9.2.1 Direct synthesis of large-scale graphene films

Ni films were e-beam evaporated on SiO₂/Si substrates, and thermally annealed. After exposing the Ni surface to a CH₄, H₂ and Ar mixture in atmospheric conditions, the system was rapidly cooled to room temperature. This fast cooling rate is critical in suppressing formation of multiple layers and for separating graphene layers efficiently from the substrate in the later process.³

A scanning electron microscope (SEM) image of graphene films on a thin nickel substrate shows clear contrast between areas with different numbers of graphene layers (Fig. 9.1a). Transmission electron microscope (TEM) images (Fig. 9.1b) show that the film consists mostly of less than a few layers of graphene. After transfer of the film to a silicon substrate with a 300 nm thick SiO₂ layer, optical and confocal scanning Raman microscope images were made of the same area (Fig. 9.1c,d).⁴ The brightest area in Fig. 9.2(d) corresponds to monolayers, and the darkest area is composed of more than ten layers of graphene. Bilayer structures appear to predominate in both TEM and Raman images for this particular sample, which was prepared from 7 minutes of growth on a 300 nm thick nickel layer. It is found that the average number of graphene layers, the domain size and the substrate coverage can be controlled by changing the nickel thickness and growth process, thus providing a way of controlling the growth of graphene for different applications.

Atomic force microscope (AFM) images often show the ripple structures caused by the difference between the thermal expansion coefficients of nickel and graphene (Fig. 9.1c).⁵ It is believed that these ripples make the graphene films more stable against mechanical stretching,⁶ making the films more expandable. Multilayer graphene films are preferable in terms of mechanical strength for supporting large-area film structures, whereas thinner graphene films have higher optical transparency. It is found that a ~300 nm thick nickel layer on a silicon wafer is the optimal substrate for the large-scale CVD growth that yields mechanically stable, transparent graphene films to be transferred and stretched after they are formed, and that thinner nickel layers with a shorter growth time yield predominantly mono- and bilayer graphene film for microelectronic device application.

9.2.2 Transfer processes for large-scale graphene films

The transfer of CVD-derived graphene films to non-specific substrates is enabled by FeCl₃ etching of the underlying Ni film. Use of buffered oxide etchant (BOE) or hydrogen fluoride solution removes silicon dioxide layers, so the patterned graphene and the nickel layer float together on the solution surface. After transfer to a substrate, further reaction with BoE or hydrogen fluoride solution completely removes the remaining nickel layers.

A dry-transfer process has been developed for the graphene film using a soft substrate such as polydimethylsiloxane (PDMS) stamp.⁷ Here the PDMS stamp is first attached to the CVD-grown graphene film on the nickel substrate (Fig. 9.2d). The nickel substrate can be etched away using FeCl₃ as described above, leaving the adhered graphene film on the PDMS substrate (Fig. 9.2e). By using the pre-patterned nickel substrate (Fig. 9.2c), we can transfer various sizes and shapes of graphene film to an arbitrary substrate. This dry-transfer process turns out to be very useful in making large-scale graphene electrodes and devices without additional lithography processes (Fig. 9.2f–h). Microscopically, these graphene films often show linear crack patterns with an angle of 60°C or 120°C, indicating a particular crystallographic edge with large crystalline domains. In addition, the Raman spectra measured for graphene films on nickel substrates show a strongly suppressed defect-related D-band peak. This D peak grows only slightly after the transfer process (Fig. 9.1e), indicating overall good quality of the resulting graphene film. Further optimization of the transfer process with substrate control makes possible transfer yields approaching 99%.

9.2.3 Optical and electrical properties of the graphene films

For the macroscopic transport electrode application, the optical and electrical properties of 1 × 1 cm² graphene films were measured by ultraviolet-visible spectrometer and four-probe Van der Pauw methods, respectively (Fig. 9.3a,b). The transmittance is measured using an ultraviolet-visible spectrometer after transferring the floating graphene film to a quartz plate (Fig. 9.3a). In the visible range, the transmittance of the film grown on a 300 nm thick nickel layer for 7 minutes is ~80%, a value similar to those found for previously studied assembled films.^8,9 Because the transmittance of the graphene layer is ~2.3%,¹⁰ this transmittance can be increased to ~93% by further reducing the growth time and nickel thickness, resulting in a thinner graphene film. Ultraviolet/ozone etching is also useful in controlling the transmittance under ambient conditions (Fig. 9.4a, upper inset). Indium electrodes were deposited on each corner of the square (Fig. 9.3a, lower inset) to minimize contact resistance. The minimum sheet resistance is ~280 Ω per square, which is ~30 times smaller than the lowest sheet resistance measured on assembled films.^8,9 The values of sheet resistance increase with the ultraviolet/ozone treatment time, in accordance with the decreasing number of graphene layers (Fig. 9.3a)

For microelectronic application, the mobility of the graphene film is critical. To measure the intrinsic mobility of a single-domain graphene sample, the graphene samples are transferred from a PDMS stamp to a degenerate doped silicon wafer with a 300 nm deep thermally grown oxide layer. Monolayer graphene samples were readily located on the substrate from optical contrast¹⁰ and their identification was subsequently confirmed by Raman spectroscopy.⁴ Electron beam lithography was used to make multi-terminal devices (Fig. 9.3b, lower inset). Notably, the multi-terminal electrical measurements showed that the electron mobility is ~3750 cm² V^− 1 s^− 1 at a carrier density of ~5 × 10¹² cm^− 2 (Fig. 9.3b). For a high magnetic field of 8.8T, the half-integer quantum Hall effect (Fig. 9.3b) is observed, indicating that the quality of CVD-grown graphene is comparable to that of mechanically cleaved graphene.

In addition to the good optical and electrical properties, the graphene film has excellent mechanical properties when used to make flexible and stretchable electrodes (Fig. 9.3c,d)⁷ The foldability of the graphene films, transferred to a polyethylene terephthalate (PET) substrate (thickness, ~100 μm) coated with a thin PDMS layer (thickness, ~200 μm; Fig. 9.3c), is evaluated by measuring resistances with respect to bending radii. The resistances show little variation up to the bending radius 2.3 mm (approximate tensile strain of 6.5%), and are perfectly recovered after unbending. Notably, the original resistance can be restored even for the bending radius of 0.8 mm (approximate tensile strain of 18.7%), exhibiting extreme mechanical stability in comparison with conventional materials used in flexible electronics.¹¹

The resistance of graphene films transferred to pre-strained and unstrained PDMS substrates was measured with respect to uniaxial tensile strain ranging from 0 to 30% (Fig. 9.3d). Similar to the results in the folding experiment, the transferred film on the unstrained substrate recovers its original resistance after stretching by ~6% (Fig. 9.3d, left inset). However, further stretching often results in mechanical failure. Thus, it was tried transferring the film to pre-strained substrates¹² to enhance the electromechanical stabilities by creating ripples similar to those observed in the growth process (Fig 9.1c, inset). The graphene transferred to a longitudinally strained PDMS substrate does not show much enhancement, owing to the transverse strain induced by Poisson’s effect.¹³ To prevent this problem, the PDMS substrate was isotopically stretched by ~12% before transferring the film to it (Fig. 9.3d). Surprisingly, both longitudinal and transverse resistance (R_y and R_x) appear stable up to ~11% stretching and show only one order of magnitude change at ~25% stretching. It is supposed that further uniaxial stretching can change the electronic band structure of graphene, leading to the modulation of sheet resistance. These electrochemical properties thus show graphene films to be not only the strongest¹⁴ but also the most flexible and stretchable conducting transparent materials so far measured.¹⁰

9.3.1 Synthesis of 30-inch graphene films and roll-to-roll transfer of graphene films

Figures 9.4a–c present photographs of the roll-based synthesis and transfer process. An 8-inch-wide tubular quartz reactor (Fig. 9.4a) is used in the CVD system, allowing a monolayer graphene film to be synthesized on a roll of copper foil with dimensions as large as 30 inches in the diagonal direction (Fig. 9.4c). A temperature gradient usually exists that depends on the radial position inside the tubular reactor. In our preliminary work, this sometimes resulted in inhomogeneous growth of the graphene on the copper foils. To solve this problem, a ~7.5-inch quartz tube wrapped with a copper foil was inserted and suspended inside the 8-inch quartz tube. In this way, the radial inhomogeneity in the reaction temperature could be minimized. In the first step of synthesis, the roll of copper foil is inserted into a tubular quartz tube and then heated to 1000°C with flowing 8 standard cubic centimetres per minute) (sccm) H₂ at 90 mtorr. After reaching 1000°C, the sample is annealed for 30 minutes without changing the flow rate or pressure. The copper foils are heat-treated to increase the grain size from a few micrometers to ~100 μm, as we have found that the copper foils with larger grain size yield higher-quality graphene films, as suggested by Li and colleagues.¹⁵ The gas mixture of CH₄ and H₂ is then flowed at 460 mtorr at rates of 24 and 8 sccm for 30 minutes, respectively. Finally, the sample is rapidly cooled to room temperature (~10°C s^− 1) with flowing H₂ under a pressure of 90 mtorr.

After growth, the graphene film grown on copper foil is attached to a thermal release tape by applying soft pressure (~0.2 MPa) between two rollers. After etching the copper foil in a plastic bath filled with copper etchant, the transferred graphene film on the tape is rinsed with deionized water to remove residual etchant, and is ready to be transferred to any kind of flat or curved surface on demand. The graphene film on the thermal release tape is inserted between the rollers together with a target substrate (Fig. 9.4b). By repeating these steps on the same substrate, multilayered graphene films can be prepared that exhibit enhanced electrical and optical properties, as demonstrated by Li and colleagues using wet-transfer methods at the centimeter scale.¹⁶ Figure 9.4(c) shows the 30-inch multilayer graphene film transferred to a roll of 188 mm thick polyethylene terephthalate (PET) substrate. Figure 9.4(d) shows a screen-printing process used to fabricate four-wire touchscreen panels¹⁷ based on graphene/PET transparent conducting films. After printing electrodes and dot spacers, the upper and lower panels are carefully assembled and connected to a controller installed in a laptop computer (Figs 9.4e and f), which shows extraordinary flexibility. The scalability and processability of CVD graphene and the roll-to-roll methods presented here are expected to enable the continuous production of graphene-based electronic devices at large scales.

The graphene films seem to be predominantly composed of monolayers when analysed using Raman spectra (Fig. 9.5a). However, atomic force microscope (AFM) and transmission electron microscope (TEM) images often show bilayer and multilayer islands. As the graphene layers are transferred one after another,¹⁶ the intensities of the G- and 2D-band peaks increase together, but their ratios do not change significantly. This is because the hexagonal lattices of the upper and lower layers are randomly oriented, unlike the graphite, so the original properties of each monolayer remain unchanged, even after stacking into multilayers;^18,19 this is clearly different from the case of multilayer graphene exfoliated from graphite crystals.⁴ The randomly stacked layers behave independently without significant change in the electronic band structures, and the overall conductivity of the graphene films appears to be proportional to the number of stacked layers.¹⁶ The optical transmittance is usually reduced by ~2.2–2.3% for an additional transfer, implying that the average thickness is approximately a monolayer (Fig. 9.5b).¹⁰

The unique electronic band structure of graphene allows modulation of the charge carrier concentrations, depending on an electric field induced by gate bias²⁰ or chemical doping,²¹ resulting in enhancement of sheet resistance. We tried various types of chemical doping methods, and found that nitric acid is very effective for p-doping of graphene films. Figure 9.5(c) shows Raman spectra of the graphene films before and after doping with 63 wt% HNO₃ for 5 minutes. The large peak shift (Δu= 18 cm^–1) indicates that the graphene film is strongly p-doped. The shifted G peak is often split near the randomly stacked bilayer islands, as shown in Fig. 9.5(c). We hypothesize that the lower graphene layer, which is screened by top layers, experiences a reduced doping effect, leading to G-band splitting. In X-ray photoelectron spectra (XPS), the C 1s peaks corresponding to sp² and sp³ hybridized states are shifted to lower energy, similar to the case for p-doped carbon nanotubes.²¹ However, multilayer stacking results in blueshifted C1s peaks. We suppose that weak chemical bonding such as π-π stacking interaction causes descreening of nucleus charge, leading to an overall increase in core electron binding energies. We also find that the work functions of graphene films as estimated by UV photoelectron spectroscopy (UPS) are blueshifted by ~130 meV with increasing doping time (Fig. 9.5d, inset). The multiple stacking also changes the work functions (Fig. 9.5d, inset) which could be very important in controlling the efficiency of photovoltaic²² or light-emitting devices based on graphene transparent electrodes.²³

9.3.2 Optical and electrical properties of graphene films

The electrical properties of graphene films formed using layer-by-layer staking methods were also investigated. Usually, the sheet resistance of graphene film with 97.4% transmittance is as low as ~125 Ω/sq (Fig. 9.6a) when it is transferred by a soluble polymer support such as polymethyl methacrylate (PMMA).^16,24,25 The transferrable size achievable using a wet transfer method is limited to less than a few inches of wafer because of the weak mechanical strength of spin-coated PMMA layers. However, the scale of roll-to-roll dry transfer assisted by a thermal release tape is in principle unlimited. In the process of roll-to-roll dry transfer, the first layer sometimes shows approximately two to three times larger sheet resistance than that of the PMMA-assisted wet transfer method. As the number of layers increases, the resistance drops faster compared to the wet transfer method (Fig. 9.6a). We postulate that the adhesion of the first layer with the substrate is not strong enough for complete separation of the graphene films from the thermal release tape. As a result, there can be mechanical damage on the graphene films, leading to an increase in the overall sheet resistance. Because additional layers are not directly affected by the adhesion with the substrate surface, the sheet resistance of multilayers prepared by the roll- to-roll method does not differ much from that for the wet transfer case. The p-doping with HNO₃ clearly enhances the electrical properties of graphene films, and is more effective in roll-to-roll processes. The sheet resistance of the p-doped four-layer graphene film with ~90% optical transmittance is as low as ~30 Ω/sq, which is superior to common transparent electrodes such as indium tin oxide (ITO) and carbon-nanotube films (Fig. 9.6b).²⁶

Standard electron-beam lithography has been used to fabricate graphene Hall bars on conventional 300 nm SiO₂/Si substrates (Fig. 9.6c). The left inset of Fig. 9.6c shows the four-terminal resistance of such samples as a function of backgate voltage (V_bg) at both room temperature (solid curve) and at a low temperature (T = 6K) and zero magnetic field. We observe the graphene specific gate bias dependence of the resistance with a sharp Dirac peak and an effective Hall mobility of 7350 cm² V^–1 s^–1 at low temperatures. This allows the observation of the quantum Hall effect (QHE)²⁷ at 6K and a magnetic field of B = 9T (Fig. 9.6c, right). The fingerprint of single-layer graphene, the half-integer quantum Hall effect, is observed with plateaux at filling factors of u = 2,6 and 10 at R_xy = 1/2, 1/6 and 1/10 (h²/e²), respectively. Although the sequence of the plateaux remains for both the electron side and the hole side, there is a slight deviation from the fully quantized values on the hole side.

Finally, the electromechanical properties of graphene/PET touchscreen panels were tested (Fig. 9.6d). Unlike an ITO-based touch panel, which easily breaks under ~2–3% strain, the graphene-based panel resists up to 6% strain; this is limited not by the graphene itself, but by the printed silver electrodes (Fig. 9.6d).²⁸

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