raman scattering of underlying Si and luminescence features, while the D–G band dominates the corresponding region of C:Ni films grown at similar temperatures. Raman scattering in carbons is a resonant phenomenon mediated by electronic transitions. Further, the sp2 cluster size is inversely proportional to the band gap [40,41,46]. At low temperatures, the low intensity from near-infrared excitation, which corresponds to a small band gap, indicates that larger clusters are essentially absent in the C reference films. In contrast, they deliver a significant signal from C:Ni films. At 488 nm where the clusters with larger band gaps (lower sizes) are probed, both carbon reference and C:Ni films show a significant Raman intensity. Concluding, carbon reference films deposited at low temperatures contain only small carbon clusters, while C:Ni films grown at low temperatures show the presence of both small and large clusters. The presence of the D band indicates the presence of aromatic rings in the sp2 phase, while the ID/IG ratio reflects the size of these aromatic clusters in disordered carbons [43]. The C:Ni films at both excitation energies show a significant contribution in the D band region, and significantly higher ID/IG ratios at RT-100 C (see Fig. 5(b) and (d) and 6). This indicates that the clusters probed with 785 nm C:Ni films are indeed 6-fold ring clusters and again confirms the previous statement that the average aromatic cluster size is significantly higher in the C:Ni films than in the carbon reference films. This has been reported in [29,48].

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From the G peak dispersion, the information about the topological disordering and presence of olefinic structures can be drawn [43], while its absence indicates a graphitelike structure. The lower dispersion of the G band for C:Ni films indicates that these films exhibit a higher degree of topological ordering with a lower concentration of olefinic structures than carbon reference films, even at low temperatures. Strong differences among both the types of films at low temperatures are also observed in the D peak position behavior. For carbon films deposited at low temperatures, the low D peak position dispersion measured by both excitation light energies indicates that the D peak represents more a density-of-states feature. The similar observation in sputtered amorphous carbon films from [60] was associated with the high degree of disorder present in the sp2 phase. The non-dispersive behavior of the D band contradicts both the double-resonance [47,42] and ‘molecular’ approaches [41,46]. From the structural point of view the latter approach is more relevant as the broad features in Raman spectra points more towards clustered and disordered material than towards the defected graphite. The ‘molecular’ approach assumes the uniform distribution of the aromatic cluster sizes and the contribution to the D line intensity only of the clusters which satisfy the resonant conditions [46]. According to this model, different excitation energies probe different clusters of different sizes thus of different D-mode frequencies. A higher D frequency is associated with smaller clusters. According to the calculations, this results in an apparant shift of the D line as a function of excitation light energy of 50 cm1 eV1 [46]. The absence or low D peak dispersion in carbon reference films grown at low temperature strongly indicates non-uniform cluster size distribution. According to the ‘molecular’ approach, in the case of non-uniformal distributions, the influence of the clusters with the dominant size will give the most significant contributions. Moreover, these contributions can be observed in the pre-resonant conditions. This would explain both the non-dispersive nature and low intensity at 785 nm. It should be noted that carbon films with high sp3 content prepared by pulsed filtered cathodic vacuum arc show no Raman response in the D– G band region at 785 nm1, which indicates that the ring structures if present are indeed too far away even from the pre-resonance conditions. Concerning the C:Ni films deposited at low temperatures, the D line is strongly dispersive, and this behavior does not change strongly with temperature. According to the ‘molecular’ approach this indicates that a more homogeneous distribution of cluster sizes is present in the C:Ni films as compared to the C reference films. One could argue that the presence of Ni could influence the optical process in such a way that even small clusters could be brought into resonance conditions for the NIR light. However, the D mode frequency close to that of

1

M. Krause, G. Abrasonis, L. Ryves, M.M.M. Bilek, unpublished.

graphite observed in all C:Ni thin films is more consistent with the assumption that large clusters are present in these films in sufficient amounts to yield strong Raman scattering intensities. The presence of the D peak dispersion of 50 cm1 eV1 for C:Ni films grown at low temperatures might indicate that the material consists either of 6-fold ring clusters with broad cluster size distribution or of defected graphitic-like structures. However, the significant G peak position dispersion eliminates the second possibility. Thus in summary, the intensity of the D–G band, the ID/IG ratio, and the G and D peak dispersion indicate that at lower growth temperatures the carbon films are highly disordered and contain only small aromatic clusters while the carbon sp2 phase of the C:Ni films consists of an ensemble of aromatic clusters with broad cluster size distribution, thus containing both large and small 6-fold ring clusters. The decrease in the G peak dispersion at increasing temperature indicates the graphitization of carbon in both types of films. At 500 C, the D and G positions are very close and similar to those of graphite. The G peak is almost dispersion free while the D peak shows a dispersion of 50 cm1 eV1. This indicates the dominance of graphite-like structures in both types of films. However, this state is achieved differently for the carbon reference and C:Ni films. For the former, the D peak dispersion gradually increases while the G peak dispersion gradually decreases, whereas for the latter the G peak dispersion decreases without significant variation of D peak dispersion. This indicates that the nucleation and growth of 6-fold ring clusters are significantly retarded in pure carbon films, while it becomes promoted by the presence of Ni. Following this, it can be concluded that Ni nanoparticles act as effective nucleation and growth centers of 6-fold ring clusters at low temperatures. However, at higher temperatures the carbon graphitization rate is determined mostly by the processes within the carbon phase. One could possibly assess and compare the total amounts of graphite-like phase in both types of films by comparing the total intensities of the D–G band for different excitation wavelengths. However, the optical properties (reflectivity, absorption) would have to be known also for the C:Ni films. Such data are not available to our knowledge. The observation that in C:Ni thin films the D peak dispersion does not change significantly while material evolves from an ensemble of clusters with different sizes towards a graphite-like material indicates that there is a transition from the regime where ‘molecular’ approach [41,46] is valid towards the regime where ‘solid-state’ [47] regime is valid. The ID/IG ratio which is associated with an average cluster size in highly disordered carbons [43] or the inter-defect distance in a graphite-like material [47] also changes significantly. In general, the ID/IG ratio (i) increases when the distance between defects (or cluster size) L decreases in graphitic materials as 1/L [44] and (ii) decreases concomitantly with degree of disorder as L2 for more disordered materials [43]. We do not attempt here to quan-

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tify L as our peak fitting procedure slightly differs from that reported in [43,60], but we mostly focus on the observed tendencies. The ID/IG ratio measured with 785 nm increases at increasing temperature for both types of films. This indicates that the observed tendencies should be associated with the second regime, where ID/ IG  L2 nm, i.e., higher ordering results in higher ID/IG. This regime is valid if the interdefect distances or cluster sizes are <2 nm [43]. This also agrees with the observation that the ID/IG dispersion increases with temperatures, which is associated with higher degree of ordering in disordered carbons [60]. However, the TEM images (Figs. 3 and 2) show that the graphitic planes are significantly larger than 2 nm already at 200 C. Therefore, the ID/IG ratios corresponding to distances <2 nm have to be attributed rather to defects within the graphitic planes. As the ID/IG ratio increases with temperature, this should be associated with the ordering within graphitic planes. The ID/IG ratio for lower temperatures (<200 C) might be associated with cluster sizes as TEM does not show any graphite-like structure. At higher temperatures, the ID/IG ratios become similar for both types of films at both excitation wavelengths, which indicates similar average cluster sizes for both types of films. In addition, the ID/IG dispersion saturates at higher temperatures. One can expect that at even higher temperature the ID/IG will start to decrease as a consequence of further ordering. Finally, in the case of C:Ni the G peak widths are similar and almost independent of the excitation light at 500 C, while the D peak widths are significantly lower (see Fig. 6). The widths of the D peak can be associated either to phonon life-times or to a distribution of D peak frequencies [43]. In view of the large D peak width (200 to 300 cm1), latter interpretation appears to be more appropriate. This is consistent with the picture that the presence of nickel at 500 C results in a formation of a more homogeneous graphitic structure with a lower amount of local distortions such as defects or odd-membered rings. This points out that Ni not only promotes the 6-fold ring clustering at lower temperatures, but also enhances the ordering within the graphitic phase at higher temperatures. 4.2. Growth regimes of C:Ni composites The morphology of a single component thin film is determined by nucleation and grain growth, which are dependent on the type of atoms and the growth conditions such as temperature and growth rate [61,62]. At low temperatures (or high growth rates), when surface and bulk diffusivity is low, nucleation dominates, thus the structure consists of fine grains. When surface diffusivity increases, lateral crystal growth sets in and the thin films exhibit the so-called T zone columnar structures [62]. At sufficiently high temperatures, bulk diffusion sets in and allows grain growth through grain boundary movement in lateral dimensions as well. It is known from the literature that at the initial stages of thin film growth, Ni forms islands

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which serve as a sink for the subsequently deposited Ni atoms [63,64]. Present TEM results show that at lower temperatures repeated nucleation of Ni prevents the formation of rod-like structure. XRD indicates very low degree of crystallinity accordingly. This is consistent with the low degree of mobility of the as-deposited metal atoms at lower temperatures. Thus, newly adsorbed Ni atoms tend to renucleate rather than to diffuse towards the existing islands. This is consistent to the so-called Zone 1 fine equiaxed grain structures [61]. The presence of defects in the carbon phase may also play a role as the investigations of the metallic cluster formation on amorphous carbon at RT show that as-deposited metallic atoms at low thin film coverages do not tend to form islands. At higher coverages the clusters grow by random adsorption of metals [65–67] and not due to surface diffusion. The former is saturated with defects which form the traps for the as-deposited metal atoms. Carbon mobility might also be low and results in the interruption of Ni3C nanoparticle growth. The increase in temperature results in the formation of columnar structure where the height of crystalline columns is similar to the film thickness which has been also reported by some of us [12]. This should be associated with the increase in the diffusivity of the adsorbed nickel atoms on the carbon phase surfaces which is sufficient to prevent the re-nucleation. The increase in temperature might also result in higher diffusivity of carbon atoms over the surface of Ni3C or Ni nanoparticles driving the excess carbon atoms to a surrounding C phase rather than creating the C phase on the growth surface of Ni3C. Besides, the carbon phase becomes more graphitic, which on its turn can further enhance adsorbed Ni atom mobility [63]. In this case, the nano-rod density and diameter are pre-determined by the nucleation and growth of Ni islands at the initial film growth stages. The thin film nucleation theory for single phase thin films [61] predicts lower nucleation density and formation of larger islands with the increase in substrate temperature which is consistent with the nanocolumn coarsening with increasing temperature observed in the present work (see Fig. 4). Following this, the kinetics of Ni phase morphology is qualitatively consistent with the single phase film growth. However, some additional factors come into play during composite structure formation. In particular for C–Ni system, nickel acts as a catalyst for the formation of graphitic structures [68,69]. Due to the non-negligible affinity of carbon to nickel, the constituents are not completely immiscible. Thus carbon atoms which land on Ni nanoparticles can form covalent bonds and participate in the formation of carbide phases. This phase is metastable and decomposes into its components at 700 C [70]. However, during dynamic conditions as in this work this phase becomes unstable at temperatures >300 C (see also [12]). This might be related with the increase of surface diffusivity and weak nature of Ni–C bond. Independently of the Ni particle growth regime (dominated by nucleation or surface diffusion), the carbon phase

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shows higher degree of 6-fold ring clustering or graphitization than that of the carbon reference films. At lower temperatures (<200 C), even the high resolution TEM observations do not show any long range ordering of carbon, Raman spectroscopy observations strongly indicate that some local 6-fold ring ordering takes place which is promoted by Ni. Thus, although surface diffusion is strongly limited and Ni nanoparticle growth is dominated by nucleation, it still enhances 6-fold ring clustering of the carbon phase significantly. At higher temperatures when surface diffusion becomes significant, both Raman spectroscopic analysis and TEM imaging show graphitic features in the carbon phase. Besides, Raman spectroscopy clearly shows higher degree of ordering within the graphitic phase in C:Ni than in carbon reference films. It is well known that Ni as well as other transition metals catalyze the graphitization [68,69,71]. However, the mechanism of the bulk ‘supersaturation–precipitation’ mechanism proposed in the literature [68,71] is not relevant to the results of the present study as the significant bulk diffusion of carbon atoms and precipitation on the nickel/nickel carbide nanoparticle surfaces would prevent the columnar growth which contradicts the present TEM observations. Besides, at lower temperatures, the surface diffusivity of Ni atoms is strongly limited thus bulk effects can be neglected. This strongly suggests that the surface diffusion is responsible for the observed enhanced 6-fold ring clustering. Recent in situ TEM investigations of the Ni nanoparticle dynamics deposited on the surface of amorphous carbon shows that Ni particles induce the graphitization while there is no evidence of carbon diffusion in the bulk of Ni nanoparticles [72]. The surface effects have also been underlined in the recent studies on the catalyst driven formation of carbon nanotubes [49–51]. At present, the exact mechanisms of such enhancement of carbon aromatic clustering and the function of Ni catalyst is still debatable. However, our present results clearly indicate that the enhanced formation of graphitic nanostructures occurs independently of the Ni phase state (carbidic or metallic) or independently of Ni nanoparticle size and shape. Besides, the graphenic planes follow the boundaries of Ni nanoparticles (see Fig. 2). This in combination with the non-interrupted nanocolumnar Ni or Ni3C structure points to migration of the excess of carbon atoms adsorbed on metal surfaces towards the boundaries of metal based nanoparticles where carbon precipitates in the form of graphitic phase. So-formed graphitic planes might provide a template of graphitic structure, and subsequent layers can build-up laterally, which is confirmed by TEM observations (see Fig. 2(b)) [10]. The fact that graphitic planes follow the surface of Ni in different phases (carbidic (or hexagonal Ni) and metallic) and the common observation in the carbon–metal and BN–metal composites that graphitic or BN planes follow the boundaries of nanoparticles for such different metals like Ni [11,12,29], Co [9], Fe [17,30,31], Cu [32] or Ag [10,19] points out that metal nanoparticles act more like a platform or a template for

the atoms of layered material [52]. This is also confirmed by recent investigations on Ti/C multilayers grown by pulsed filtered cathodic vacuum arc (FCVA) method. Carbon deposition onto an underlying titanium layer induces carbon graphitization (c-axes normal to the film plane) at the interface while subsequent deposition onto the carbon layer results in the formation of a fully amorphous structure2. In the same experiments, carbon deposition on SiO2 results in the formation of an amorphous structure, implying that formation of oriented graphite layers is associated with the presence of the titanium underlayer. The deposition of titanium on the carbon layer failed to induce graphitization of pre-deposited carbon layers, indicating that the catalytic action of the metal occurs only when carbon adsorbs onto the metal surface. At higher temperatures than those used in this study (>500 C) the formation of granular structure takes place and fcc Ni nanoparticles become completely encapsulated in curved graphitic sheets [11,12]. One can argue that bulk processes become dominant in that temperature range. However, the granular structure has been observed for silver [10] or copper [32] even at lower growth temperatures than in this work. The formation of graphite-like shell of several layers encapsulating the nanoparticles is accompanying such growth. At these temperatures, bulk diffusion effects are negligible. This phenomenon also occurs during annealing of deposited metal particles on the amorphous carbon surface [71,72]. The experimental data at the present do not allow unambiguous identification of the mechanisms responsible for such behavior. The low miscibility of encapsulating material and metal or the free surface energy might play a role. However, these assumptions remain speculative at the present and require further investigations. 5. Conclusion The C:Ni (30 at.%) composite thin films grown by ion beam co-sputtering have been investigated by the means of ERDA, XRD, TEM and Raman spectroscopy employing two different excitation wavelengths. The combination of these analytical methods allows characterization of both constituents – metal nanoparticles and embedding carbon. The obtained results show that in the temperature range of RT-500 C there are three different Ni growth regimes: • at low temperatures (RT-200 C) the formation of Ni nanoparticles is dominated by re-nucleation of subsequently deposited Ni atoms resulting in fine grained and interrupted growth of metal nanoparticles; • at 300 C the growth of Ni3C (or hexagonal Ni) nanoparticles is dominated by surface diffusion resulting in the formation of well-defined nanocolumns with high degree of crystallinity;

2 ˚ . Persson, L. Ryves, M. Tucker, D.R. McKenzie and M.M.M. P.O.A Bilek, Private communication.

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• at >300–500 C the growth is predominantly controlled by surface diffusion while Ni3C (or hexagonal Ni) becomes unstable and the fcc Ni phase becomes predominant. The use of two wavelengths for Raman scattering excitation has allowed revealing the following tendencies in the carbon phase. The comparison of Raman spectra of the thin carbon films with and without Ni clearly indicates that the presence of Ni enhances significantly the 6-fold ring ordering process at temperatures as low as RT, while at higher temperatures it favors the ordering within the 6-fold ring clusters. The enhancement occurs independently on Ni nanoparticle size, shape or phase. Both types of systems – carbon and carbon–nickel – approach the state at 500 C with the Raman spectroscopic features close to that of disordered graphite. However, the C:Ni approaches this state through intermediate states with higher topological ordering. The comparison with the results obtained on Ni nanoparticle growth regimes strongly suggests that the enhancement of carbon graphitization is pre-dominantly due to surface processes. Acknowledgements This work has been carried out inside the integrated EU project ‘Fullerene-based Opportunities for Robust Engineering: Making Optimised Surfaces for Tribology’ and supported by the EU contract no. NMP3-CT-2005515840. We thank Dr. Dieter Fischer (Institute of Polymer Research Dresden, Germany) and Prof. Dr. Lothar Dunsch (Leibniz Institute for Solid State and Materials Research Dresden, Germany) for the opportunity to use the Raman equipment in their laboratories and Dr. F. Munnik (Forschungszentrum Dresden-Rossendorf, Germany) for providing the ERDA depth profiles. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at doi:10.1016/ j.carbon.2007.09.044. References [1] Ajayan PM, Iijima S. Capillarity-induced filling of carbon nanotubes. Nature 1993;361(6410):333–4. [2] Ruoff RS, Lorents DC, Chan B, Malhotra R, Subramoney S. Singlecrystal metals encapsulated in carbon nanoparticles. Science 1993;259(5093):346–8. [3] Seraphin S, Zhou D, Jiao J, Withers JC, Loutfy R. Selective encapsulation of the carbides of yttrium and titanium into carbon nanoclusters. Appl Phys Lett 1993;63(15):2073–5. [4] Seraphin S, Zhou D, Jiao J, Withers JC, Loutfy R. Yttrium carbide in nanotubes. Nature 1993;362(6420). 503–503. [5] Seraphin S, Zhou D, Jiao J. Morphology and yield of carbon clusters in arc-discharge deposits. Carbon 1993;31(7):1212–6. [6] Dravid VP, Host JJ, Teng MH, Elliot B, Hwang JH, Johnson DL, et al. Controlled-size nanocapsules. Nature 1995;374(6523). 602–602.

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