Insights into the Early Growth of Homogeneous Single-Layer

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Insights in to the Early Growth of Homogeneous Single-Layer Graphene over Ni-Mo binary substrates Mark H. Rümmeli, Mengqi Zhen, Svetlana Melkhanova, Sandeep Gorantla, Alicja Bachmatiuk, Lei Fu, Chenglin Yan, Steffen Oswald, Rafael Gregorio Mendes, Denys Makarov, Oliver G. Schmidt, and Jürgen Eckert Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/cm4020783 • Publication Date (Web): 30 Aug 2013 Downloaded from http://pubs.acs.org on September 3, 2013

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Chemistry of Materials

Insights in to the Early Growth of Homogeneous Single-Layer Graphene over Ni-Mo binary substrates Mark H. Rümmeli*1,2, Mengqi Zeng3, Svetlana Melkhanova4, Sandeep Gorantla4, Alicja Bachmatiuk1,2, Lei Fu3, Chenglin Yan5, Steffen Oswald4, Rafael G. Mendes4, Denys Makarov5, Oliver Schmidt5, Jürgen Eckert4,6 1. IBS Center for Integrated Nanostructure Physics, Institute for Basic Science (IBS), Daejon 305-701, Republic of Korea 2. Department of Energy Science, Department of Physics, Sungkyunkwan University, Suwon 440-746, Republic of Korea 3. College of Chemistry and Molecular Science, Wuhan University, 430072 Wuhan, China 4. IFW Dresden, Institute of Complex Materials, P.O. Box 270116, D-01171 Dresden, Germany 5. IFW Dresden, Institute for Integrative Nanosciences, P.O. Box 270116, D-01171 Dresden, Germany 6. TU Dresden, Institute of Materials science, 01062 Dresden, Germany

Correspondence: [email protected]

Abstract The employment of Ni-Mo films has recently been shown to yield strictly homogenous single layer graphene. In this study we systematically investigate the different stages of nucleation and growth of graphene over Ni-Mo layers. The studies reveal that the Ni film breaks up and diffuses into the underlying Mo foil forming a Ni-Mo intermetallic. Nucleation only occurs from Ni sites and thus the nucleation density can be controlled by the Ni film thickness. Both nucleation and growth of the graphene are shown to be susceptible to very efficient selftermination processes to the formation of molybdenum carbide and this guarantees the formation large area graphene that consists entirely of monolayer graphene.

Keywords: Graphene, Synthesis, Nucleation, Growth

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Introduction

The intense interest in graphene's exciting properties has spurred a wealth of studies in graphene preparation. The primary routes are mechanical exfoliation [1, 2], chemical exfoliation [3, 4], reduced graphene oxide [5, 6], the bottom-up synthesis from molecular precursors [7], chemical vapor deposition (CVD)[8-10] and epitaxial growth over SiC. [11,12] In terms of preparing graphene for device applications, there is a need for graphene to be grown over large areas. Moreover, it should be highly crystalline with negligible defects and this should be achievable in a cost effective manner. [13] With the above prerequisites in mind, the most promising routes are the CVD fabrication of graphene and epitaxial graphene formed from the decomposition of SiC. In terms of quality, SiC has the edge over CVD. However, CVD is extremely well established in industry, is cost effective and is ideal for synthetic graphene fabrication over large areas. In the early days of the current drive in graphene it was the well-established transition metals known for graphite and carbon nanotube formation that were exploited. The primary transition metals of choice were Ni, Co and Fe. Early on there was a strong interest in Ni because apart from its well-known potential to yield high quality graphite [14] and form carbon nanotubes [15], the lattice mismatch between Ni (111) surfaces and graphene is less than 1 % viz. graphene is commensurate with the substrate lattice. Despite intense studies using Ni as a substrate to grow single and fewlayer graphene by CVD a fundamental limitation of Ni was soon established. In short, single and few-layer graphene is obtained over tens of microns and it is not homogenous across the substrate surface. The lack of control over layer number is attributed the high C solubility of Ni (0.6 weight % at 1326 oC). [16] Carbon and nickel form a solid solution above 800 oC and upon cooling (below 800 oC) carbon diffuses out. Moreover carbon segregation is rapid within grains and heterogeneous at grain boundaries making control of the carbon layer and layer homogeneity difficult. Copper has a significantly smaller carbon solubility (0.008 weight % at 1084 oC) [17] making control of layer number and homogeneity far easier and hence with the emergence and success of copper as a substrate for the CVD fabrication of graphene, Ni was soon abandoned as a major substrate for synthetic graphene. However, recently, Dai et al [18] designed a Ni-Mo combinatorial system that fully overcomes the fundamental limitations of Ni and yields homogenous single layer graphene over large areas. In this catalyst/substrate design, a thin layer of Ni is deposited over Mo. During the reaction the two metals blend. The technique is not only remarkable for its ability to yield homogeneous single layer graphene (better than Cu) with full coverage over the substrate


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surface, but also shows remarkable tolerance to variations in growth conditions. In this sense it may well be one of the easiest CVD routes to obtain graphene single layer graphene. In this study, we systematically investigate the blending of the Ni film with the base Mo substrate during the CVD process and correlate this with the graphene formation.

Experimental Substrate preparation: Ni films of 200 nm and 75 nm were grown onto 25 µm Mo foils using dc magnetron sputter deposition at room temperature. Ar was used as a sputter gas at a pressure of 10-3 mbar (base pressure: 10-7 mbar). Low Ni deposition rate of 0.6 A/s was applied to assure homogeneous growth of the metal film. The purity of the Ni target material was 99.95%. The CVD reaction: The CVD synthesis of graphene was conducted in a purpose built horizontal tube furnace. Initially the substrate was heated up to 500 °C in Ar (100sccm) then a flow of H2 (100 sccm) was additionally introduced. Heating to 1000 °C a rate of 20 ° per minute followed and this state was maintained for another 25 min at 1000 °C. After this CH4 was added with a flow of 50 sccm for different periods as discussed in the text. During cooling CH4 was maintained at a flow of 20 sccm with Ar and H2 for given temperature windows after which the CH4 was stopped and cooling thereafter was only with Ar and H2. The cooing rate was 20 °C per minute. An overview of synthesis steps are provided in figure 1 below. Graphene transfer: Transfer of the graphene on to Si/SiO2 and standard lacey carbon TEM grids was accomplished by first spin coating the target substrates with poly(methylmethacrylate) (PMMA) to form a film. The PMMA film was then released by immersing in to an aqueous solution of iron (III) chloride (FeCl3) for a few hours to etch away the Ni-Mo film. After transfer on to the target (Si/SiO2 or TEM grid) the PMMA was removed by exposure to hot acetone vapor. Finally the sample was annealed in high vacuum (10-6 mbar) at 200 oC to minimize residual surface material on the graphene. Characterisations: Raman spectroscopic measurements were obtained from graphene samples transferred on to Si/SO2 wafer using a Thermo-Fisher Smart Raman DXR Raman spectrometer using an excitation wavelength of 532 nm. These transferred samples were also examined using Atomic force microscopy on a Digital Instruments Veeco, NanoScope IIIa, in the tapping mode and a scanning electron microscope (FEI, NOVA NanoSEM 200). Samples

transferred on to standard lacey carbon grids were examined using a double Cs Corrected JEOL JEM-2010F working with an acceleration voltage of 80 kV equipped a Gatan UltraScan 1000



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slow-scan CCD camera, a Gatan DigiScan II digital beam control system, and a Bruker XFlash 5030T silicon drift detector.

Results In our CVD process prior to the introduction of the carbon source the Mo-Ni substrate is heated up to the reaction temperature over a period of 55 min in a Ar/H2 environment (see methods section for details) to 1000 oC. After the annealing treatment the carbon feedstock (CH4) is introduced and the CVD process begins. Initially we explored how the reaction time at a stable temperature affected the growth (∆t region as illustrated in figure 1a).

Figure 1. Schematic showing the two synthesis windows explored in the CVD process. a. Variation of time (∆t) for a stable upper temperature (1000 oC). b. Variation of cooling temperature window (∆T) during which CH4 continues to be supplied.

This was done for two different Ni film thicknesses (200 nm and 75 nm) deposited on a 25 µm high purity Mo foil. In all cases, including ∆t = 0, single layer graphene was found as determined by transmission electron microscopy (TEM) and Raman spectroscopic investigations (c.f. figure 2a). No trend could be discerned in the Raman signatures for graphene, that is, the D mode, the G mode and the 2D modes. The 2D mode showed a symmetric single Lorentzian line shape with a full width at half maximum less than 35 cm-1, and the intensity ratio of the 2D to G modes was ca. 2 (+/- 0.1). These values are concomitant with single layer graphene. The fact that no changes were observed in the obtained graphene for changes in ∆t from 30 min down to no stable temperature reaction window (∆t = 0), indicates that growth does not occur in this region but takes place during the cooling process.


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As the next step, we explored different cooling windows (∆T in figure 1b) before switching off the carbon feedstock (CH4). The cooling rate was 20 oC per min. Clear differences are observed that, as will be shown, confirm nucleation and growth of the graphene occurs during the cooling cycle and provide important insight into the dynamics of graphene CVD growth over binary Ni-Mo systems.

Figure 2. Raman Spectroscopic data. a. Typical Raman spectrum from homogenous monolayer graphene grown over Ni-Mo layer systems. b. Evolution of graphene Raman modes as the cooling temperature range during which CH4 is applied increases. The temperatures at which the CH4 flow is switched off after cooling down from 1000 oC are indicated. In this case the Ni film thickness was 200 nm. c. The same as for b, but for a Ni film thickness of 75 nm. d. variation of G to D mode with respect to cooling window for Ni films of 200 nm (black) and 75 nm (red). The ratio is determined from the area of the G and D modes. Note: the G to D ratio is measured as a function of peak area, not amplitude as is often found in the literature We begin by looking at the Raman spectra for samples cooled from 1000 oC in steps of 50 oC. In the case of a 200 nm film being deposited on Mo foil, no signatures for graphene are observed until a cooling difference of 150 oC is reached (i.e. from 1000 oC down to 850 oC) as shown in figure 2b. In the case of a 75 nm film, a cooling change of ∆T = 200 oC is required



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before the graphene modes appear in the Raman spectrum (figure 2c). In both cases the early formation of the graphene seems rather defective as can be seen by the large D mode (ca. 1350 cm-1) relative to the G mode (ca. 1590 cm-1). With longer cooling windows the D mode shrinks relative to the G mode as can be seen in figure 2d. This trend can be understood when looking at the graphene formation as it evolves over the different cooling windows.

Figure 3. AFM images of graphene evolution for different cooling windows. For narrow cooling windows only small graphene islands are observed. As the cooling window increases the islands grow and for sufficiently large cooling windows a full homogeneous layer of graphene can be observed.

Figure 3 shows atomic force microscopy images of the graphene samples for different cooling windows (∆T = 150, 200 and 300 oC) after transfer on to Si/SiOx substrates. For ∆T = 150 o

C only tiny graphene islands on the surface can be seen (in the case of the 75 nm Ni film

they are just visible). For ∆T = 200 oC, one can clearly observe larger graphene islands and by the time the cooling window reaches ∆T = 300 oC, in case of the 200 nm Ni film, full coverage is observed whilst for the 75 nm Ni film, near full coverage is observed. The flake size evolution for the samples (see table 1) explain the evolution of the observed G/D ratios with cooling window variation. In essence, the edges of graphene flakes contribute to the D mode [19,20]. As the flakes increase in size the relative signal contribution from edges decreases in a linear manner, exactly as we observe in our Raman data.


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Table 1 Average flake size Active temp. drop 200 nm Ni film 75 nm Ni film

ΔT = 300 °C Full coverage > 90% full coverage

ΔT = 200 °C 50 – 1000 nm flakes 50 – 700 nm flakes

ΔT = 150 °C 25 – 500 nm flakes 25 – 150 nm flakes

We now turn to the crystallinity of the graphene once full coverage has been achieved. To do this we evaluate the graphene through TEM after transfer to standard lacey carbon grids. Figure 4a shows a large region of transferred graphene. Figure 4b shows a selected area electron diffraction (SAED) image from the graphene shown in panel a. To confirm the observed reflexes are from mono-layer graphene and not from AB Bernal stacked graphene, we examine the relative intensity of the inner {10-10} and outer {11-20} spots. For monolayer the inner spots are more intense than the outer spots as we observe (figure 4c) confirming monolayer graphene [21]. We further confirm the presence of monolayer graphene by finding holes in the membrane to count the layers [8]. Only single layers were identified as for example shown in panels d and e from figure 4.

Figure 4. TEM characterisation of transferred graphene. a. Large region of (transferred) graphene resting on lacey carbon. b. Selected area electron diffraction pattern from the graphene. c. intensity profile over the {11-20} (outer) and {10-10} (inner) spots from the SAED pattern (in panel b) confirming monolayer graphene. d. and e. HRTEM micrographs of monolayer graphene with holes, confirming single layer graphene. f. magnified region of graphene from d. showing the typical honeycomb lattice from graphene. (note: the additional ACS Paragon Plus Environment


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material on the surface if the graphene e.g. in panels d and e is attributed to remnant material from the transfer process.)

One can analyse the graphene in greater detail to obtain statistical information on the grain range and mean size that form the monolayer graphene. This is done by piecing together TEM micrographs and looking at the diffraction information in the Fourier domain to determine the relative crystal orientations. The study shows that the graphene is turbostratic. Examples are provided in figure 5 in which the grains are highlighted in false colour. The grain size distributions are asymmetric (figure S2) and are analyzed using a positive skew function. The range and mean width are provided in table 2. The data reveal that the average grain size is larger for the graphene formed on substrates using a 75 nm Ni film and the basis for this is discussed further on.

Figure 5. Example of grain mapping of the graphene. The grains are determined from the reflexes obtained from the Fourier domain which indicate relative changes in crystallographic orientation. The grains are mapped in false color. Table 2 (median and average grain sizes) Ni film thickness 200 nm Ni film 75 nm Ni film

Grain size median (nm) 120 +/- 15 180 +/- 25


Grain size range (nm) 25 – 450 25 - 600

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We now turn to investigating the substrate surface to better comprehend the interaction between the Ni and the Mo. We do this by implementing Auger spectroscopy which is a powerful surface analytical technique which allows one to explore the first few atomic layers of a specimen with high spatial resolution. In this case we focus on the C KLL Auger spectra from the samples as this can provide important information on the bonding environment of the C species (e.g. sp2 carbon and carbides). Figure 6 shows the spatial variation for C, Ni and Mo over the surface of two post synthesis samples after cooling to 750 oC (left) and 850 oC (right) providing a sample with full graphene coverage and partial coverage, respectively.

Figure 6. Auger spectroscopy data. Comparison of a sample from a small cooling window (left) in which only graphene islands lie on the Ni-Mo intermetallic surface and a sample with a long cooling window (right) with full graphene coverage over the surface. Top: C KLL Auger spectra for different locations on the surfaces as indicated in the scanning electron images (SEM) (middle). The lower panels shows the elemental content along the red lines indicated in the SEM images.

The spatial data shows that regions richer in Mo tend to be deficient in Ni but with high contents of C. However, in places where the Ni content increases the Mo is seen to decrease,



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at the same time the carbon content decreases. These trends can be better understood by looking at the Auger peak profiles at these different regions. In Mo rich regions a negative peak just above 270 eV can be observed. In addition, three positive peaks at 247 eV, 255 eV and 263 eV can be identified. The peaks at 255 eV and 263 eV correspond to molydenum carbide. The peak at 247 eV also arises from molydenum carbide and overlaps with that from graphene or graphite. However, when graphene or graphite is present the intensity of the peak at 247 eV increases relative to that at 263 eV [22]. We observe this in the case of full graphene coverage (left) but not in the case of partial coverage. This suggests that no graphene resides over the Mo in the partially covered sample. For Ni rich regions the positive Auger spectra show a peak at 247 eV and no peaks at 255 eV and 263 eV concomitant with graphene over Ni [23, 24]. We can go yet further with our understanding by examining focussed ion beam prepared cross-section lamellas of the samples using energy dispersive Xray spectroscopy in scanning TEM (STEM). Figure 7 shows STEM images along with elemental mapping images (in false color). The cross-section images shows that the Ni film breaks up an diffuses in to the underlying Mo foil. During this process the two metals form an alloy in which there are clear boundaries between the Ni and Mo phases i.e. Ni islands embedded in the Mo form. The formation of these island explain the changes in Ni and Mo content observed in the spatially acquired Auger spectra. Moreover, the elemental mapping data (figures 7 and S3) show that with a thicker (200 nm) Ni film the number of Ni islands at the surface increases and the diffusion depth of these islands is greater as compared to the thinner (75 nm) Ni film. However, as discussed below, it is the number of Ni islands at the surface that is important and not the diffusion depth that is important, because the islands at the surface provide the nucleation sites for graphene growth and hence play a direct role in the grain density per unit area of produced graphene.


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Figure 7. STEM investigations of sample cross-sections (lamellas). Furthermost left: STEM images of sections of lamellas. Next to these are elemental EDS maps (in false color). We also explored the effect of the upper growth temperature for three temperatures, namely, 825 oC, 775 oC and 700 oC. In each case, the methane feedstock was supplied for 6 min as the system cooled down to 700 oC. In the case of the sample produced at 700 oC, no cooling is applied, instead the system was supplied CH4 for 6 min and then the methane flux was shut off and cooling in H2/Ar as for all other experiments was applied. In the case of the sample produced at 825 oC incomplete graphene coverage was obtained with 75% of the surface showing graphene coverage. For the sample produced at 775 oC, only small nucleation islands with dentritic morphology were obtained while for the sample produced at 700 oC only a few very small nucleation islands were obtained. Raman spectroscopic investigations of these samples shows a degradation in the quality of the material which is mainly attributed to the increase in open edges. In short that data (provided in the supplementary information figure S4) indicates one can possibly produce graphene at lower temperatures however the time required would need to increase significantly. This may also come at the cost of reduced quality.

Discussion The collective data enables a clear model of the nucleation and growth of the homogenous monolayer graphene to be formed. Spatial Auger spectroscopy and cross-sectional elemental analysis of the surface shows that during the synthesis route the Ni layer breaks up and intermixes through inward diffusion into the underlying Mo foil, thus forming a Ni-Mo intermetallic compound at the surface. This will take place during the initial heating up cycle ACS Paragon Plus Environment


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in hydrogen (see figure 1). With the thicker (200 nm) Ni films, the penetration depth and the lateral density of Ni grains embedded in Mo is greater than for the thin (75 nm) films. Once the reaction temperature has been reached and the feedstock supply initiated, there is no need to maintain a stable temperature, one can simply start the cooling cycle whilst keeping a flow of methane. By controlling the methane supply for different cooling windows one can explore the early stages of graphene growth, viz. nucleation through full coverage, over the substrate. The AFM and supporting Raman spectroscopic investigations show initially small graphene flakes or islands form on the surface (nucleation). Spatially resolved Auger spectroscopy studies showed these early nucleation islands of graphene form preferentially over Ni island rather than Mo. This is because the nickel carbide phase is unstable and so precipitates carbon on cooling whilst molybdenum carbide is stable and remains a carbide. With continued growth these graphene islands expand laterally and merge with other islands eventually forming a homogeneous single sheet of graphene. The crystal orientations of the Ni grains embedded in the Mo will vary and thus the relative orientation of the graphene nucleation islands differs and as they grow, meet and eventually merge this leads to the turbostratic graphene that we observe (figure 5). In the case of the thinner Ni film formed over the Mo foil the number of Ni grains at the surface is reduced in which leads to fewer graphene nucleation sites and hence larger grain sizes as we observe experimentally (table 2). The initial formation of a single graphene layer forming over Ni occurs because any excess carbon is soaked up by the Mo thus preventing multilayer graphene at nucleation. Once the island start to grow, the now graphene coated Ni grain cannot absorb any further carbon through the catalytic decomposition of methane since the graphene layer blocks this process. In effect it is self terminating. After nucleation, the catalytic decomposition of the methane feedstock occurs over unexposed Mo surfaces. As the methane is decomposed the produced carbon species are absorbed by the molydenum forming a carbide phase, or if in proximity of an open graphene edge add to the edge thus growing the graphene sheet. Once all the Mo surface is coated with graphene the catalytic decomposition of the methane ceases. Again, as with embedded Ni surfaces, the growth of graphene over Mo surfaces is self-terminating. In effect it is the selftermination of graphene formation over both the Mo and Ni sites that guarantees real homogeneous single layer. This is distinct to CVD growth of graphene over Cu, which although claimed as self-terminating more often than not results in a mix of 1, 2 and even 3 layer graphene [8]. It is also worth noting that once a graphene seed has been formed on the surface of the Ni-Mo(carbide) substrate, stable homogenous growth over the Mo (carbide) surface occurs with relative ease. Indeed, studies from the 1990s have shown that a variety of


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early transition metal carbides are suitable for the formation of graphitic layers suggesting a large family of carbides may be suitable for homogeneous graphene synthesis (see ref 25 and references within).

Conclusion

Systematic evaluations of as-produced graphene and the underlying substrates from CVD grown graphene over Ni-Mo layers confirm only single layer graphene is obtained over the entire substrate surface. The Ni film deposited over Mo foils are shown to break up and diffuse inwardly in the Mo foil forming Ni island embedded in Mo. Graphene nucleation is shown to occur exclusively over these Ni islands. Only single layers of graphene nucleate because the Mo which surrounds each Ni island soaks up any excess carbon species so the graphene formation of the Ni island is in effect self-terminating at one layer. These islands grow as methane is catalytically cracked over exposed Mo regions providing a source of carbon that is either absorbed to for molybdenum carbide, or if in proximity of a graphene edge add to the edge and hence grow the graphene. Thus, the growth of graphene over the Mo is also self-terminating. These self-terminating processes are responsible for the formation of homogenous single layer graphene.

Acknowledgements This work was supported by the Institute of Basic Science (IBS). MHR and LF acknowledge the

support from the Sino German Center (project GZ 871). RGM thanks the DFG (RU1540/8-1). We thank C. Krien (IFW Dresden) for deposition of the metal films.

Supporting Information Available: Supporting SEM data, statistical data and elemental mapping data are available. This material is available free of charge via the Internet at http://pubs.acs.org.

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Insights into the Early Growth of Homogeneous Single-Layer

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