Structural morphology of carbon nanofibers grown on different substrates

Carbon 98 (2016) 343e351

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Structural morphology of carbon nanofibers grown on different substrates S. Sainio a, H. Jiang b, M.A. Caro a, c, J. Koehne e, O. Lopez-Acevedo c, J. Koskinen d, M. Meyyappan e, T. Laurila a, * a

Department of Electrical Engineering and Automation, School of Electrical Engineering, Aalto University, Espoo, Finland Department of Applied Physics, School of Science, Aalto University, Espoo, Finland COMP Centre of Excellence in Computational Nanoscience, Department of Applied Physics, School of Science, Aalto University, Espoo, Finland d Department of Materials Science, School of Chemical Technology, Aalto University, Espoo, Finland e Center for Nanotechnology, NASA Ames Research Center, Moffett Field, CA 94035, USA b c

a r t i c l e i n f o

a b s t r a c t

Article history: Received 27 July 2015 Received in revised form 23 September 2015 Accepted 6 November 2015 Available online 11 November 2015

We present a detailed microstructural study comparing conventional carbon nanofibers (CNFs) and novel carbon hybrid CNF materials. The hybrid consists of CNFs grown on top of tetrahedral amorphous carbon (ta-C) thin films on silicon with nickel catalyst and Ti adhesion layers. The conventional CNFs were grown on silicon with nickel catalyst and Cr layers. Even though CNFs can be grown in both systems by tip growth, the micro- and nanoscale features are very different in the two systems. The crystalline structure of the CNF in the hybrid case changes from horizontal alignment to near-vertical alignment from the root to the tip and no bamboo structure is observed. The results show that micro- and nanoscale properties of CNFs grown under the same process conditions can be readily altered by using a sacrificial ta-C layer below the metallic layer to prevent the alloying of Ni with carbide-forming metals used as adhesion promoters and to act as an additional carbon source during the pre-annealing stage. The experimental results are further rationalized with the aid of assessed thermodynamic data and simulations based on density functional theory (DFT) with van der Waals (vdW) corrections. © 2015 Elsevier Ltd. All rights reserved.

1. Introduction Carbon nanofibers (CNFs) were first discovered in 1952 by Russian scientists [1] and have received renewed interest after the report on carbon nanotubes (CNTs) by Iijima in 1991 [2]. Carbon nanofibers and nanotubes have both been used in many applications such as energy harvesting [3], gas permeable electrodes in fuel cells [4], field-emitters [5], supercapacitors [6] and electrodes for biosensors [7e9]. CNFs are favorable in many applications due to their properties such as orientation, height, diameter, amount of sidewall defects and aspect ratio. Their properties can be altered by several parameters including the type of catalyst used as seed for the growth, nature of the growth method, growth temperature, carbon precursor chemistry and growth pressure. Feasibility of CNFs for the construction of biosensor electrodes has been widely studied previously [7e9]. There are also several

* Corresponding author. E-mail address: tomi.laurila@aalto.fi (T. Laurila). http://dx.doi.org/10.1016/j.carbon.2015.11.021 0008-6223/© 2015 Elsevier Ltd. All rights reserved.

studies on other carbon based materials such as tetrahedral amorphous carbon (ta-C) [10e12], ta-C þ CNT hybrid [13] and taC þ CNF hybrid [14], all reporting promising results. Thus, a detailed view of the structural features of these hybrid materials would be beneficial in trying to understand their unique properties. For example, a recent study by Robertson et al. [15] showed that MWCNTs can be grown from ta-C surface via bimetallic catalyst at around 500  C without affecting the ta-C underlayer. There was no direct evidence of the ta-C layer remaining unaffected, nor were cross-sectional transmission electron microscopy (TEM) images taken. However, there has been previous evidence through detailed cross section TEM observation that the ta-C layer is completely graphenised under very similar growth conditions [13]. Here we report results of a novel CNF growth method where the CNFs are grown on top of silicon substrates that are coated with taC thin film. Detailed investigation of the growth results along with in-depth structural characterization of the fibers with highresolution transmission electron microscopy (HRTEM), gives us the ability to discuss the structure and formation of the CNFs in

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considerable detail. Especially, the differences between the socalled bamboo type growth reported frequently in the literature and the markedly different growth mode observed here with the use of a ta-C interlayer are rationalized. Finally, we utilize assessed thermodynamic data together with detailed atomistic simulations to obtain a better understanding of the underlying reasons for the observed differences in the behavior between the two systems.

polishing. Polymer ink was used as the filler material with 70e80 nm of sputtered Pt and FIB deposited PteC to protect the CNFs during thinning. High-resolution transmission electron microscopy (HRTEM) was performed using a double-aberrationcorrected microscope JEOL 2100 (JEOL, Japan) equipped with Xray energy dispersive spectrometer (EDS) operating at 200 kV. A Gatan 4k,4k UltraScan 4000 CCD camera was employed for digital recording of the HRTEM images.

2. Experimental work 3. Results and discussion The deposition of the tetrahedral amorphous carbon (ta-C) was carried out as described in detail in Ref. [16]. Briefly, the process can be described as follows: p-type Si (100) wafer substrates (Ultrasil, USA) were cleaned by standard RCA cleaning procedure before deposition. The samples were mounted on a rotating holder (20 rpm) in a chamber with a base pressure of 103 Pa. The resulting taC film thickness was about 7 nm. Prior to deposition, the samples were etched by using a griddles argon ion source and coated with Ti using a continuous current arc source equipped with 60 magnetic filtering. The role of Ti is to ensure excellent adhesion of the ta-C layer on the substrate. Carbon nanofibers were grown on top of the ta-C by plasma-enhanced chemical vapor deposition (PECVD). First, a 20 nm thick Ni catalyst layer was deposited on the ta-C using cathodic arc deposition. The catalyst coated ta-C wafers were placed in a cold-wall PECVD reactor (Aixtron, Black Magic, Germany) with the chamber pressure pumped down to <1 Pa, and the samples were annealed at 400  C for 3 min before the growth process was started. The chamber was heated to 400  C by a ramp speed of 250  C/min. After the annealing step, NH3 buffer was used to fill the chamber (100 sccm) while the chamber pressure was maintained at 10 Pa. The temperature was increased to 750  C with a ramp speed of 300  C/min. After the temperature had reached 675  C, 150 W DC plasma was ignited while injecting the carbon precursor C2H2 to the chamber (30 sccm) and increasing the NH3 flow to 125 sccm. The growth phase lasted 10 min and produced vertically aligned fibers that were ~2 mm tall as shown in Fig. 1. After the growth process, the pressure was maintained below 20 Pa mbar until the temperature decreased to 300  C. The same process was used to grow the conventional CNFs on Si with ~80 nm Cr layer and ~24 nm Ni layer deposited via magnetron sputtering. The surface of the CNF layers was examined by scanning electron microscopy (SEM) (Hitachi S4800, Pleasanton, CA). Crosssectional TEM samples were prepared by focused ion beam (FIB) by using first 30 kV for thinning and subsequently 5 kV for final

PECVD produced films of vertically aligned CNFs on both surfaces mentioned above. For the conventional CNFs, the typical diameter of the nanofibers in the films varied between 20 and 120 nm, while their length was around 1 mm after the 10 min growth period. CNFs grown on ta-C were wider, ranging from 50 to 500 nm and shorter, being under 1 mm. The observed sidewall structure on the conventional CNF is consistent with the literature showing a clear, bamboo-like structure. From the TEM images in Fig. 1, the single “bamboo” units are crystalline in structure on their sidewalls and possibly hollow from the center. A notable difference can be seen for the CNFs grown on ta-C where the structure of the fiber is crystalline all the way from the substrate to the tip of the fiber, only having a small (approximately one fifth of the width of the fiber) center that is amorphous. The content is either from the growth process or from the filler material used in the TEM sample fabrication. The structure of ta-C-grown fibers appears like graphene sheets stacked on top of each other. Graphene sheet alignment remains horizontal for approximately ~80 nm from the base. Some areas of the edge of the fiber between the interface and the ~80 nm from the base show crystalline, roundshaped structures - about ~2 nm in diameter e that are aligned differently than the central part of the fiber (Fig. 2c). Following the fiber towards the tip above the ~80 nm from the base, the graphene sheet alignment changes to almost perpendicular to the substrate. TEM analysis also shows that at least some the fibers are not hollow, but filled with wrinkled graphene sheets (Fig. 3). There are also clear differences between the conventional CNFs and the ta-C-CNFs on the substrateefiber interface. In the conventional case, formation of some type of Cr-carbide layer, ~80 nm thick, on top of the Si substrate (Figs. 1a and 2a) with an uneven surface is seen and the fibers grow upwards from this layer. In contrast, there are no visible changes in the titanium layer (Figs. 1b and 2b) in the case of ta-C þ CNFs, and the EDS analysis did not

Fig. 1. Conventional CNF on silicon substrate with nickel catalyst and chromium underlayer showing clear bamboo-like structure along with the uneven chromium-carbon-layer on the siliconefiber interface. (b) CNFs grown on ta-C showing smooth sidewalls and no evidence of bamboo-like structure.

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show Ti-carbide formation. The titanium layer retained its original thickness (about 20 nm) throughout the process, but the ta-C-layer on top of the titanium completely vanished. The titanium layer, which works as an adhesion promoter for the ta-C layer coated on the Si, reacts with C less than Cr as seen from the above figures. This somewhat surprising result may be due to following reasons: (i) It is known that Ti has a high affinity towards oxygen and there is thus a large driving force for the dissolution of oxygen into Ti. (ii) Consequently, there is definitely enough oxygen in the deposition chamber during growth to completely saturate the Ti layer with oxygen, as the partial pressure needed for this at the growth temperatures used here is ~1038 Pa [17]. (iii) The presence of oxygen in the Ti[O]ss solid solution layer most probably stabilizes the film against carbide formation. Related to the last statement, we have previously shown that a TaC layer can be partly decomposed by the residual oxygen present in the deposition chamber; Ta is very similar to Ti with respect to its affinity towards oxygen, thus emphasizing the stable nature of the oxygencontaining solid solutions of these metals [18]. The higher magnification images of the fibers in Figs. 3 and 4 reveal that (i) the sidewalls of the single “bamboo unit” are crystalline and the rest of the fiber is amorphous, (ii) the ta-C þ CNFs are built from stacked crystalline sheets, having a completely different structure, and (iii) the CNFs grown on top of ta-C are not hollow. CNFs on both surfaces show tip growth as Ni particle can be found from the tip of the fibers. Based on EDS measurements, there are no clear traces of Ni remaining on the fiberesurface interface. The Ni particle is not completely covered with carbon in both cases as seen in Fig. 5. Thus, there is no poisoning of the catalyst particle and the growth could probably be continued for prolonged times. It is to be noted that when growing CNFs on ta-C films (with the Ni catalyst) with the same process parameters as used for the CNFs discussed above, but leaving the titanium layer out, the CNFs do not grow on the substrate at all. Thus, at least with the given set of parameters here, it is essential to have some metal layer on top of the Si substrate to promote adhesion of the CNFs to grow and stay on the surface. This is also the case when MWCNTs were grown on top of ta-C thin films [13,19]. In addition, the Ti film deposited on the Si substrate also improves the ta-C adhesion to the surface [20]. As seen from the extensive relevant literature, successful CNF growth can be achieved with a large parametric range for the growth variables and material combinations [21]. However, the resulting CNF structure from these growth conditions is often unclear. Based on the work by several groups [21e24] and especially that by Helveg et al. [25], the catalyst particle shape affects the resulting fiber structure significantly. Observation of the in situ growth of a CNF indicates that the catalyst particle protrudes from the surface after formulation, then stretches while protruding further away from the substrate and contracts back to a more rounded shape after the CNF reaches a critical length, when it becomes energetically too costly for the particle to gain anymore length [25]. When the elongation process is dominant, the resulting graphene sheets are vertically aligned towards the substrate and when the contraction of the particle takes place, the process produces the top portion finishing up the bamboo-like structure. The results here support that the growth of the conventional

Fig. 2. .TEM images. (a) shows ~80 nm thick, rough chromium-carbon compound formation on the Si-fiber interface. The chromium-carbon layer is about as thick as the original Cr layer but it has completely reacted during the process. In (b) the original titanium layer of ~20 nm is preserved. The ~7 nm thick ta-C layer has completely disappeared and there are no traces of amorphous carbon on the surface of the titanium. In (c) edge of ta-C-CNF, where there are small portions of the edge (marked with red arrows) aligned differently than the central area of the fiber. (A colour version of this figure can be viewed online.)

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Fig. 3. TEM images. Image on the left shows a fiber (on the center) that has been “burst open” from the interface-CNF area. Image on the right shows higher magnification image from the bursted area. It seems that the fiber is filled with graphitized, but wrinkled sheets.

Fig. 4. TEM images. The image on the left shows bamboo-like structure on the two fibers on the two sides and the larger fiber in the center showing highly defective sidewalls. From the higher magnification image on the right, it is possible to see the crystalline sidewalls of the larger fiber, the center of the fiber showing amorphous structure. Red lines on top of the image are: leftmost is marking the edge of the fiber and the two following the amorphous region are inside the fiber. (A colour version of this figure can be viewed online.)

Fig. 5. TEM images. (a) shows the nickel particle at the tip of a conventional CNF. The tip is little over half-way covered by graphene sheets. (b) shows the nickel particle at the tip of a CNF grown on ta-C. Here the graphene sheets cover less than half of the nickel particle. In both cases, the very top-most portion of the nickel catalyst particle seems to be free of carbon.

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CNFs proceeds approximately in the way explained above. However, the process appears to take a different route in the case of taC-grown CNFs, as the resultant fibers do not show bamboo-like structure. It is hypothesized that the presence of the ta-C layer under the Ni catalyst layer prevents the Ni particles from forming elongated shapes for the first ~80 nm, allowing them to remain more closely to the original shape immediately after formation (a semicircle droplet on the surface). This subsequently results in the graphene sheets forming a structure where they are horizontally aligned with respect to the surface. This issue will be discussed in more detail below. Helveg et al. [25] hypothesized that when the catalytic particle moves further away from the surface, the nickel particle starts to elongate, turning the forming graphene sheets to take more vertical alignment towards the substrate. However, in the case of ta-C, the particle never goes through this cycle, which would turn the particles first perpendicular and then back to the horizontal alignment with respect to the surface orientation of the forming graphene sheets. This was confirmed by imaging several resultant fibers, where all appear to have a structure where the first ~80 nm of the fiber is built by horizontally aligned graphene sheets and then the angle of the sheets turn to a more vertical alignment. Fig. 6 shows that the graphene sheets coming from the sides of the fiber to the center form a cone shaped structure and the middle part of the fiber has an amorphous (possibly hollow) center. The amorphous center area observed in the micrographs in Fig. 6 continues to around 80 nm height from the surface ending in a set of graphene sheets forming a semicircle. Below the semicircle, the graphene sheets are aligned horizontally towards the surface (Fig. 6) as already discussed. The structural differences observed between the ta-C grown CNF and conventional CNF are suggested to be mainly caused by the following reasons: (i) The Ni layer on top of ta-C will not form compounds or be alloyed with the Ti underlayer as there is a ta-C layer between them, which essentially acts as a sacrificial diffusion barrier. Thus, Ni particles remain relatively free from any major alloying elements or impurities except for carbon. However, it is expected that during the annealing stage (before the growth stage), Ni will dissolve at least some carbon from the underlying ta-C layer and becomes largely or perhaps completely saturated with respect to carbon, as also suggested in Ref. [26]. Solubility of ta-C to fcc-Ni is not known, but based on the stable solubility of graphite to fcc-Ni it can be estimated to be around 2e3 at-% at the present temperature

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range. This is because as ta-C is a metastable phase with respect to graphite under these conditions it has to have higher solubility to fcc-Ni than that of graphite. Based on the experimental information it is generally found that metastable solubilities are 2e3 times higher than stable ones [27]. Thus, this gives the above rough estimate. It has to be noted however, that these values are for bulk phases and somewhat higher values can be realized when dealing with nanoscale particles as here (see also [26]). It is known that the solubility of dissolving component (under equilibrium conditions) increases with decreasing particle size according to the OstwaldFreundlich type relations [28]. However, as long as the particle radius is not less that 5e10 nm this effect is not very strong. Formation of metastable Ni3C has also been reported when supersaturated fcc-Ni[C]ss samples have been splat quenched [29]. However, it is not very likely that such a phase would form under the present conditions. Finally, it is also known that the interaction between Ni and C is weak (as Ni is not a carbide forming metal) or even repulsive as Ni and C form an eutectic system. Based on what has been stated above we propose the following sequence of events during the growth process: When temperature is ramped up to growth temperatures there are two opposing effects: (a) solubility of ta-C to fcc-Ni will increase as temperature increases but at the same time (b) some of the ta-C starts to transform into graphite (so effectively to stacks of graphene sheets), which in turns decreases the amount of dissolved carbon in fcc-Ni, as the interface between the Ni-particles and underlying carbon phase tries to establish the local equilibrium conditions. Thus, it is suggested that some of the dissolved carbon is precipitated out of the Ni particles as “additional” graphene sheets underneath the Ni particles and contribute to the growth and morphology of the CNFs formed in this process. A slightly similar scenario has also been discussed in Ref. [26]. As the Ni particle is supersaturated with respect to carbon it is also expected that during the initial steps of the growth (when the carbon source is introduced to the chamber) carbon will mostly diffuse along the surface of the particles to contribute to the further growth of CNFs underneath the Ni particles. It is to be noted that as the temperature is high during these phenomena it can be assumed that the kinetics is not an issue here. After the local equilibrium conditions are reached at the Ni-particle graphene sheet interface (ie. most of the extra carbon has been precipitated out) the growth continues by the dissolution of carbon to fcc-Ni from the gas phase and the subsequent diffusion and formation of graphene sheets around the

Fig. 6. The micrograph on the left shows the lower portion of the Ni tip, on the left and right sides of the tip the graphene sheets are angled as the side of the catalyst is angled (around 20e30 towards the surface) and the micrograph on the right shows center area of the fiber where the graphene sheets are aligned as in (a), showing the amorphous center area. Red arrows point to the edges of the amorphous area. (A colour version of this figure can be viewed online.)

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particle under “steady-state” conditions. As the interaction of carbon saturated Ni particle with the graphene sheets around it is weak it will try to minimize its interaction with these. This can be achieved quite conveniently by orienting the (111) close packed planes of Ni to face the graphene sheets resulting more or less into the shape of the particles seen in the experimental results. (ii) In contrast, in the case of conventional fiber, the Ni and Cr thin films are not separated by a diffusion barrier and they can readily interact with each other. Perhaps more importantly, there is no carbon source to saturate fcc-Ni with respect to carbon during the annealing stage. Instead there will be interdiffusion between the metallic elements and a formation of Ni[Cr]ss solid solution phase. This is driven by the large solubility of Cr in Ni as shown in the corresponding binary phase diagram.1 The presence of Cr in Ni particles located at the fiber tips in this case was confirmed by EDS analysis. When the catalytic particles are Ni[Cr]ss, the growth process is seen to be the stretching-contracting sequence proposed by Helveg et al. [25]. We, however, propose a slightly different sequence of events. Stretching of the seed-particle in CreNi case is expected to be caused by the seed-particle adhering to the surface or to the newly formed “bamboo-unit” root because of the favorable interaction between Cr and C. The particle is “dragged” along from its other end by the growing fiber and by electrostatic forces induced by the plasma. The same phenomenon has been discussed in Ref. [22]. The particle is elastically and plastically deformed during the process reaching relatively high elongations. Thus, one is most likely dealing with superplastic behavior. It is known that Ni as well as NieCr alloys can experience superplastic behavior when the particle size is below 35 nm and the temperature is above 420  C [30,31]. Both conditions are readily fulfilled in the present case. Elongations up to 875% have been reported under these circumstances [30]. However, at some point, the adhesion strength at the lower particle/graphene interface will give up as there is a force asymmetry because of the plasma, and the particle will contract back to the more rounded (not the original one though as there has been permanent deformation of the particle) form finishing the cycle. When the seed-particle has lower interaction with carbon, as is the case with Ni, there is no elongation-contraction process as the adhesion strength is not high enough and the Ni just attempts to minimize the interaction with the resulting carbon fiber by changing its shape. Helveg et al. used MgAl2O4 as a substrate, and both Al and Mg are carbide forming metals; thus, there is a possibility that the Ni particles in Ref. [25] were not pure and contained some carbide forming metal resulting into the bamboo type of growth. It should also be noted that as there was no “extra” carbon source in Ref. [25] Ni particles would not be so easily saturated with respect to carbon. This in turn would increase the interaction (although weak) between Ni particles and defective graphene sheets in comparison to the case where Ni is fully saturated with carbon to start with. To verify and have a more quantitative support for the above arguments, we have utilized a combination of assessed thermodynamic data and detailed simulations based on density functional theory (DFT) [32,33]. From the thermodynamically assessed binary systems NieC, CreC and NieCr, the following data can be extracted. All the values are calculated at 750  C i.e. at the experimental temperature. The data shows that pure Cr has a very high attractive interaction with C (graphite in this case), which is close to 7.1 eV [34]. In contrast, Ni has a repulsive interaction with graphite (about 0.28 eV) [35]. This very large difference in the interaction between the two metals and graphite indicates that differences in adhesion

1 ASM Alloy Phase Diagram Database, http://www1.asminternational.org/ asmenterprise/apd/BrowseAPD.aspx.

strength between Cr and C and Ni and C are likely expected. Further, Ni and Cr have a slightly attractive interaction of about 0.06 eV and this value increases as the temperature is increased, as evident from the relevant phase diagram [36]. Based on these values, we can state that adding Cr to Ni should have a relatively clear effect by increasing the interaction with carbon atoms. Naturally, the values shown above are for bulk phases and they will not represent the interfacial nanoscale system very well. However, they clearly show a trend that is consistent with the arguments presented above. To deal with the “other extreme end” of energetic considerations, detailed atomistic calculations were carried out. The model system designed in order to obtain a good compromise between computational efficiency and representation of reality is a graphene sheet that interacts with a Ni or NiCr slab. The metal slab orientation is along the (111) direction of the facecentered cubic (fcc) crystalline phase. The evolution of the system's total energy as a function of graphene/slab separation is monitored for trends to model the interaction between graphene and metal. Finally, the experimental system at hand is hardly ideal, for example, the graphene sheets surrounding the metal particles are highly defective in the practical situation. Two different alignments are considered: i) “on top”, where the superficial metal atoms are placed right above the C atoms and ii) “interstitial”, where the superficial metal atoms are placed right above the middle of the interstitial space inside the C6 rings in the graphene sheet. The graphene sheet is a periodic structure made up of 4  4 primitive unit cells (32 C atoms in total) and the metal slab consists of 4 periodic hexagonal close-packed (hcp) planes with 16 metal atoms each, stacked with an ABCA sequence, corresponding to the (111) plane of the fcc lattice. The atoms in the graphene sheet and the top two layers (AB) of the slab are fixed, and the bottom two layers are allowed to relax freely as the distance between graphene and slab is changed. Three different Ni[Cr]ss slabs with 20 at-% of Cr have been generated, where the Cr atoms have been randomly substituted in the fcc lattice. This random sampling gives some insight about possible effects of Cr segregation. A suitable amount of vacuum is added to the simulation cell to prevent spurious interaction with periodic replicas along the direction perpendicular to the surface. This model system is schematically depicted in Fig. 7. The calculations were carried out using the plane wave-based implementation of DFT available from the VASP code [37] using the projector augmented wave (PAW) formalism [38,39]. To save computational time we did not include semicore states: only the 3d94s1 and 3d54s1 valence electrons of Ni and Cr, respectively, were included in the calculation. The 2s22p2 valence electrons were included for C. For the exchange-correlation part of the energy, we used the functional of Perdew, Burke and Ernzerhof (PBE) [40] for pre-relaxation of the system and subsequently we used the van der Waals-corrected optB88-vdW functional to obtain the geometry of the final system and total energy [41,42]. Sampling in reciprocal space was done using a Monkhorst-Pack [43] 4  4  1 mesh centered at the G point. A fully spin-polarized calculation was performed to account for the possible effects of unpaired d orbitals in Ni and Cr. The cutoff energy for plane waves was chosen as 400 eV. It has been previously shown [44e46] that the choice of DFT energy functional can have a significant impact on the results of adsorption energy of graphene on Ni surfaces; for example, the PBE functional tends to underbind whereas the local-density approximation (LDA) overestimates binding between graphene and a Ni (111) surface. Applying van der Waals corrections can help bring the calculations to better qualitative agreement with more sophisticated methods [44]. The optB88-vdW functional results have been tested to be in good agreement with more computationally

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Fig. 8. Interaction energy per metal surface atom as the distance between the Ni(Cr) metal slab and the graphene sheet is varied. (A colour version of this figure can be viewed online.)

Fig. 7. Simulation structures used in this work. A, B and C indicate the stacking sequence of the (111) planes in the Ni(Cr) fcc lattice. (A colour version of this figure can be viewed online.)

demanding schemes for this kind of systems, where the dispersive corrections due to the interaction between the p orbitals of graphene and the d orbitals of Ni lead to important contributions to the total energy. Therefore, we performed our calculations for the graphene/metal slab interaction using the optB88-vdW functional. The results from the calculations are shown in Fig. 8, which indicate that the alignment between the graphene sheet and the Ni(Cr) (111) surface strongly influences the interaction energy. The “on top” configuration provides a much less repulsive system than the “interstitial” configuration. This corresponds to the preference of the Ni atoms to avoid the interaction with C neighbors, as there is mutual repulsion. Recall that the NieC coordination is one and six for “on top” and “interstitial” configurations, respectively. This is in contrast to the CreC interaction: it can be seen from the bottom panel of Fig. 7 (which corresponds to the relaxed NiCr sample #3) that an increase in coordination is favored as Cr atoms at the surface relax to positions closer to the graphene sheet, compared to the

neighboring Ni which remain further apart. This behavior can be taken to indicate that segregation of Cr towards the graphene/alloy interface could be favored in the experimental system. The effect is so large that the last (interfacial) layer of NiCr atoms in “interstitial” configuration rearranges as part of the relaxation process so that it lies “on top”. Given the constraints present in the simulation, this gives rise to a modification of the stacking sequence where the stacking sequence ABCA of Fig. 7 is modified to ABCB, which is unrealistic. In the actual experimental setup two situations seem more plausible: either i) the alignment of graphene and metal would change at short distances so that the two systems “glide” with respect to each other and the “on top” configuration is fully retrieved (that is, with the original stacking sequence), or ii) the “interstitial” alignment would be maintained and the distance between the two systems would shift corresponding to the position of the minimum at approximately 3.3 Ångstrom (Fig. 8 bottom). We have calculated the energy cost of a stacking fault formation for the isolated pure Ni slab to be 24.6 meV per surface atom (obtained with the PBE functional). However, as seen from Fig. 8 (bottom), this relaxation is not observed in the pure Ni slab, although it would be energetically favorable. We attribute this discrepancy to the fact that the conjugated-gradient algorithm used to relax the metal slab might get stuck in a local minimum due to the higher slab symmetry which is broken by introducing randomly placed Cr atoms in the other cases. Another observation relates to enhanced binding as Cr is introduced in the slab. The three samples used, NiCr #1, #2 and #3, correspond to three different random configurations where the only constraint is to keep the overall composition of the four slab layers at 20 at-% Cr. This leads to a different number of surface Cr atoms in each sample, denoted by Cr* in Fig. 8. The use of different random samples with the same nominal composition gives insight to the possible effect of Cr segregation towards the surface, represented by configurations with more superficial Cr atoms. The increasing number of Cr atoms at the interface perfectly correlates with the lowering of the local energy minimum at around 2.2 Å observed for the “on top” configurations. Therefore, the binding

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between Ni[Cr]ss and graphene is expected to increase with Cr content, especially because the experimental growth conditions with a high temperature ~750  C readily allow Cr segregation towards the interface. Thus, based on both the assessed thermodynamic data from the corresponding systems and the atomistic simulations, one would expect the interaction between the Ni[Cr]ss and graphene sheets to be much more attractive than that of pure Ni and graphene. Finally, at some locations of the ta-C grown fiber edges below the semicircle formation, small ~2 nm diameter areas on the edges are not aligned horizontally, unlike the rest of the graphene sheets. These small portions of the edge are aligned roughly at 45 from the surface (Fig. 2c). Thus, it is possible that the Ni particle starts to change its shape on the fiber edges prior to the center of the fiber. Another interesting feature revealed by the TEM analysis is the seemingly full removal of the ta-C layer after the CNF growth process. The reasons for this are not unambiguously known at present. Some of the ta-C layer will be used for the growth of CNFs as discussed above. However complete disappearance of the ta-C layer is hard to explain only by its dissolution to fcc-Ni as this would require solubilities reaching tens of atomic percentages. Based on the earlier experiments with MWCNT growth on ta-C we know that NH3 alone at 550  C [13] does not affect ta-C layer. However, the combination of higher temperature and the presence of plasma in the present case may partly destabilize the ta-C layer, but this requires further studies. 4. Conclusions We have carried out a detailed study of the factors affecting the morphology of the novel ta-C þ CNF hybrid carbon material and compared that to CNF grown by a “conventional process”. It is shown that (i) the synthesis of the CNFs on the top of ta-C layer is feasible, (ii) nano- and microscale structure of the resulting fibers are significantly different between the substrates that have Ti þ taC þ Ni when compared to the Cr þ Ni, (iii) the Ti adhesion layer remains almost unaffected by the growth process and in particular there is no titanium-carbon compound formation, and (iv) the ta-C layer is fully vanished during the growth of CNFs. It is likely that taC layer is partly used for the CNF growth, thus resulting in different graphene sheet orientation between the two systems. The observed differences between the growth behavior in the two systems were rationalized by utilizing thermodynamic calculations combined with detailed atomistic simulations within the density functional (DFT) theory with appropriate van der Waals (vdW) corrections. We show that the alloying of Ni with Cr in the absence of ta-C sacrificial barrier leads into the conventional bamboo structure of the CNFs. However, when ta-C is present, it (i) saturates Ni particles with respect to carbon already during the pre-annealing and (ii) prevents the alloying of Ni particles with carbide forming metal(s) and the resulting CNF structure is distinctly different with no bamboo type features. Acknowledgments The authors acknowledge funding from the Academy of Finland (grant numbers 285015 and 285526) and the Finnish Funding Agency for Innovation (grant number 211488). Michael E. Salmon at Evans Analytical is acknowledged for the FIB sample preparation. Dr. V. Protopopova is acknowledged for fabrication of the Ti þ taC þ Ni substrates for the CNF experiments. This work made use of the Aalto University Nanomicroscopy Center facilities. The computational resources for this project were provided by the Finnish Center for Scientific Computing (CSC) through the Sisu € rn Bjo €rkman for supercomputer. M.A.C. would like to thank Torbjo

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