Catalyst-Free Direct Growth of Triangular Nano-Graphene on All

Jun 28, 2011 - The direct synthesis of graphene onto the desired substrates without using any metallic catalyst is free from the hurdle related to the...
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Catalyst-Free Direct Growth of Triangular Nano-Graphene on All Substrates Ki-Bum Kim, Chang-Mook Lee, and Jaewu Choi* Department of Information Display, Kyung Hee University, 1 Hoegi-Dong Dondaemoon-Gu, Seoul 130-701, Korea ABSTRACT: To epitaxially grow graphene, metallic catalysts or carbon containing silicon carbide have been typically utilized. The embedded metallic catalyst between graphene and the substrate as well as the expensive silicon carbide substrate create hurdles in the development of graphene-based devices. However, what is inevitably necessary is not a metallic catalyst but a flat plane able to hold the carbon species and to mediate their interaction on the plane. The plane needs neither to hold a large amount of carbon species nor be a highly efficient catalyst because one monolayer of carbon on the plane may be enough to grow graphene. In this study, graphene was grown directly on various substrates such as transparent substrates, insulators, and semiconductors without any catalyst. The directly grown graphene is triangular nanographene with sides of 100 200 nm in length. This study suggests that graphene can be directly grown on all substrates.

’ INTRODUCTION Chemical vapor deposition (CVD) methods have been utilized for thin film development because it is advantageous for the development of large-scale devices. Graphene also has been grown by CVD methods; typically metallic catalysts such as transition and noble metals are employed.1,2 However, in growing graphene on various metal films on substrates, the sandwiched metallic layer becomes a serious obstacle for diverse device developments because the metallic catalyst should be removed before utilizing the unique physical properties of graphene for device applications, or the grown graphene should be transferred to desirable substrates, such as semiconducting silicon, insulating silicon dioxide, or transparent substrates for the fabrication of electronic and optical devices. Removing the underlying metallic catalyst layer through the narrow gap between the graphene and the substrates requires the time-consuming wet-chemical etching processes. Additionally, it is unrealistic to apply transfer technology to large-scale applications even though large-scale graphene can be grown on the metal substrate because the required time for the wet-chemical etching proportionally increases with the size of the growth pads. Graphene is a one atomic layer hexagonal structure of carbon atoms.3 5 To grow graphene layers, a metal catalyst may not be needed because monolayer graphene films do not require a higher growth rate perpendicular to the surface even though a higher growth rate along the surface may be desirable. Whenever a gas molecule touches the surface of a solid, at least there is a van der Waals attractive interaction, but the attractive interaction strength may vary between molecules over a wide range according to the origin of the interaction between them. As a result, gases adsorb on the surface as long as the sticking coefficient is not zero. This suggests that graphene can be grown on the surface of any material including insulators, r 2011 American Chemical Society

semiconductors, and metals even though a high temperature may be required to enhance surface reactivity as well as diffusivity to increase the lateral growth rate and the extent of order. This is supported by the recently published work.6 8 Even with excellent physical properties of graphene, such as quasi-particle behavior near the Dirac point with zero-effective mass4,5,9 11 and extremely high carrier mobility,5,12 the intrinsic semimetallic properties of graphene with the lack of an intrinsic energy band gap have become a critical hurdle for the development of devices such as electronic switches, sensors, optical devices, and logic gates because these devices generally require semiconducting materials with a substantial energy band gap compared to environmental thermal energy.13 18 Various efforts have been made to open the energy band gap of graphene by controlling the interaction between graphene and its substrates19 or among graphene layers20 and by controlling the size and edge of the graphene, which has been demonstrated by employing chemical methods,21 unzipping carbon nanotubes,22 using atomic force microscopy23 and metal nanoparticles,24 and shaping with costly labor-intensive, time-consuming physiochemical e-beam lithographical methods.25 Additionally, the large energy band gap was obtained by transforming the sp2 graphene network to the sp3 hybridization by hydrogenation or oxidation, but this leads to a large distortion into the nonplanar structure.26 Band gap opening by controlling the size and the edge of graphene at nanoscale is promising and practical because the planner structure of nanographenes may maintain the intrinsically excellent physical property of graphene while the energy band gap can be opened up to a few electronvolts.21,27 Received: February 22, 2011 Revised: June 28, 2011 Published: June 28, 2011 14488

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Figure 1. Optical images of graphene directly grown on quartz substrates at growth temperatures of (a) 800, (b) 900, (c) 1000, and (d) 1100 °C at 10 Torr for 1 h. (e) The sheet resistance vs the transmittance of graphene grown on quartz substrates at various temperatures (800 1100 °C) and pressures (2 100 Torr) for various growth periods (1.5 60 min).

Further, the control of the edge structure allows for the tuning of the physical properties of the nanographene. Thus, the control of the size and edge structure of graphene at nanoscale becomes important in device applications. Among the methods for the size and shape control of graphene, the mostly practical one is by cutting graphene on a substrate using e-beam lithography.25 Before cutting, graphene must be transferred to the desired substrates from highly oriented graphite,25 from graphene grown on a metallic catalyst layer,1 or from graphene grown at high temperatures such as ∼1500 °C on polar crystals such as expensive silicon carbide.28 These transferring processes require multiple, time-consuming, and costly procedures that risk damaging the graphene during the process and are inadequate for the deposition of graphene on various technologically important substrates for the large scale device application of graphene. This study shows catalyst free growth of nanographene on technologically important representative substrates such as silicon (semiconductor), silicon oxide (insulator), and quartz (transparent substrate). The direct synthesis of graphene onto the desired substrates without using any metallic catalyst is free from the hurdle related to the metallic catalysts as well as is able to control the size of the nanographene by self-assembling methods. This will open the opportunity to fully utilize graphene in the development of various types of devices including thin film transistors, light emitting diodes, solar cells, touch pads, displays, etc.

’ EXPERIMENTAL METHODS Graphene was grown directly on technologically important substrates, such as semiconducting 500 μm thick silicon (100) wafers, 300 nm thick insulating silicon dioxide on Si(100) wafers, and 1000 μm thick transparent quartz plates. Before these substrates were loaded into the growth system, they were sequentially cleaned by acetone, isopropyl alcohol, and deionized water. The cleaned substrates were inserted into a 3 in. quartz reaction tube of a thermal chemical vapor deposition (CVD) system. The system was pumped at 1 mTorr, and then the

reaction tube was heated to the desired growth temperatures. Once the pressure reached 1 mTorr at the growth temperature, acetylene (C2H2) (25 sccm) with argon (50 sccm) was introduced. Graphene growth was conducted by varying the temperature (800 1100 °C), pressure (2 100 Torr), and growth period (1.5 60 min.). The sheet resistance was measured by a four-probe setup with a Keithley 2400 sourcemeter, and the transmittance was measured by a Scinco S-4100 UV vis spectrometer. The topography of graphene was investigated by a Bruker AXS N8 NEOS atomic force microscope. The Raman spectra were obtained from the directly grown graphene samples on the various substrates using a JY LabRam HR micro Raman spectrometer with backscattering geometry using a 514.5 nm wavelength Ar-ion laser.

’ RESULTS AND DISCUSSION Without using any metallic catalyst, graphene directly grew not only on quartz but also on silicon (100) and silicon oxide (300 nm)/silicon (100) using acetylene as a carbon source at temperatures of 800 1100 °C. Optical images of the graphene-transparent electrodes, which consist of graphene directly grown on quartz substrates at 800 (a), 900 (b), 1000 (c), and 1000 °C (d) for 1 h at 10 Torr, are shown in Figure 1. With increasing growth temperature, the graphene-transparent electrodes become increasingly darker due to the thicker graphene film grown on the substrates. Graphene was directly grown simultaneously on both sides of the quartz. The experimental data of transmittance (T) vs sheet resistance (Rs) were taken from graphene directly grown on quartz by varying temperature (800 1100 °C), pressure (2 500 Torr), and growth time (1.5 60 min). The transmittance and the sheet resistance decreased with growth temperature, pressure, and time. Figure 1e shows the single side transmittance of graphene-transparent electrodes versus sheet resistance. The single side transmittance of graphene grown on the quartz was extracted from the measured transmittance using a simple relationship of T = T1T2 with assumption of T1 = T2, and plotted in Figure 1e. The sheet resistance at the transmittance of 80% was 14489

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Figure 2. AFM images of nanographene directly grown on silicon substrates for 1 h at 800 °C (a and b) and at 900 °C (c and d), 1000 °C (e and f), and 1100 °C (g and h) using the topography mode except (b) (phase mode). The scan size is 5 μm  5 μm for (a), (c), (e), and (g) and 1 μm  1 μm for (b), (d), (f), and (h).

Figure 3. AFM topographical images of nanographene grown on quartz substrates at 900 °C (a and b) at 50 Torr for 5 min, and at 900 °C (c and d), 1000 °C (e and f), and 1100 °C (g and h) at 10 Torr for 1 h. The insets of panels a and b are line profiles. The scan size is 5 μm  5 μm for (a), (c), (e), and (g) and 1 μm  1 μm for (b), (d), (f), and (h).

∼4 kΩ. The sheet resistance is 1 order of magnitude higher than the large-scale graphene itself but is lower than the samples prepared by spin coating or Langmuir Blodgett monolayer deposited graphene films from graphene dispersed in solutions.29,30 The transmittance data of Figure 1e were fitted with a function of T(%) = (1 R) exp( A/RS)  100 and shown by a solid line where R and A are reflectance and a coefficient related to the resistivity and absorption coefficient of the material,31 and they are ∼0.027 and 0.90 Ω, respectively. The reflectance of ∼0.027 is very close to the quantum optical opacity rather than the reflectance.32

Atomic force microscopy (AFM) images of graphene grown on silicon (100) and quartz are shown in Figures 2 and 3, respectively. The AFM images (Figure 2a f) of graphene grown on silicon at the growth temperature range of 800 1000 °C show nanoscale triangle-shaped planar graphitic carbon structures, triangle nanographenes (TNGs). However, at 1100 °C, spherical nonplanar carbon clusters were grown as shown in panels g and h of Figure 2. The AFM topographical (Figure 2a) and phase images (Figure 2b) of the TNGs grown at the growth temperature of 800 °C, which are shown in panels a and b of Figure 2, 14490

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Figure 4. The Raman spectra taken from nano-graphene grown for 1 h at 800, 900, 1000, and 1100 °C on (a) silicon and (b) quartz substrates, respectively. (c) Raman signal intensity of substrate Si (975 cm 1) and quartz (435 cm 1) as a function of temperature. (d) The ratio of the D band intensity to the G-band intensity grown on silicon (blue) and quartz (red), respectively.

respectively, show that the whole silicon surface was covered with the TNGs, which are right triangles. The sides of TNGs are ∼80 and ∼95 nm, respectively. The phase image (Figure 1b) of the TNGs shows rough edges, which indicate that the edge may be nonplanar. The nonplanar structure of edges can originate from the sp3 hybridization with hydrogen, which likely happened due to the relatively low growth temperature in the hydrocarbon environment. The right triangular shape of nanographene is anisotropic and a good indication that TNGs are crystalline rather than amorphous because the crystal growth largely depends on the crystalline orientation owing to orientation-dependent formation energy. This was supported by their Raman spectra showing the crystalline sp2 characteristics of nano-graphene with a high and sharp G band at 1600 cm 1 and a D band at 1350 cm 1 features as shown in Figure 4a. The sides of the right triangle could be highly symmetric structures such as armchair (A) or zigzag (Z). If one of the sides is armchair, the perpendicular side should be zigzag. Therefore, the possible highly symmetric edge structure of the sides perpendicular to each other is armchair (A) zigzag (Z) while the crystalline structure of the hypotenuse depends on the angle or the length ratio between neighboring edges. When the 30° angle is between a side and the hypotenuse of the triangle, the side is armchair (zigzag) and the hypotenuse is zigzag (armchair). When the angle is 60° between a side and the hypotenuse, the side and the hypotenuse are either armchair or zigzag. However, the measured angles between the side and the hypotenuse are neither 30° nor 60°, but they are close to these. This suggests that the hypotenuse of the right triangle could be the combination of the zigzag and armchair. However, two types of TNGs were observed to have grown on silicon at 900 °C: (right triangles and isosceles triangles) while the majority of TNGs are right triangles as shown in panels c and d of Figure 2. This is evidence that thermal energy plays an important role in determining the shape of the nano-graphenes by competition between the orientation-dependent formation

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energy and the thermal energy. The typical perpendicular sides of TNGs grown at 900 °C are 100 and 135 nm, respectively. The size of the TNGs became larger (80 nm  95 nm) than that at 800 °C, but the population of the TNGs became less compared to that at 800 °C. This indicates that the surface diffusion was enhanced by the relatively high growth temperature. As the growth temperature increased to 1000 °C, TNGs became larger and the perpendicular sides of the TNGs are 160 and 160 nm, respectively, as shown in panels e and f of Figure 2. They are of the isosceles right triangular nano-graphene on silicon, and the number density of the triangles is even lower. This strongly suggests that isosceles triangle formation on silicon is more favorable as the growth temperature increases. Thus, at high growth temperatures, the growth depends less on the diffusion than on the crystal orientation. In the range of 800 1000 °C, the size of TNGs increases with temperature while the number density of nanographene becomes lower with temperature. However, at the growth temperature of 1100 °C, the carbon structure grown on silicon was no longer planar but a spherical cluster with the diameter of ∼200 nm as shown in panels g and h of Figure 2. That carbon cluster growth on silicon was more favorable than the planar structure growth at the high growth temperature can be attributed to the high thermal stress driven by the large thermal expansion coefficient (TEC) with a significant lattice mismatch between graphene and silicon.33,34 It is known that the TEC of silicon is relatively higher than the in-plane TEC of graphene. Unlike TNG growth at 800 1000 °C and spherical cluster growth at 1100 °C on crystalline silicon substrates, the growth behavior of graphene on thick thermal oxide or quartz substrates was distinct in several aspects as shown in Figure 3. First, the AFM studies show that carbon films grown on quartz and silicon oxide on silicon at the low growth temperature of 800 °C were continuous (not shown here). However, the Raman spectrum shown in Figure 4b indicates that the continuous carbon films correspond to nano-graphene rather than a large-scale graphene sheet. This suggests that the continuous graphitic carbon film can be considered as a mosaic of nanographenes rather than a single domain large-scale film. The graphene grown on quartz at 900 °C for 5 min shows TNG islands as well as vacancies as shown in panels a and b of Figure 3. The nanographenes grown on quartz at 900 °C for 5 min are isosceles right triangular nano-graphene in shape and 190 nm  190 nm in size. TNG islands and vacancies resemble each other in shape and size. The typical thickness of the TNG islands and vacancies are 1 2 nm as shown by the line profile in the inset of panels a and b of Figure 3. This strongly suggests that TNG islands were detached from the continuous layer and left TNG vacancies, indicating how the growth mode transition occurred from the continuous-mosaic growth mode to discrete growth mode at 900 °C. The one corner of the isosceles right TNGs became rounded with the longer growth time as shown in panels c and d of Figure 3 (at 900 °C for 1 h). In particular, by increasing the growth temperature to 1000 °C, all three corners of the isosceles right TNGs became rounded as shown in panels e and f of Figure 3. Additionally, a dent was observed in the middle of each rounded triangle nanographene as in panels e and f of Figure 3 (1000 °C). The dent might correspond to a nucleation site. At 1100 °C (quartz), the number of prominent triangles was reduced, but TNGs were merged as shown in panels g and h of 14491

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The Journal of Physical Chemistry C Figure 3 by enhanced diffusion at higher growth temperature. Unlike the spherical formation on silicon substrate, the merging of TNGs is attributed to the relatively low TEC of silicon oxide compared to that of silicon. Panels a and b of Figure 4 show temperature-dependent Raman spectra taken from the nano-graphene grown on silicon and quartz substrate, respectively. Features that originated from substrates (Si features at 520, 980 cm 1 and quartz feature at 480 cm 1) decrease in intensity with growth temperature compared to graphitic carbon features at 1350 (D), 1600 (G), 2710 (2D), 2950 (D + G), and 3200 (2G) cm 1. Over the employed growth temperature region, the position of the G band is located at 1600 cm 1, which is higher than the typical largescale graphene or graphite (1580 cm 1), and this is the typical fingerprint of the nano-graphene.35 As mentioned above, with an increase in the growth temperature, panels a and b of Figure 4 show that the Raman intensity of silicon or quartz features from substrates exponentially decays with temperature because the graphitic carbon film is thicker with higher growth temperature. The exponential decaying intensity variation with temperature is clearly shown in Figure 4c. The silicon signal intensity decays faster than the quartz signal intensity with growth temperature. This indicates that the growth rate of graphitic carbon film on silicon is higher than that on quartz because the substrate signal intensity is inversely proportional to the growth rate for the fixed growth period. It is believed that the higher density of the dangling bond on silicon compared to that on quartz plays an important role in enhancing the growth rate. On silicon, the relatively low growth rate at 900 °C compared to that at 1000 °C suggests that the silicon dangling bond may not be fully activated at 900 °C caused by the dimerization of the dangling bonds. However, on quartz, the quartz signal intensity decays monotonically and exponentially with temperature because the density of the dangling bond increases with temperature, but the dimerization of the silicon dangling bond on quartz did not occur due to spared dangling bonds. The ratio of the intensity of the D band (ID) to the intensity of the G band (IG) in Figure 4d can represent the crystalline property of nano-graphene.28 The intensity ratio for nanographene grown on silicon slightly increases with the growth temperature until 1000 °C. This indicates that the crystalline property of graphene grown on silicon becomes worse with higher growth temperature even though the size of the nanographene becomes larger with temperature as shown in Figure 2. This implies that defect density became larger with growth temperature. It is believed that the higher TEC of silicon played a significant role in the increase of defect density. However, at 1100 °C, the ratio of the intensity of the D band (ID) to the intensity of the G band (IG) was recovered as shown in Figure 4d. This suggests that the thermal stress relaxation by forming spherical graphitic clusters reduces the defect formation unlike the case for the planar structure. However, the crystalline property of graphene grown on quartz is significantly improved with growth temperature. At low temperature, the ratio is relatively high due to the lower activity and dangling bonds. This indicates that the nanographene grown at low temperature has a large number of defects due to the lack of nucleation centers. The density of the active silicon dangling bonds acting as a nucleation center becomes less at lower growth temperature. This can cause large defects. However, as the temperature increases, the number and activity

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of the dangling bond on the quartz becomes higher, and this leads to the growth of a highly crystalline nano-graphene. Overall, the crystalline property grown on silicon is better than that grown on quartz over the employed growth temperature region except at 1000 °C. The Raman study clearly indicates that TNGs correspond to the sp2 two-dimensional graphitic carbon. The temperature-dependent growth behavior of nanographene on silicon was distinct from that on silicon oxide (both polycrystalline quartz and amorphous thick thermal oxide on silicon). These distinct behaviors of graphene growth on two types of substrate can be summarized as follows: discrete (on Si) versus semicontinuous (on SiO2) at 800 °C, decreasing (on Si) versus increasing (on SiO2) the number density of TNGs in the growth temperature range of 800 1000 °C, increasing the size of nanographene (on Si) versus smaller in size and rounded in shape (on SiO2) (800 1000 °C), and clustering (on Si) versus ripening (on SiO2) at 1100 °C. The crystalline properties of TNGs are worse on silicon but better on quartz with growth temperatures. The growth rate increases with temperature except at 900 °C on silicon due to the dimerization of dangling bonds. The clear growth mode transition on silicon oxide was observed at 900 °C by the simultaneous formation of TNG islands and vacancies. The growth behavior of graphene on amorphous silicon oxide was similar to that on polycrystalline quartz. This indicates that the dominant parameter determining the graphene growth behavior is the nucleation density rather than the crystalline property itself.

’ CONCLUSIONS Catalyst-free direct growth of triangular nano-graphenes on semiconductors (silicon), insulators (silicon oxide on silicon), and transparent substrates (quartz) using a simple thermal chemical vapor deposition method was demonstrated. This suggests that nano-graphene can grow on any kind of substrate and the direct growth method can be free from the timeconsuming, costly, labor-intensive multiple processes required for graphene transfer and size control. The shape of the directly grown nano-graphene is an isosceles right triangle, and its size depends on the growth temperature and substrates. This unique growth behavior can be utilized to control the size, shape, and edge of nano-graphene by varying growth temperature and substrates. Thus, it is not difficult to predict that these direct growing methods and graphenes grown on various technologically important substrates can be widely utilized for the development of passive components as well as active components in the development of transparent electrodes, electronic devices, photonics devices, sensors, and spintronic devices. ’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected].

’ ACKNOWLEDGMENT We thank Professor H. K. Park and his research group for allowing us to utilize the AFM facilities. This work was supported by the Korea Research Foundation Grant funded by the Korean Government (KRF-2008-313-C00316) and by the Basic Science Research Program through the National Research Foundation of 14492

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The Journal of Physical Chemistry C Korea (NRF) funded by the Ministry of Education, Science and Technology (MEST) (2010-0005706).

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