Heteroepitaxial nucleation and growth of graphene nanowalls on silicon

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Heteroepitaxial nucleation and growth of graphene nanowalls on silicon Chia-Hao Tu a, Waileong Chen b, Hsin-Chiao Fang a, Yonhua Tzeng

b,c,* ,

Chuan-Pu Liu

a,d,*

a

Department of Materials Science and Engineering, National Cheng Kung University, No. 1, University Road, Tainan City 701, Taiwan Institute of Microelectronics, National Cheng Kung University, No. 1, University Road, Tainan City 701, Taiwan c Advanced Optoelectronics Technology Center, No. 1, University Road, Tainan City 701, Taiwan d Center for Micro/Nano Science and Technology, No. 1, University Road, Tainan City 701, Taiwan b

A R T I C L E I N F O

A B S T R A C T

Article history:

Heteroepitaxial nucleation of {0 0 2} graphene sheets on {1 1 1} facets of plasma treated (1 0 0)

Received 3 October 2012

silicon by direct-current plasma enhanced chemical vapor deposition in methane–

Accepted 15 November 2012

hydrogen gas mixtures is confirmed by high-resolution transmission electron microscopy.

Available online 27 November 2012

Lattice mismatch by 12% is compensated by tilting the graphene {0 0 2} with respect to silicon {1 1 1} and matching the silicon lattice with fewer graphene layers. The interlayer spacing of graphene sheets near the silicon surface is 0.355 nm, which is larger than that of AB stacked graphite and confirmed as AA stacked graphitic phase. Subsequent growth of standing graphene nanowalls is characterized by scanning electron microscopy and Raman scattering (633 and 514 nm excitation). The Raman peaks of D-band, G-band, and 2D-band are discussed in correlation with SEM images of graphene nanowalls. A strong Raman peak corresponding to silicon–hydrogen stretch vibration is detected by 633 nm excitation at the early stage of graphene nucleation, indicating the silicon substrate etched by hydrogen plasma. With these analyses, the growth mechanism is also proposed in this paper.  2012 Elsevier Ltd. All rights reserved.

1.

Introduction

Single-layer graphene is a stable one-atom-thick 2D sp2 hybridized carbon crystal in nature. Since graphene was first separated from graphite by exfoliation [1], a number of exciting properties of graphene have been discovered and experimentally confirmed, such as room temperature anomalous quantum Hall effect and high electron mobility [2]. Due to its unique electronic band structure [3], the band gap of graphene is tunable by electric field, which is promising for many future electronic application [4]. Besides mechanical exfoliation [1], many synthesis processes have been developed to grow graphene, including reduction from graphite oxide, thermally induced epitaxial growth on silicon carbide (SiC) [5], and chemical vapor

deposition on transition metals [6]. Mechanisms for epitaxial growth of graphene on silicon-terminated SiC have been rationalized as a two-step process. Firstly, surface reconstruction of SiC resulting from thermal decomposition of SiC and desorption of Si from the surface form a carbon buffer layer p p which is identified by a (6 3 · 6 3)R30 LEED pattern. Secondly, further annealing at a higher temperature leads to epitaxial growth of monolayer or bilayer graphene on top of the carbon interfacial layer [7]. Based on this growth mechanism, Suemitsu and Fukidome developed a technique to grow epitaxial graphene on a silicon substrate by depositing an ultrathin (100 nm) SiC film on a Si substrate followed by annealing at about 1500K in vacuum [8]. Besides in-plane growth, standing structure of multi-layer graphene (MLG) on a variety of substrates by catalyst-free

* Corresponding authors at: Department of Materials Science and Engineering, National Cheng Kung University, No. 1, University Road, Tainan City 701, Taiwan. Fax: +886 62346290. E-mail addresses: [email protected] (Y. Tzeng), [email protected] (C.-P. Liu). 0008-6223/$ - see front matter  2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.carbon.2012.11.034

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plasma enhanced chemical vapor deposition (PECVD) has been reported [9]. The MLG sheets stand roughly vertical to the substrate surface forming a wall-like network with corrugated surfaces in nature. As a result, this MLG network is also known as carbon nanowalls. A number of possible growth mechanisms for these standing MLG nanowalls synthesized by PECVD and sputtering deposition have been proposed. Among them, a three-step process was proposed as follows: firstly, fragmentary graphite base layers are formed in parallel to the substrate surface; secondly, many upward curling edges of graphite layers were formed by internal stress due to temperature gradient or lattice mismatch; finally, MLG sheets nucleate from these edges, radically changing their growth direction from parallel to vertical to the substrate surface [10]. The same group later also found that by angleresolved X-ray photoelectron spectroscopy (XPS), the presence of carbide and amorphous carbon in the initial stage of growth serves as an interlayer between graphite base layer and silicon substrate. This led them to revise the growth mechanism for growing MLG on top of SiC layers [11]. Another mechanism proposed that MLG sheets grow laterally first on silicon followed by a transition from lateral to vertical growth when two MLG sheets push each other under high pressures [12]. Other groups claimed the importance of built-in electric field between the plasma and the substrate surface, which enhances vertical growth of graphene layers [9]. Furthermore, epitaxial nucleation of MLG on {1 1 1} diamond facets was reported evidenced by high-resolution transmission electron microscopy (HRTEM) imaging [13]. Because of the complication of PECVD conditions involving bombardment by ions and electrons and reactions with radicals, atoms, and molecules, multiple routes to the growth of standing graphene cannot be ruled out. The growth mechanism depends on precise conditions of specific substrate and plasma. Multiple mechanisms leading to the growth of standing MLG nanowalls on the same substrate are also possible considering that the pristine substrate surface will be modified by the plasma and the deposition of different phases of carbon and its compounds before MLG nucleation is completed. Due to its ultra-sharp edges and high aspect ratio, MLG nanowalls are promising for electron field emission applications [14]. Yet, the characteristics of vertical arrays with high effective surface area enable MLG nanowalls to be especially ideal for applications in electrodes of fuel cells, electrochemical capacitors, and catalyst supports. In order to optimize the structure of MLG nanowalls for better performance, investigation into the interface between MLG nanowalls and the substrate is desirable. Direct imaging techniques with TEM is applied to elucidate the microstructure of the silicon–graphene interface. We unambiguously demonstrate the heteroepitaxial nucleation of MLG nanowalls on crystalline silicon where nanosized silicon {1 1 1} facets are exposed. The nanometer size of silicon {1 1 1} facets provides lateral space for mismatched graphene lattice to be compensated gradually with respect to the silicon lattice and form stable MLG nanowalls. MLG sheets nucleate at an optimal tilted angle with respect to the silicon surface initially and eventually grow into standing MLG nanowalls. In areas where conditions for heteroepitaxial nucleation of graphene on silicon are not favorable, amorphous carbon and clusters of silicon carbide are

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grown. Instead of heteroepitaxial nucleation, different graphene nucleation mechanisms suitable for the highly diversified silicon surface with different carbon and carbide deposits begin to play their roles. As a result, a high density wall-like standing MLG is grown on silicon.

2.

Experimental

2.1.

Growth of MLG nanowalls

MLG nanowalls are grown on 1 cm · 1 cm p-type silicon substrates by direct-current (dc) PECVD without catalysts. Prior to deposition, (1 0 0) silicon substrates are cleaned by acetone, ethanol, and de-ionized water in an ultrasonic bath. The plasma is ignited in hydrogen at 1 torr gas pressure. After about 3 min, it is followed by a CH4/H2 gas mixture at the flow rate ratio of 8/100 sccm at 120 Torr gas pressure. A dc voltage of 750–800 V is applied to produce a discharge current of about 1 A. This plasma is sufficient to heat the substrate to 800– 850 C as measured by an optical pyrometer. The substrate is placed on top of a molybdenum anode as shown by the schematic of the PECVD apparatus in Fig. 1.

2.2.

Characterization of MLG

Surface morphology of MLG is analyzed by scanning electron microscopy (SEM, Hitachi SU-8000), operating at 5 kV. Microstructure characterization is performed by HRTEM (JEOL JEM 2010F) operating at 200 kV. TEM samples are prepared by dual-beam focus ion beam milling. Chemical bonding of asgrown MLGs is studied by XPS (PHI 5000 Versa Probe). The carbon derivatives are verified by Raman spectroscopy measurements, performed with 633 and 514 nm lasers, a 100·

Fig. 1 – Schematic of a dc PECVD reactor.

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objective lens, and a liquid-nitrogen-cooled CCD camera detector.

3.

Results and discussion

The evolution of chemical bonding in MLG nanowalls grown for 3, 15, and 60 min is revealed by XPS C1s and Si 2p3/2 spectra in Fig. 2(a) and (b), respectively. The XPS spectra in Fig. 2(a) can be deconvoluted into four peaks: carbon in SiC (283.3 eV) [15], graphitic carbon (284.7 eV) [16], amorphous carbon (285 eV) [17], and carbon bonding to oxygen (286.2 eV) [15]. The Si 2p3/2 spectra in Fig. 2(b) can also be deconvoluted into four peaks, including elemental silicon (99.8 eV) [15], silicon– carbon bonding (101.3 eV) [15], silicon with two oxygen bonds (102.3 eV) [18], and silicon with four oxygen bonds (103.7 eV) [18]. Fig. 2(a) and (b) indicate that not only amorphous carbon and graphitic carbon, but also silicon carbide are formed on the silicon substrate in the early stage of MLG nanowalls deposition. The oxidization of the sample is responsible for the oxygen bonding with Si in the early stage and with carbon during the entire growth period. SEM images in Fig. 2(c) (e) shows the morphology of MLG nanowalls after different peri-

ods of growth. Fig. 2(c) reveals many nano-scale pedal-like structures as indicated by red arrows after 3-min deposition consistent with previous reports [10]. We suggest that the pedal-like structures keep on growing to become larger MLG sheets. As the growth proceeds, many sheet-like structures form and become bigger as shown in Fig. 2(d). Fig. 2(e) clearly exhibits similar morphology of MLG nanowalls to those reported in [9]. Thin MLG nanowalls are almost electrontransparent with a height of about several micrometers. Raman spectra excited by 633 and 514 nm lasers for the MLG nanowalls grown for 3, 15, and 60 min, corresponding to the SEM images in Fig. 2 on (1 0 0) silicon are shown in Fig. 3(a and b). After the silicon has been exposed to the CVD plasma for 3 min, the exposed fresh silicon surface is terminated by abundant atomic hydrogen ready for heteroepitaxial growth of graphene. A strong Raman peak (633 nm) at 2256 cm1 is clearly shown in the Raman spectrum. Raman scattering with Si–H stretch vibration appears to be dependent on the wavelength of the excitation laser [19]. Under excitation by 514 nm laser, the Si–H stretch vibration peak is not measured. Repeated measurements under excitations by lasers of these two wavelengths exhibit the same different

Fig. 2 – XPS spectra of MLG nanowalls grown for 3, 15, and 60 min: (a) C1s, and (b) Si 2p3/2 spectra; SEM images of MLG nanowalls grown for (c) 3 min, (d) 15 min, and (e) 60 min.

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Fig. 3 – Raman spectrum of MLG nanowalls grown for 60 min excited by (a) 633 nm and (b) 514 nm lasers.

Raman scattering due to Si–H stretch vibration. At this stage, only scarcely distributed nanometer sized MLG sheets are present on the silicon surface as shown in Fig. 2(c). Raman spectra (633 nm) mainly consist of three characteristic peaks, including D band at around 1330 cm1 from defective graphite such as graphene edges, G band at round 1580 cm1 from the radial breathing mode of sp2 carbon, and 2D band at around 2660 cm1 from double resonance with two phonon interactions [20]. As the growth time increases, the intensity of the G- and 2D-band becomes stronger while that of the D-band becomes smaller. The intensity ratio of the D- to G band, I(D)/I(G), is inversely proportional to the crystalline quality and should be zero for highly ordered pyrolitic graphite [21]. The domain size for 60 min sample is about 30 nm from I(D)/I(G), which is attributed to many nanometer sized graphene domains stacking together to form the vertical MLG nanowalls and is also partially contributed by the exposed edges of graphene sheets. However, a small diamond Raman peak at 1333 cm1 overlaps with the D-band and can be observed in the earlier stage of the MLG nanowalls growth. The PECVD conditions for MLG nanowalls are only

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slightly deviated from those for PECVD diamonds. As a result, diamond clusters composed of nanodiamond grains are present between MLG nanowalls, as indicated by red circles in Fig. 2(c). These diamond clusters contribute to the measured diamond Raman peak. As MLG nanowalls grow higher, these diamond clusters become surrounded by MLG nanowalls. The existence of these diamond clusters also complicates the analysis based on the I(G)/I(D) ratio. Because of the different resonance Raman scattering of diamond clusters and the MLG by 633 and 514 nm laser excitations, and the overlapping of Raman scattering signals from diamond clusters and MLG nanowalls, the estimation of the domain size of MLG nanowalls based on I(G)/I(D) is not precise. Furthermore, the I(G)/ I(D) ratios for 633 and 514 nm excitations are quite different because of different resonant scattering cross-sections by the G- and D-band by 633 and 514 nm excitations. Nevertheless, the low intensity of the D-band in comparison with that of the G-band indicates that the heteroepitaxially nucleated and grown graphene nanowalls highly graphitized. Unlike typical graphite, the shoulder peak at 2680 cm1 around 2D-band is missing. The 2D-band is symmetric, but with a much larger full-width at half-maximum (FWHM) of 60 cm1 compared to 34 cm1 for single-layer graphene. The broad 2D band is in agreement with MLG structures without AB stacking of 3D graphite. It might be of different stacking order such as AA or random stacking. Near the silicon surface, the heteroepitaxially nucleated graphene sheets have an interlayer spacing of 0.355 nm close to that of AA stacked graphite [22]. However, individual graphene sheets do not continue to grow vertically without defects. Instead, many defects cause the termination and re-nucleation of graphene sheets making the grown graphene nanowalls to consist of many small graphene domains of tens of nanometers in size. Edges of these nanometer sized graphene domains breaks the hexagonal symmetry required for 2D-band Raman scattering and, therefore, the low 2D-band intensity and the large FWHM value. Fully grown MLG nanowalls consist of curling graphene due to internal stress induced by defects in the MLG and the stacking disorder of graphene sheets. Without the guidance of a flat surface like the growth of few-layer graphene on the copper surface, the growth of standing graphene nanowalls relies on diffusion of carbon species from 3D space. The incoming carbon species find favorable spots to nucleate new graphene, while defects cause graphene growth to terminate. This process results in curling graphene nanowalls to be composed of non-AB stacked graphene sheets of nanometers to tens of nanometers in sizes. This is consistent with the I(G)/I(D) ratio and the wide FWHM of the 2D band of the Raman spectra for MLG nanowalls. HRTEM is employed to further investigate the interface between MLG nanowalls and the silicon substrate. Fig. 4(a) shows a bright-field cross-sectional TEM image of the MLG nanowalls grown on Si for 60 min. Apparently, the silicon surface becomes rough due to preferential etching of the Si (1 0 0) surface by atomic hydrogen and ion bombardment, resulting in nanoscaled Si {1 1 1} facets. Fig. 4(b) shows a higher magnification image of Fig. 4(a) at the interface, where bunches of curled MLG sheets surrounded by amorphous carbon can be observed to stick out from the silicon surface. Fig. 4(c) is the

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Fig. 4 – (a) A bright-field cross-sectional TEM image of the MLG nanowalls grown on Si for 60 min; (b) interface between MLG nanowalls and silicon substrate revealing that the silicon substrate was etched; (c) SAED pattern of the area indicated by white circle in (b); (d) HRTEM image of the enlarged area in (b) labeled by the white square, showing that MLG nanowalls grown on the surface of silicon substrate is embedded in amorphous carbon. (e) Schematic of AA stacked MLG nanowalls nucleated on Si (1 1 1) plane. (f) HRTEM image of a SiC cluster embedded in amorphous carbon.

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selected area electron diffraction (SAED) pattern from the area marked by a white circle. From the SAED pattern, the ˚ , these value spots yield d-spacing of 3.55, 2.12, 1.78, and 1.05 A are similar to the simulated diffraction data [22], based on the structure of AA graphite, corresponding to (0 0 2), (1 0 0), (1 1 0), and (2 0 0) planes. It indicates the presence of AA stacked graphite phase near the silicon surface, instead of AB stacked ˚ in (0 0 2) planes. The HRTEM graphite with d-spacing of 3.35 A image shown in Fig. 4(d), which is a higher magnification image of the white square region in Fig. 4(b), shows the MLG ˚ , corsheets containing lattice fringes with a spacing of 3.55 A responding to the {0 0 2} lattice spacing of graphite calibrated internally by using the lattice fringes of the silicon substrate. Moreover, Fig. 4(d) clearly reveals that the graphene {0 0 2} lattice fringes join directly with the silicon {1 1 1} lattice fringes. In order to accommodate a large lattice mismatch of about 12% between Si (1 1 1) planes and AA stacked MLG (0 0 2) planes, there is a transition area of 1 nm-thick between MLG sheets and silicon as indicated by two white arrows in Fig. 4(d). The transition area contains dislocations needed to accommodate the larger lattice spacing of C (0 0 2) than that of the Si (1 1 1) by 12%. For large-area (1 1 1) Si, such as a (1 1 1) wafer, the accumulated lattice mismatch of the very large number of graphene sheets would be very large and the internal stress would be more than what the interfacial transition layer can withstand. Therefore, unlike the demonstrated heteroepitaxial nucleation of graphene sheets on nanoscaled facets of {1 1 1} Si, heteroepitaxial growth of graphene nanowalls on large (1 1 1) Si wafer has never been successful. Nanometer sized Si {1 1 1} facets allows few layers of graphene sheets to heteroepitaxially nucleate on the (1 1 1) silicon surface. The number of graphene layers is thus controlled by the size of the Si {1 1 1} nano-facets. Fig. 4(e) is a schematic showing the AA stacked MLG nucleated on Si (1 1 1) plane with graphene layers tilted from the Si(1 1 1) plane by 163 to compensate the lattice mismatch. The HRTEM images and diffraction patterns indicate that graphene sheets in the MLG nucleate directly from the exposed {1 1 1} Si surface without SiC interlayer. Although the XPS data (Fig. 2) show the silicon–carbon bonding in the beginning of the growth, SiC clusters are found embedded in amorphous carbon in regions without MLG. Fig. 4(f) demonstrates such a typical example of a SiC cluster. The lattice ˚ , corresponding to cubic fringe spacing is measured as 2.5 A SiC {1 1 1} planes. The large lattice mismatch (34%) between SiC and graphene makes it unfavorable to retain SiC growth between graphene and Si. A similar case was found in ultrananocrystalline diamond films, where SiC was surrounded by amorphous carbon and the epitaxial relationship with silicon is SiC(1 0 0)[0 1 1] || Si(1 0 0)[0 1 1] [23]. Thus, we believe that SiC phase is formed of etched silicon and has nothing to do with the epitaxial growth of graphene on silicon. Based on the experimental observation, one of possible growth mechanisms for MLG nanowalls on silicon substrates is rationalized as follows: (i) Etching of silicon and formation of silicon {1 1 1} oriented nano-facets by hydrogen etching and ion bombardment during the initial hydrogen plasma treatment due to preferential etching of the Si (1 0 0) plane; (ii) AA stacked MLG nanowalls nucleate on selected silicon {1 1 1} nano-facets while the remaining silicon surface is

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deposited with amorphous carbon and some silicon carbide clusters. Because the lattice mismatch is relatively small between graphene (0 0 2) plane and silicon (1 1 1) plane, it is energy favorable for the observed heteroepitaxial growth. (iii) The growth of MLG nanowalls is enhanced because hydrogen etches amorphous carbon faster than graphite [24]. This phenomenon is also observed in the growth of diamond-like carbon film. (iv) Finally, the MLG nanowalls continue to grow at a much higher rate than that the surrounding amorphous carbon layer. The growth of MLG nanowalls is also enhanced by the high incident carbon flux from the plasma and the plasma induced electrical field between the growing MLG and the dc plasma, rendering MLG sheets to grow vertically to the substrate and stick out. The same growth mechanisms will be justified further by growing carbon nanostructures on different substrates in the future.

4.

Conclusions

In the conclusion, we have demonstrated a heteroepitaxial growth mechanism for MLG nanowalls to nucleate and grow on (1 0 0) Si substrate by dc PECVD. We attribute the heteroepitaxial growth of graphene nanowalls on Si to the fact that (1 0 0) Si substrate is etched by hydrogen plasma at a high temperature to expose Si {1 1 1} planes, which result in many dangling bonds for reacting with carbon-containing radicals. Raman scattering spectrum exhibits a wide and symmetrical 2D band. SEM images show curled standing graphene nanowalls. The synthesized graphene nanowalls are composed of graphene domains which stack on each other without following the AB stacking order of 3D graphite. Defects in graphene domains and those due to disordered stacking of graphene sheets induce internal stress which leads to the curling of the standing graphene nanowalls. The interface between MLG nanowalls and silicon has also been inspected by HRTEM, showing a transitional area with dislocations to compensate the mismatch between silicon and graphene. MLG nanowalls grown directly on silicon heteroepitaxially shows AA stacked graphitic phase in the region very near the nucleation site on {1 1 1} silicon facets. The epitaxial orientation relationship is graphene [0 0 2] || Si[1 1 1]. Thus, heteroepitaxial nucleation of graphene on silicon {1 1 1} facet is one of the possible graphene nucleation mechanisms for the synthesis of high-density standing graphene nanowalls on (1 0 0) silicon.

Acknowledgments The authors would like to thank Orlando H. Auceillo, Yuzi Liu, and Mengchun Pan for useful opinions, the Center for Micro/ Nano Science and Technology, National Cheng Kung University, Tainan, Taiwan, for access to equipment and technical support, and the NSC Core Facililies Laboratory for NanoScience and Nano-Technology in the Kaohsiung-Pintung Area. This work was supported in part by the National Science Council of Taiwan (Grant Nos. NSC-100-2221-E-006169-MY3, 101-2221-E-006-140-MY3, 100-2120-M-006-001 and 101-2911-I-006-517) and Ministry of Education, Taiwan. Use of the Center for Nanoscale Materials was supported by the US Department of Energy, Office of Science, Office of Basic

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Energy Sciences, under Contract No. DE-AC02-06CH11357. The electron microscopy was accomplished at the Electron Microscopy Center for Materials Research at Argonne National Laboratory, a US Department of Energy Office of Science Laboratory operated under Contract No. DE-AC0206CH11357 by UChicago Argonne, LLC.

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