Impurities produced during the synthesis process of a material pose detrimental impacts upon the intrinsic properties and device performances of the as-obtained product. This effect is especially pronounced in graphene, where surface contamination has long been a critical, unresolved issue, given graphene’s two-dimensionality. Here we report the origins of surface contamination of graphene, which is primarily rooted in chemical vapour deposition production at elevated temperatures, rather than during transfer and storage. In turn, we demonstrate a design of Cu substrate architecture towards the scalable production of super-clean graphene (>99% clean regions). The readily available, super-clean graphene sheets contribute to an enhancement in the optical transparency and thermal conductivity, an exceptionally lower-level of electrical contact resistance and intrinsically hydrophilic nature. This work not only opens up frontiers for graphene growth but also provides exciting opportunities for the utilization of as-obtained super-clean graphene films for advanced applications.
Surface contamination has long been a great challenge in the whole society of carbon materials1, and still unresolved in graphene2,3. The surface contamination has been intensively highlighted as a major hurdle in probing intrinsic properties of graphene3,4,5,6, and strongly hinders device performance and applications of graphene, such as surface chemistry7,8,9,10 ultrahigh speed electronics2, and transmission electron microscopy (TEM) support11, where clean surface is highly needed. Among the various methods for graphene synthesis, chemical vapour deposition (CVD) approach, especially on Cu substrate, holds great potentials in the scalable and cost-efficient production in a controllable fashion12. Despite recent advances in dictating the grain size and scalability13,14 in CVD approaches, the growth of clean graphene films by eliminating surface contamination remains a daunting challenge2. Although much attention has been paid to post-growth processing aiming for higher degree of cleanness, with an emphasis on the suppression of transfer-related impurities15,16,17,18 and airborne contaminants5, the variations in reported cleanness indicate that the dominant processes and intrinsic root of contamination still remain to be unravelled2,15.
Herein, we prove that the contamination on graphene surface is primarily introduced during the high-temperature CVD growth. The growth of metre-scale, super-clean graphene with advanced performances is facilely realised through the continuous supply of Cu vapour, via an ingenious substrate design using alternating stacks of Cu foil and foam.
Intrinsic contamination on graphene surface during growth
During the high-temperature catalytic growth of graphene on Cu, the graphene surface becomes simultaneously contaminated due to the generation of amorphous carbon, which has been widely reported to be stable in CVD conditions (Fig. 1a)19. The competition between the formation of sp2 crystalline carbon (graphitisation process) and amorphous carbon during a CVD reaction primarily determines the cleanness of the graphene surface, and this phenomenon has been discussed intensely in academic and industrial settings in relation to the preparation of synthetic graphite and diamond19,20. In the course of graphene growth, Cu would catalyse the decomposition of hydrocarbons as well as the graphitisation process. However, the catalytic activity of Cu would be limited gradually upon the increase of graphene coverage, presumably leading to the formation of amorphous carbon19,21.
Figure 1b shows an atomic force microscopy (AFM) view of graphene on the Cu substrate, right after growth. Sites of discontinuous surface contamination with a thickness of ~1 nm are clearly visible. The presence of surface contamination immediately after graphene growth seems quite ubiquitous, which was also observed in samples from other research groups (Supplementary Fig. 1). The high-resolution transmission electron microscopy (HRTEM) image of obtained graphene membrane transferred without the aid of polymer scaffold15,22 clearly manifests the universal distribution of amorphous carbon contaminants (Fig. 1c), leaving the clean graphene areas at only tens of nm2 (Fig. 1c, inset), in good agreement with previously reported observations2,15. Detailed elemental inspections further reveal that the contaminated regions are enriched with carbon and copper species (Supplementary Fig. 2). In addition, X-ray photoelectron spectroscopy (XPS) analysis identifies that the carbon species within the surface contamination contain sp3-carbon (Supplementary Fig. 3).
Tip-enhanced Raman spectroscopy (TERS) with improved lateral resolution and sensitivity23 was utilised to explore the composition and origin of the surface contamination. The representative TERS spectra of some graphene regions show prominent D (1350 cm−1) band signal (blue curve in Fig. 1d; Supplementary Fig. 4b), while the 2D-band mapping over the same area exhibited a high 2D band density, indicating the intrinsically high quality of the graphene (Supplementary Fig. 4c). The prominent D band in TERS was also observed in samples from other research groups (Supplementary Fig. 5). Therefore, the prominent D band intensity here is originated from the surface contamination rather than graphene itself, and was taken as the signature of amorphous carbon24. Note that, due to considerably higher spatial resolution and selective enhancement of D band intensity than that of conventional Raman techniques23, the observation of visible D band would be only possible in TERS (previous reports regarding TERS of graphene, see Supplementary Table 2), in contrast with in-situ far-filed result (dark cyan) and results in conventional Raman spectroscopy (Supplementary Fig. 6 and Supplementary Table 1). Furthermore, the TERS mapping using the D band (Fig. 1d, inset) shows a similar distribution to that of contamination obtained by AFM. In this way, through TERS, we can differentiate between the clean graphene regions (with nearly no D band, red curve in Fig. 1d) and contaminated graphene regions (high D band intensity, blue curve in Fig.1d).
To clarify the origin of the surface contamination, a 12C/13C isotope-labelling technique21 was employed (Fig. 1e). The evident difference in the contamination-related peaks between the 12C and 13C samples indicates that amorphous carbon has already been isotopically labelled by the carbon feedstock employed in the CVD procedure. In addition, time-of-flight ion mass spectrometry (ToF-SIMS) analysis verifies that the carbon source fuels the dual formation of graphene and amorphous carbon during CVD (Supplementary Fig. 7). Altogether, our results unambiguously reveal that surface contamination is mainly introduced during the high-temperature synthetic process.
The direct and large-area visualisation of amorphous carbon on graphene/Cu foil is the key to the exploration of contamination mechanism towards the growth of clean graphene. This can be readily achieved through the exposure of as-grown graphene/Cu to TiCl4 vapour in humid air, where TiO2 nanoparticles were in-situ deposited on the graphene/Cu25. Interestingly, these TiO2 nanoparticles would be preferentially and selectively adsorbed onto the contaminated regions of graphene (Supplementary Fig. 8). Thanks to the strong Mie-scattering ability of TiO2 nanoparticles to visible light, the as-grown, large-area graphene surface turns multi-coloured after TiO2 deposition, in sharp contrast to the unexposed graphene/Cu sheets (Fig. 1f), thereby indicating the adsorption of large amounts of TiO2 particles on a predominantly unclean surface (Fig. 1f, inset). This urges us to pay further attention to the advancement of CVD synthetic protocols leading to the production of super-clean graphene films at a large scale.
Growth of metre-scale, super-clean graphene
Based on the aforementioned insights, we devise a substrate architecture via stacks of Cu foil and Cu foam (Fig. 2a) (Cu foil-to-foam distance of ~15 μm, Supplementary Fig. 9) to realise the controlled fabrication of super-clean graphene films. The key aspect is an ingenious catalyst design utilising Cu foam mediation leading to the growth of super-clean graphene without detectable contaminations, as witnessed by AFM (Fig. 2b). The as-obtained graphene was further characterised by HRTEM (Fig. 2c), showing continuous, clean sheets with negligible amorphous carbon regions and easily observable atomic lattices (Fig. 2c, inset). Indeed, our prepared graphene possesses much higher degree of cleanness (accounting for >99% clean regions) as compared with the previously reported data and methods (Supplementary Table 3 and Supplementary Fig. 10)15,26,27,28. This confirms the feasibility and significance of our protocol design in obtaining super-clean graphene, with emphasis on suppression of contamination during growth stage. In addition, the quantitative measurement of cleanness of graphene sample is presented in Supplementary Fig. 11.
Folding the copper foil to form an alternating foil/foam stacked roll leads to the growth of large-scale (0.3 m × 1 m), super-clean graphene (Supplementary Fig. 12). Upon exposure to the vaporised TiO2, the graphene/Cu shows no colour change to the naked eyes (Fig. 2d), suggesting a significantly lower-level adsorption of TiO2 nanoparticles and hence a clean surface of grown graphene as compared with the contaminated films prepared via traditional route (Supplementary Fig. 8b, c). Notably, the design of stacked architecture, rather than carbon precursor and/or surface morphology of the substrate, stays as the key parameter in determining the realisation of super-clean graphene. In this sense, sub-centimetre-sized, single-crystalline graphene with super-clean surface can be achieved by using controllably limited carbon supply (Supplementary Fig. 13)29.
During a CVD reaction, Cu substrate would catalyse the decomposition of carbon feedstock and graphitisation process, whose catalytic ability, in turn, would determine crystalline quality of carbon material, i.e., the generation of crystalline graphene or amorphous carbon. In such case, gradually inhibited catalytic ability of Cu due to the increasing coverage of graphene enables the formation of amorphous carbon onto the graphene sheet during the growth (Supplementary Fig. 14a, b). Produced by the decomposition of carbon precursor, a large quantity of carbon species would desorb from Cu substrate, being abundant in the boundary layer to induce the generation of amorphous carbon to contaminate graphene (Fig. 2e; Supplementary Fig. 14c–f and Supplementary Fig. 15)30,31. Hence, the catalytic activity of Cu is relatively insufficient in a conventional CVD reaction. In contrast, due to its high specific surface area, the use of Cu foam mediation is expected to provide sufficient supply of Cu vapour, to unintermittently catalyse the decomposition of carbon species. Then, the formation of amorphous carbon would be suppressed, enabling the growth of crystalline graphene with super-clean nature (Fig. 2e; Supplementary Fig. 14g, i). Note that at low pressure, the Cu foil itself can provide Cu vapour, such as previous reported Cu envelop structure32,33, which can be also used for the remote catalytic growth of graphene on oxide substrates34. However, the amount of Cu vapour is very limited, and especially it would be reduced with the increasing coverage of graphene on Cu foil. Thus, a continuous supply of sufficient Cu vapour is very important for suppressing the formation of amorphous carbon. In this regard, the lower growth rate of graphene on Cu foam than that on Cu foil guarantees a continuous supply of Cu vapour during the entire growth process, which is the key factor in obtaining clean graphene (Supplementary Fig. 16).
Super-clean graphene surface after transfer onto target substrates
Transferring graphene onto target substrates, such as SiO2/Si substrate, is needed for further device fabrication and applications. However, after transfer, the common CVD-grown graphene suffers from an abundant distribution of polymer residues, such as polymethyl-methacrylate (PMMA) residues (Fig. 3a). Interestingly, the availability of super-clean graphene surface on Cu foil ensures a clear reduction of polymer impurities after transfer, as confirmed in the AFM image, where polymer residue and the amorphous carbon are clearly invisible (Fig. 3b). The surface cleanness is comparable with mechanically exfoliated counterpart with only substrate-induced small fluctuation, evident from the similar height histogram (Fig. 3c). In addition, this small fluctuation becomes invisible after transferring super-clean graphene onto atomically flat mica substrate (Supplementary Fig. 17). For a large-area evaluation of cleanness, graphene was transferred with the assistance of deuterated PMMA (2H-PMMA). Therefore, it was possible to quantify the amount of PMMA residues using mass signature of isotope 2H (Fig. 3d, inset). After the removal of 2H-PMMA by acetone, ToF-SIMS was conducted to detect the amount of the 2H on graphene surface. For unclean graphene sample after transfer, a clear 2H− peak is observed at the m/Z of 2, indicating the abundant PMMA residues. In contrast, 2H− peak is invisible in clean sample, confirming clear reduction of transfer-related impurities on graphene surface after the transfer (Fig. 3d). The large-scale cleanness is additionally confirmed by the statistic of intensities of 2H− peaks of unclean and super-clean graphene surface after transfer (Fig. 3d, inset).
Optical and electrical properties of super-clean graphene
Our super-clean graphene, almost devoid of amorphous carbon (>99.0% clean regions), represents the quality of CVD-derived graphene with optical and electrical properties comparable with that obtained by mechanical exfoliation. The large-area transfer of clean and unclean samples onto transparent polyethylene terephthalate (PET) substrates was achieved by means of a thermal release tape method14, where super-clean graphene sample exhibits a lighter contrast than that of the unclean counterpart (Fig. 4a, inset). The elimination of amorphous carbon and low-level of polymer residues after the transfer altogether contribute to the enhanced light transparency in both monolayer and multilayer graphene films (Fig. 4a; Supplementary Fig. 18).
From the point of view of device applications, the construction of high-performance graphene transistors relies heavily upon the reliable fabrication of electrical contact and graphene channel with high carrier mobility35. In our study, we found that commonly used contact metals (Au, Cr and Pd) were preferentially adsorbed on the contaminated regions upon deposition (Supplementary Fig. 19). This would result in small grain sizes of the deposited metal and less effective contact between the metal and graphene. Furthermore, transfer-related polymer residues would lead to a larger coupling length and reduce the coupling strength between the metal and graphene, giving rise to larger contact resistances36. In contrast, using a graphene transistor array (six transistors for transfer length method (TLM) measurements) (Fig. 4b, inset)37, the extracted contact resistance of super-clean graphene devices is 115 ± 19 Ω μm with minimum value of 96 Ω-μm at room temperature when the channel is slightly p-doped (Fig. 4b; Supplementary Fig. 20), substantially lower than the previously reported values (Supplementary Table 4) and that of unclean graphene with similar work function, confirming the importance of improving cleanness for better contact36,37 (Supplementary Fig. 21).
Furthermore, the measured field-effect transistor mobility of the transferred graphene on SiO2/Si substrate ranges from 14,900 to 18,500 cm2 V-1 s-1 at room temperature and ~31,000 cm2·V-1·s-1 at 1.9 K (Fig. 4c), higher than those of the unclean samples and previously reported values of CVD graphene on SiO2 substrates (Supplementary Fig. 22, Supplementary Tables 5, 6). Besides the suppression of amorphous carbon, the reduction of polymer residue was reported to result in an enhanced carrier mobility4. The carrier mobility of graphene is also sensitive to the transfer process and water doping. Thus, the transfer process is carefully controlled to obtain a convincing comparison (see the Methods section). The exceptional quality of obtained graphene was further corroborated by the evolution of the quantum Hall effect with magnetic field in devices, where the prominent quantum Hall platform is visible at low magnetic field intensities (Fig. 4c, inset).
Encapsulation of graphene would significantly improve the carrier mobility due to the reduced transfer-related doping, substrate scattering and squeezing effect of h-BN to enhance the interface cleanness38,39,40. Following reported encapsulation techniques38,39, in encapsulated as-grown clean graphene samples, the carrier mobility can be further improved to 625,000 cm2 V-1 s-1 for hole side and 1,083,000 cm2 V-1 s-1 for electrons side, along with observation of ballistic transport at 1.9 K (Supplementary Fig. 23). The reduced and uniform full-width at half maximum (FWHM) of 2D in Raman results were also observed, as an indicator of enhanced carrier mobility in encapsulated graphene (Supplementary Fig. 24)41. Consequently, the improved cleanness and optimisation of transfer techniques, such as significant improvement in the fabrication of large and thin h-BN flakes would be both important for future electrical application of graphene38.
Furthermore, super-clean graphene also exhibits an intrinsically hydrophilic property (Supplementary Fig. 25) and higher thermal conductivity (Supplementary Fig. 26) and electrical conductivity (Supplementary Fig. 27) than that of unclean graphene.
In essence, our results elucidate the origins of surface contamination of CVD graphene and outline a possible path towards the large-scale production of super-clean graphene with reliable high-quality equivalent to that of the mechanically exfoliated counterpart by designing a Cu foam mediator. The method presented here not only opens up avenue for advanced applications of graphene at the industrial level but also promotes further studies on the tailored synthesis of other two-dimensional crystals with super-clean nature.
Growth of super-clean graphene films
The super-clean graphene was grown by the low-pressure CVD method. The Cu foil (25 μm-thick, 99.8%, Alfa Aesar) after electrochemical polishing was placed below a Cu foam (840 g1·m−2 in areal density and 1.0 -mm thickness) (Supplementary Fig. 9) and then they were loaded into CVD system equipped with a 2.5-cm-diameter quartz tube. The system was heated to 1030 °C in 1 h with H2 (200 sccm, ~102 Pa), followed by annealing in H2 (200 sccm, ~102 Pa) for 1 h to eliminate the surface oxygen and contamination. Subsequently, CH4 was introduced to initiate the graphene growth for different duration. Note that, the molar ratio of CH4 and H2 was carefully controlled to produce large graphene single crystals. The flow rate of CH4 is ~0.1–1.0 sccm with pressure from 0.3 Pa to 2.8 Pa. After growth, the system was quickly cooled down to room temperature while still under the same flow. The 13C-isotope-labelled methane is purchased from the Sigma-Aldrich company (production number #490229) with 13C atom ratio of 99%. The 2H-PMMA is purchased from Polymer source company (production number #P100226-dsPMMA) with Mn = 820,000 and Mw = 1,500,100.
The domain size of graphene single crystals grown using the vertical stacking structure of Cu foil and Cu foam can be tuned from tens of micrometres to large than three millimetres by carefully controlling the carbon source supply (H2:CH4 molar ratio). The domain size distribution of graphene grown using different H2:CH4 molar ratio is 22.5 ± 1.2 μm for H2:CH4 molar ratio of 200, 125 ± 14 μm for H2:CH4 molar ratio of 300, 350 ± 11 μm for H2:CH4 molar ratio of 400, 750 ± 80 μm for H2:CH4 molar ratio of 500, 1750 ± 152 μm for H2:CH4 molar ratio of 1000, 2500 ± 575 μm for H2:CH4 molar ratio of 2000, respectively.
The graphene was transferred onto quartz or SiO2/Si substrate by the PMMA method. The graphene film was spin-coated with PMMA and baked at 170 oC for 5 min. Then, 1 M Na2S2O8 solution was used to etch Cu foil away. After being washed by deionized water, PMMA/graphene was subsequently placed onto target substrates. Especially, for reducing transfer-related doping, a dry transfer method was used, where a window cut in scotch tape was applied to the backside of Cu before etching Cu substrate. After etching, the resulting suspended graphene/PMMA/membrane across the tape window can be rinsed by isopropanol and then dried overnight. Subsequently, the membrane was adhered onto a heated SiO2/Si substrates at 150 oC. Finally, PMMA was dissolved by acetone. The super-clean graphene on Cu was transferred onto TEM grid without the assistance of PMMA to endure no interference from polymer residue. In detail, by the evaporation of isopropanol drop between graphene on Cu and TEM grid, a TEM grid is directly contacted with graphene as a supporting layer in the subsequent etching and drying process.
TERS measurements were performed on a modified upright TERS system (NTEGRA Spectra, NT-MDT) in scanning tunnelling microscopy (STM) mode. A long working distance objective (×100, NA 0.7) was used for both excitation and collection of the backscattered light from the sample. The exciting wavelength was 632.8 nm and the laser power on the sample was less than 0.5 mW to avoid sample degradation. All the TERS spectra were collected with an acquisition time of 1 s.
Annealing experiments were performed in the chamber of the ToF-SIMS spectrometer (ToF-SIMS V, ION-TOF GmbH, Munster, Germany) before characterisation. The samples were analysed at 25 °C and after annealing at 100 °C for 1 h. ToF-SIMS spectra were acquired at the annealing temperature using a Bi3+ beam operating at 25 keV. The scanning area was 200 μm × 200 μm with an acquisition time of 40 s. Negative ion spectra were collected for each sample. The software used for peak analysis was SurfaceLab 6.0 from ION-ToF.
Raman spectra were obtained with LabRAM HR-800 with 514nm laser and ×100 objective. Optical microscopy images were obtained with an Olympus BX51 microscopy. Optical transmittance spectra were collected by a Perkin-Elmer Lambda 950 UV-vis spectrophotometer.
SEM images were obtained with SEM (Hitachi S-4800, acceleration voltage 5–30 kV). The graphene on TEM grids was characterised by TEM (FEI Tecnai F30, acceleration voltage 300 kV). HRTEM imaging was performed on an aberration-corrected and monochromated G2 cubed Titan 60–300 electron microscope under 80 kV. Energy-dispersive X-ray (EDX) spectra, bright-field (BF) and high-angle annular dark-field (HAADF) images were acquired using scanning TEM (STEM) mode at 60 kV in a double corrected FEI Titan Themis G2 electron microscope.
The graphene samples were transferred onto SiO2/Si substrates with marks for alignments and then subjected to AFM (Veeco dimension 3100) imaging to determine whether they were flat without wrinkle. Next, each graphene sample was etched into a Hall bar geometry or graphene transistor array using a PMMA etching mask (PMMA 950 K A2 @ 4000 rpm) designed by electron-beam lithography (EBL) (Raith 150 2nd) and reactive-ion etching (RIE) with O2 (Trion Technology Minilock III). Finally, after using EBL to design a PMMA mask (PMMA 950 K A4 @ 4000 rpm), 10-nm Pd and 80-nm Au were deposited on the samples using an electron-beam evaporator (Kurte J. Lesker AXXIS) and then a standard metal lift-off technique. Electrical characterisation at room temperature was performed in a vacuum probe station (Lakeshore TTP-4) using a Keithly Semiconductor Characterisation System (Model 4200-SCS). Electrical-transport and magneto-transport measurements at low temperatures were performed using a lock-in amplifier (Stanford Research 830) at 17 Hz with a source current of 10–100 nA.
The element analysis was performed by XPS (Kratos Analytical AXIS-Ultra with monochromatic Al Kα X-ray). AFM characterisation of graphene on Cu was carried out on Bruker dimension icon microscopy using the Scanasystg mode. The water contact angles were measured on a Dataphysics OCA 20 contact-angle system at room temperature.
The authors declare that the data supporting this study are available within the article and its Supplementary Information files. Further information is also available from the corresponding authors upon reasonable request.
Journal peer review information: Nature Communications thanks the anonymous reviewers for their contribution to the peer review of this work.
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We thank Prof. Lutao Weng in Hong Kong University of Science and Technology for his help in TOF-SIMS. This work was financially supported by the National Basic Research Program of China (Nos. 2016YFA0200101, 2013CB932603 and 2014CB932500), the National Natural Science Foundation of China (Nos. 51432002, 51520105003, 21525310 and 51362029), and Beijing Municipal Science and Technology Commission (Nos. Z161100002116002 and Z161100002116021), and Suzhou Key Laboratory for Advanced Carbon Materials and Wearable Energy Technologies, Suzhou 215006, China.