This observation suggests that two-dimensional layered structure of graphene and its silicon carbide-free integration with silicon can serve as a prototype in advancing silicon anodes to commercially viable technology.Silicon is receiving discernable attention as an active material for next generation lithium-ion battery anodes because of its unparalleled gravimetric capacity. However, the large volume change of silicon over charge–discharge cycles weakens its competitiveness in the volumetric energy density and cycle life. Here we report direct graphene growth over silicon nanoparticles without silicon carbide formation. The graphene layers anchored onto the silicon surface accommodate the volume expansion of silicon via a sliding process between adjacent graphene layers. When paired with a commercial lithium cobalt oxide cathode, the silicon carbide-free graphene coating allows the full cell to reach volumetric energy densities of 972 and 700 Wh l The volume expansion of silicon is a big problem in lithium-ion batteries with silicon anodes. Here, the authors report direct graphene growth on silicon nanoparticles, which effectively mitigates the problem, leading to excellent electrochemical performance. This unparalleled value has stimulated the battery community to invest considerable research efforts because the high gravimetric capacity enables one to increase the energy densities of lithium-ion batteries (LIBs) significantly, and thus bring future LIB applications, such as electrical vehicles, to a reality123467891011121314151718The theoretical gravimetric capacity of silicon (Si) reaches almost 4,000 mAh gAlthough a variety of Si morphologies and their composites with other conductive materials are currently available, for an immediate solution to actual manufacturing that requires high-standard quality control, an easy and scalable synthesis of active Si components is crucial. From a practical application standpoint, among a vast number of viable candidates, the use of commercial Si nanoparticles (NPs) with a simple and efficient conductive surface coating would be highly desirable. Various conductive materials including amorphous carbon (AC) have been investigated as coating materials for Si NP anodes820a priori in the as-made electrode. Also, the graphene-coated Si NP pellet exhibits 12.8 S cm−1 at a marginal graphene content of 1 wt% through a high percolation network. To this end, we overcome a challenge of growing graphene directly on Si surfaces without Si carbide (SiC) formation by developing a chemical vapour deposition (CVD) process that involves a mild oxidant. SiC formation is fatal in Si anode operations because SiC is an electrical insulator with poor defect characteristics. Moreover, SiC is inactive in reacting with Li ions and consequently hinders Li ion diffusion into the Si phase. With the assistance of the graphene interlayer sliding process and enhanced conductivity, the graphene-coated Si NPs reach a volumetric capacity of 2,500 mAh cm−3 , the highest value among those reported to date for any LIB anodes while exhibiting excellent cycling and rate performance.In an attempt to address the limitation of previous conductive coatings as well as to achieve good cycling performance with a significantly higher volumetric energy density, in this study, we adopt multilayer graphene directly grown on the Si surface as a coating material. The two-dimensional (2D) layered character of graphene provides a unique and efficient operation of Si anodes since multilayered graphene can accommodate Si volume expansion via a sliding process between adjacent layers without the need to provide void space 4) as a carbon precursor mixed with H2 was indeed unable to achieve a graphene growth on the Si NPs (2), a mild oxidant, in the CVD process along with CH4  surfaces using CH4 and H2  was monitored in real time during their volume expansion using ate them . Once coate them . Greaterate them , the diaate them , suggestate them , upon swate them . The obsate them . Also, rate them .K-edge near 55 eV as well as Si L-edge at 100 eV . 5 wt%-Gr–Si, 1 wt%-Gr–Si, 2 wt%-AC–Si and pristine Si exhibited lithiation/delithiation capacities of 2,820/2,340, 2,341/1,629, 3,143/1,160 and 3,632/741 mAh g−1, respectively, when Si and coated graphene are taken into account, leading to initial Coulombic efficiencies (CEs) of 83%, 69%, 37% and 20%. The low reversible capacity and initial CEs of AC–Si are attributed to the defects of the AC layers that trap Li ions.The electrochemical performance of the Gr–Si NPs was examined by preparing coin-type half-cells in which Li metal serves as the counter/reference electrode. For comparison, in addition to Gr–Si with different graphene contents, AC–Si and pristine Si were also investigated. The graphene content was controlled by the growth time while the growth temperature was fixed to 1,000 °C. All samples exhibited characteristic lithiation and delithiation plateaus at 0.1 and 0.4 V versus Li/Li−2. Under this constraint, the capacity fading of both 2 wt%-AC–Si and pristine Si was significant even for early cycles, indicating that a simple AC coating is not operational at all. Between both Gr–Si samples, 5 wt%-Gr–Si exhibited better performance than that of 1 wt%-Gr–Si (85% versus 70% retention after 200 cycles), suggesting that depending on the areal capacity, the graphene content becomes critical. The lower retention of 1 wt%-Gr–Si might be explained by occasional defect formation as observed in our in situ TEM investigation based on the same graphene content  and TEM images after 200 cycles. Even after 200 cycles, the SEM image of 5 wt%-Gr–Si visualized individual NPs with clear particle boundaries (−2 increased by 51.3% (15.0→22.7 μm) after the first cycle, which is substantially smaller than that  of 5 wt%-Super P-pristine Si even with half the areal specific capacity of 1.5 mAh cm−2.The graphene coating makes a drastic difference in the cycling performance . For a f content . For thetent  of the pristine Si particles on the assumption of their ideal packing (details in figure caption). The present values are far higher than that (550 mAh cm−3) of the current commercial graphite-based anodes. To further assess the practical viability of 5 wt%-Gr–Si, full cell performance was examined by pairing with a commercial cathode, LiCoO2. By consideration of the average cell voltage of 3.5 V, the areal capacity of 3.0 mAh cm−2, and a total electrode thickness of 108 μm , the volumetric energy density of the present full cell reaches 972 Wh l−1. This value is remarkable since it is 1.8 times as large as that (550 Wh l−1) of widely used current commercial LIBs calculated based on the same metric( http://www.samsungsdi.com/lithium-ion-battery/overview).The volumetric capacity of Gr–Si was elucidated in comparison with those of the theoretical and commercial cases . The vol−1 at its initial cycle, and this value decreased to 700 Wh l−1 after 200 cycles (−1) of the graphite-based control cell. These volumetric energy densities of the graphene-based full cells are converted to 242.0 and 169.6 Wh kg−1 at the first and 200th cycles, which are ∼22% and 0% times higher than those of the graphite-based control full cell, respectively. In addition, despite the high volumetric energy density due to the dense graphene-coated Si particle packing, safety is unlikely to be an issue because the volume change of the anode is <10% of the entire cell volume.To see the improvement in the volumetric energy density in a more tangible manner, both the 5 wt%-Gr–Si (108 μm thick) and graphite electrodes (168 μm thick) were packed into commercial 18650 cylindrical cells with all other conditions (cathode and separator) fixed. When wound into a roll using a commercial taper tension controller, the roll diameter of the 5 wt%-Gr–Si full cell was 14.2 mm, which is significantly shorter than that (17.4 mm) of the commercial graphite counterpart , providi0 cycles . It is a2 as a mild oxidant. The individual graphene layers are directly anchored to the Si surfaces and thus remain integrated with Si during large volume changes of Si over repeated lithiation–delithiation cycles. The robust growth of graphene layers uniformly around each particle as well as over a large number of particles is attributed to the addition of CO2 in the CVD process that generates catalytic sites on the particle surface. The current approach involving CO2 is expected to hold true for other oxide surfaces that suffer from similar inefficient graphene growth and carbide formation.In conclusion, we have developed a CVD process for the direct growth of graphene layers over Si NPs without SiC formation by utilizing COThe layered structure of graphene allows interlayer sliding upon the volume expansion of Si as well as a highly conductive percolating network, resulting in an unprecedented volumetric energy density of an LIB full cell with decent cycle life. Overall, the unique 2D character of graphene and the atom-level engineering of its interface with Si to avoid unwanted SiC formation will allow Si anode technology to make a meaningful step towards its wide commercialization.−1 while a gas mixture of 50 s.c.c.m. CH4 and 50 s.c.c.m. CO2 was flowed. The temperature was maintained for 10 and 20 min, which results in 1 and 5 wt% of carbon content, respectively. The system was then cooled to room temperature (25 °C) in the presence of the mixed gas flow. For the synthesis of SiC–Si, the system is first heated to 1,000 °C (23 °C min−1) while H2 gas flowed at a rate of 100 s.c.c.m. This process was maintained for 60 min to reduce the native oxide layer over Si NPs. Next, the gas mixture consisting of 50 s.c.c.m. CH4 and 50 s.c.c.m. H2 was introduced for 10 min. AC–Si was synthesized by a known pyrolysis process  using citric acid as a carbon precursorFor direct growth of graphene over the surface of Si NPs (Gr–Si), pristine Si NPs  were introduced into a fixed-bed vertical tube reactor made of quartz  in ethylene carbonate:diethyl carbonate:fluoroethylene carbonate=2:6:2=v:v:v were used as separator and electrolyte, respectively. The mass loading of the active material ranged from 0.3 to 2.5 mg. The electrochemical properties were characterized under the galvanostatic mode by scanning the potential between 0.01 and 1.5 V versus Li/Li+ at 0.5C except the first cycle at 0.05C. The full cell performance of Gr–Si was characterized by paring with a LiCoO2 cathode with a fixed areal capacity of 3.2 mAh cm−2. The N/P ratio, defined by total capacity ratio between anode and cathode, was chosen to be 1.0–1.1.For the preparation of the battery electrode, slurry was prepared by dispersing active material (Gr–Si/AC–Si) and Li-polyacrylic acid (1.1 M) in a weight ratio of 80:20 in deionized water. The bare Si electrode, a control sample, was prepared by the same procedure, but the electrode components were pristine Si NPs, Super P and 1.1 M Li-polyacrylic acid. The binder content was fixed to 20 wt%, while the contents of pristine Si NPs and Super P were varied from 65 to 79 wt% and from 15 to 1 wt%, respectively. The half-cell was then assembled by using lithium foil as both reference and counter electrodes. Polyethylene membrane (Celgard) and 1.1 M lithium hexafluorophosphate (LiPFλ=1.54 Å). The scan rate and interval were 4° min−1 and 0.02°, respectively. The X-ray photoelectron spectroscopy measurements were performed by using Physical Electronics  with an Al Kα source . The carbon content on Si NPs was quantitatively analysed  engaging nickel-filtered Cu Kα radiation . The TEM holder was moved to the TEM instrument inside a sealing box, and the holder was exposed to air for <30 s before being loaded into the TEM chamber. The manipulation of the probe tip was precisely controlled by a piezo-electric motor in order to make physical contacts between Gr–Si NPs and Li metal counter electrode. After secure contact was made, a constant bias from −0.5 to −5 V was applied. The microstructure evolution during lithiation was recorded as a movie clip.The How to cite this article: Son, I.H. et al. Silicon carbide-free graphene growth on silicon for lithium-ion battery with high volumetric energy density. Nat. Commun. 6:7393 doi: 10.1038/ncomms8393 (2015).Supplementary Figures 1-12In-situ TEM movie showing lithiation of graphene coated silicon nanoparticles by contact with a Li/LiO2 counter electrodeIn-situ TEM movie showing continuous swelling of graphene coated silicon nanoparticles during lithiationIn-situ TEM movie showing the Li/LiO2 counter electrode detached from swollen graphene coated silicon nanoparticles


This study reports a scalable and economical method to open a band gap in single layer graphene by deposition of cobalt metal on its surface using physical vapor deposition in high vacuum. At low cobalt thickness, clusters form at impurity sites on the graphene without etching or damaging the graphene. When exposed to oxygen at room temperature, oxygen functional groups form in proportion to the cobalt thickness that modify the graphene band structure. Cobalt/Graphene resulting from this treatment can support a band gap of 0.30 eV, while remaining largely undamaged to preserve its structural and electrical properties. A mechanism of cobalt-mediated band opening is proposed as a two-step process starting with charge transfer from metal to graphene, followed by formation of oxides where cobalt has been deposited. Contributions from the formation of both CoO and oxygen functional groups on graphene affect the electronic structure to open a band gap. This study demonstrates that cobalt-mediated oxidation is a viable method to introduce a band gap into graphene at room temperature that could be applicable in electronics applications. Other studies have examined Co on graphene as either adatoms or layers using experimentalK-edges of Co/graphene/Cu samples to examine the carbon and oxygen bonding environments. For Co-deposited graphene, the C = O and C−O bonds are formed and increase with Co thickness. Notably some native oxide groups  are present in low concentration on as prepared graphene/Cu, possibly due to adsorbed water on the graphene surface after exposure to ambient atmosphere. In the O K XPS there may also be a contribution from cobalt oxides in peaks in the range of 529–533 eV, but they are too small in intensity to be differentiated unambiguously using peak fitting2. At 0.12 nm thickness, the clusters become substantially larger in diameter but are spaced further apart, and do not homogeneously cover the graphene surface. This growth pattern supports nucleation at hydrocarbon or impurity sites, limiting vacancy formation. At 0.25 nm, the Co covers the surface more uniformly, presenting as many smaller clusters spaced closely together. When Co is deposited, charge transfer makes the graphene surface more energetically favorable for uniform wetting, resulting in more homogeneous coverage. This is reflected in the root-mean-square roughness showing the lowest value across all samples for 0.25 nm. Note that some local wrinkling of the graphene is evident in the roughness profile shown in The graphene surface was imaged after Co deposition using atomic force microscopy (AFM), with images of Co clusters shown in G and G’ (2D) bands. Below 0.25 nm thickness of Co, the spectra appear very similar to undeposited graphene/SiO2 while a red shift is observed in both G (6 cm−1) and G’ (3.8 cm−1) bands at 0.25 nm of Co. This small shift of the G-band to lower wavenumber is attributed to charge transfer from the Co to the graphene2324G-band is visible that develops into a shoulder peak at 1531 cm−1 for 0.25 nm Co. The similar G-band broadening has been reported for Cr-deposited on grapheneA(G’)/A(G), is calculated to be ~4 or greater for all cases (27A(D)/A(G) ratio (below 0.1) indicates nearly defect-free graphene. To verify that defects were not present over a larger area of the graphene surface, Raman mapping was performed with 120 measurements in a 10 um2 area [D-band intensity and consistent G-band signal is observed for graphene transferred SiO2, which remains consistent for Co ≤ 0.25 nm deposited on the surface. The Δ-band is also observed to be continuous in intensity for the majority of the measurements in We verified that the graphene was undamaged by solvents used in the transfer process using Raman spectroscopy measurements before and after Co deposition . After tll cases , showingum2 area . A very et al. have predicted that metal adatoms  lower the vacancy formation energy in free-standing graphene without assistance from heat31D-band signal measured on our samples, indicating that Co does not induce vacancy formation and thus does not directly manipulate the graphene band structure.With minimal defects in the graphene, damage induced by Co or Co oxides can be ruled out as a possible cause of band gap opening. This result is notable because Co and other transition metals can produce vacancy-related defects in graphene. The use of metal atoms to etch or cleave graphene in a controlled way is a well-known procedure using many metals including Ni, Cr, Pd, Al, or Ti in conjunction with oxidation or hydrogenation at elevated temperaturesK XAS and C Kα XES which probe the partial density of unoccupied and occupied C 2p states, respectively. Over the sample set the unoccupied states exhibit several π* peaks that change in relative intensity with visible trends. The main C = C π* feature (A) appearing at 285.3 eV representing the graphene conduction band (CB) edge, is observed to shift slightly to higher energy. Close inspection reveals an observable shift of A to higher excitation energy for 0.06–0.12 nm of Co compared to undeposited graphene/SiO2. Literature works describing C K XAS measurements of graphene deposited on metallic films or decorated with metal clusters show a splitting of the 285.0 eV π* feature due to metal-graphene covalent bonding3333A nor suppression of the excitonic peak occurs in our samples, suggesting an absence of metal-graphene covalent bonding. Higher energy peaks are also evident that change in intensity and shape upon Co deposition, with some also present in undeposited graphene. Peaks B (287.4 eV) and C (288.3 eV) are due to the presence of hydroxyl (C-OH) or epoxide (C-O-C), and carbonyl (C =O) or carboxyl (COOH) functional groups, some of which may be remnant from the transfer process35B does not arise from bonding to functional groups, but rather from carbon sites that have been perturbed by nearby functional groupsB) could act to break up the graphene over timeD exhibiting small or negligible intensity until 0.25 nm Co is reached, is attributed to the carbonate (CO3) structure (290.2 eV) which is verified by comparison to CoCO3 (Alfa Aesar 99.95%) reference powder. Carbonates may only develop with greater Co thickness because formation of Co-C bonding is energetically unfavorable until reaching a threshold amount of Co. Peak C has a more obvious trend of intensity increase proportional to Co thickness that is consistent with XPS measurements showing a fractional increase of C = O/C-O bonds upon Co deposition. Uniquely in our samples, oxide formation appears to be mediated by the presence of Co.The graphene electronic band structure and a band gap induction were investigated using synchrotron-based X-ray absorption (XAS) and emission (XES) measurements. K XAS and the corresponding C Kα XES spectrum were extrapolated to the baseline by fitting a linear line to the edge of each spectra to determine the conduction-band-minimum (ECBM) and valence-band-maximum (EVBM), respectivelyCBM and EVBM of undeposited graphene/SiO2 intersect at 284.4 eV, indicating a band gap of 0 eV. After depositing 0.12 nm of Co, a band gap of 0.30 eV ± 0.14 eV opens. To our knowledge this is the first reported case of band gap opening in graphene as a result of metal deposition on its surface. The error of 0.14 eV is the maximum expected one when calculating an absolute band gap in this case. Using this technique maximum errors in the range of 0.1–0.2 eV are typical depending on the energy range of the measurements38To quantitatively determine the band gap energy, the leading edges of C D) is low and that of the other oxide species is comparatively high (peak C). At Co thickness of 0.25 nm, no band gap is observed similar to undeposited graphene. 0.25 nm of Co is approximately a monolayer; at this thickness enough Co is present to revert the Co/graphene composite to a metallic state. It would seem that the amount of Co present, as well as how it bonds to the graphene lattice, has a strong influence on the degree of functionalization and hence the band gap. A band gap of 0.05–0.30 eV for Co concentration below 0.25 nm can be rationalized by considering graphene oxide and the O:C ratio. A gap of this size has been predicted to open at an O:C ratio of 27.8–50% in rGOC and D compared to peak A at 0.12 nm Co/Graphene/SiO2.Note that the band gap opening occurs when the concentration of carbonate  representing the onset of Co unoccupied states, and B* (779.4 eV), a characteristic peak of Co3O4. At first glance, Co/graphene/SiO2 shows sharper multiplet features associated with metal-oxides when compared to a 10 nm metallic reference Co film, and does not present a visible peak B*. To confirm a divalent Co (Co2+) oxidation state, multi-ligand field theory (MLFT) simulations were performed and fit to the 0.25 nm Co result, identifying Co2+ in nearly octahedral (C4h) symmetry, consistent with high spin CoO. Shown enlarged in the inset of A* shifts to lower energy by 0.15 eV due to charge transfer caused by a transition from CoO to CoCO3. This shift is identified by comparison to reference CoCO3 powder; the lineshape closely matches that of 0.25 nm Co and reproduces the shift in A*.The other main component not yet considered is the cobalt itself. To probe the bonding environment of Co sites, Co measured . Two peasp3 bond formation when a high enough oxygen concentration is reached in the form of carbonyl or hydroxyl groups.With the results presented, we can speculate that the mechanism of band opening is related to Co-mediated oxidation of the graphene . For sinK XAS spectral shape between 0.12 nm and 0.25 nm are intensity increases in peaks C and D. This suggests that CoCO3 contributes to closing a band gap that is induced by other factors, and therefore calculations were focused on isolating the effect of CoO.To further examine this system, electronic structure simulations have been used to examine the effect of CoO on the graphene band structure. Calculations were focused on CoO because a band gap is observed only when the concentration of carbonates is low compared to CoO and other oxygen functional groups. This is possibly because formation of a carbonate requires a site at a vacancy, defect or edge, meaning that carbonates initially formed are anchored at defect sites natively present on undeposited graphene. This is supported by the Raman measurements which do not show evidence of graphene etching after metal deposition. From p partial DOS. Similar calculations using GGA plus van der Waals correction with adsorbed CoO exhibit a band gap opening in the total DOS (including the carbon 2p projection) above a Co surface loading of 11% (e.g. in the case of one CoO isolated on a 3 × 3 supercell of the graphene unit cell). At lower CoO surface loading the graphene layer remains metallic. Since GGA in the PBE formulation is widely known to underestimate band gaps in materials, we have also used the Tran-Blaha modified Becke-Johnson (mBJ) potential as a correction, resulting in the band gap of 0.3 eV increasing to 0.7 eVd4s states with graphene 2p states. Electrons in hybridized orbitals are then partially localized by the Co-O bond. This localization effectively makes these electrons unavailable for conduction, resulting in partially filled states near the valence band edge and a band gap opening. Although this system was not developed experimentally, it shows that only CoO may be required to open a band gap in graphene.Simulations of a low-coverage CoO/graphene system have been performed using the all-electron WIEN2k code 41. Notab42sp3 bond formation occurs to open a band gap, but only in the local area where Co is deposited, in low enough concentration to prevent deformation or damage of the graphene. Damage and fragmentation of graphene are a weakness of other fabrication methods that rely on solution-based methods with thermal heating to functionalize graphene in a controllable way. Theoretical calculations also reveal an additional contribution from CoO that can result in a band gap opening up to 0.7 eV due to CoO alone. These results point to a dual effect of minor oxidation of the graphene combined with CoO formation that act together to open a band gap.From the results and discussion above, we have developed a low complexity scalable method to open a band gap in graphene of up to 0.30 eV. We propose a plausible mechanism of band gap opening as a two step-process of minor charge transfer to the graphene from Co, followed by formation oxide groups including CoO in a concentration mediated by the Co. The use of PVD with a very slow deposition rate and low cobalt thickness allowed for metal clusters to grow at impurity sites without etching the graphene. Exposure to oxygen forms hydroxyl and carbonyl groups on the graphene where Co has been deposited, which can lead to sufficient O:C ratio to open a semiconducting band gap. This method is unique in that Co/graphene may be attractive for use in patterned electronic devices by using masking to deposit on a selected area of a graphene substrate. This method is therefore scalable to the maximum size of graphene sheets that can be obtained by conventional growth techniques. Most importantly, the band gap produced is of sufficient size to be of use in semiconducting applications, and is stable at room temperature under exposure to ambient conditions. These results suggest that cobalt-mediated oxidation is a viable method to introduce a band gap into graphene that could be feasible in other research areas requiring controlled oxidation of organic films.2 by a standard chemical procedure  with Co metal powder  as the source material in a tungsten boat. A deposition rate of 0.02 Å/s was achieved using a crystal thickness monitor in situ. The substrates were held at room temperature prior to deposition, and allowed to heat as cobalt was deposited. During deposition the pressure did not go above 10−5 torr. Reported final thicknesses are as recorded from the thickness monitor. Afterwards samples were exposed to oxygen at atmospheric pressure, to form both carbon and cobalt oxides.Graphene was grown on polished Cu foil (100 μm thickness) using atmospheric pressure chemical vapor deposition451s XAS at 310 eV, C Kα XES at 265 eV, and Co L2,3 XAS at 810 eV. Raman Spectroscopy measurements were performed using a Renishaw inVia Raman Microscope with a 514 nm laser source at a power of 0.01 W and 1800 lines/mm grating. XPS Core-level and valence-band spectra measurements were made using PHI XPS Versaprobe 5000 spectrometer  at the Federal Ural University Lab in Yekaterinburg, Russia. The XPS spectra were recorded using Al Kα (1486.6 eV) 100 μm spot mode with X-ray power load on the sample less than 25 W and pressure of 10−7 Pa. The spectra were processed using ULVAC-PHI MultiPak Software 9.3 and the residual background was removed with the Tougard method.XAS in total electron yield mode and XES measurements were performed at the Resonant Inelastic X-ray Scattering beam line at the Canadian Light Source of the University of Saskatchewan and at Beamline 8.0.1 of the Advanced Light Source in Berkeley, CA. All spectra were measured at a 45° angle of incidence. XAS spectra were measured in total electron yield mode and normalized to the incoming photon flux as recorded by Au mesh and intensity normalized to a constant background as follows: C k-space, with a plane wave expansion of up to RK max of 5.0, and an convergence of 0.1 mRy was achieved. Muffin tin sphere sizes were C:1.28, Co:1.49, O: 1.28. The PBE functionalσ, Dτ), symmetry-dependent hopping integrals , charge transfer energy (Δ), the difference between the onsite Coulomb repulsion and the core hole potential (U−Q), and the ligand bandwidth (W) and shape48For the CoO/graphene structure, DFT simulations were performed using the WIEN2k code by placing a single Co on 3 × 3, 4 × 4 and 5 × 5 supercells of graphene . The graphene sheets were isolated from each other by 20 angstroms vacuum. The position of Co was relaxed until all forces were below 1 mRy/a.u. Similar relaxation process was done when an O was added above the Co. A gamma-point centered mesh of 20 points was sampled in How to cite this article: Bazylewski, P. F. et al. Selective Area Band Engineering of Graphene using Cobalt-Mediated Oxidation. Sci. Rep.5, 15380; doi: 10.1038/srep15380 (2015).


Few-layer nitrogen doped graphene was synthesized originating from graphene oxide functionalized by selective oxygenic functional groups  under hydrothermal conditions, respectively. Transmission electron microscopy (TEM) and atomic force microscopy (AFM) observation evidenced few-layer feature of the graphene oxide. X-ray diffraction (XRD) pattern confirmed phase structure of the graphene oxide and reduced graphene oxide. Nitrogen doping content and bonding configuration of the graphene was determined by X-ray photoelectron spectroscopy (XPS), which indicated that different oxygenic functional groups were evidently different in affecting the nitrogen doping process. Compared with other oxygenic groups, carboxyl group played a crucial role in the initial stage of nitrogen doping while hydroxyls exhibited more evident contribution to the doping process in the late stage of the reaction. Formation of graphitic-like nitrogen species was controlled by a synergistic effect of the involved oxygenic groups . The doping mechanism of nitrogen in the graphene was scrutinized. The research in this work may not only contribute to the fundamental understandings of nitrogen doping within graphene but promote the development of producing novel graphene-based devices with designed surface functionalization. Nitrogen doping in graphene has proven to be an effective way in tailoring properties of pristine graphene and extending its applications-6 becaus3 has been an important nitrogen source in doping graphene. For example, N-doped graphene could be obtained by annealing GO in NH3 atmosphere, and temperature affected the N-doped type, degree of reduction, and N-doping[3, urea is less corrosive and hazardous and is facile to handle. By replacing NH3 with urea, low N content N-doped graphene with pyridinic and pyrrolic N species was obtained at low temperature, while high N content N-doped graphene with graphitic N species was synthesized at high temperature, showing analogous doping results to the case using NH3. N-doped graphene also could be synthesized under hydrothermal conditions, and the degree of reduction and N content of graphene was dependent on the amount of urea[Currently, techniques for growing graphene mainly include chemical vapor deposition, arc-discharge, thermal annealing, and solvothermal method-8. NH3 hN-doping,10. Comp of urea-14.To the best of our knowledge, previous reports were mainly focused on the effect of experimental parameter changes  on the nitrogen doping within graphene, and few investigation has been conducted on the influence of oxygenic groups on the graphene surface which may be a non-trivial factor in controlling nitrogen doping process.Graphene oxide has been reported as an intermediate during graphene synthesis by chemical method,16. It hIn this work, we have synthesized various graphene oxides with hydroxyl, carbonyl, and carboxyl groups and reduced the graphene oxide using urea under hydrothermal condition. N-doped graphene with different content and bonding configuration of nitrogen was achieved (N content up to 9.5 at.%). The result indicated that carboxyl groups played the main role in the initial stage of nitrogen doping into graphene while hydroxyls were subsequently involved and promoted the nitrogen doping in the late stage of the reaction. It was found that graphitic-like N species was therein the most difficult N species to be obtained, whose formation could be ascribed to synergistic reactions of different oxygenic groups  and also correlated to the distribution and position of the oxygenic groups.4 , H2NCONH2 , H2SO4 , H3PO4 , HCl , H2O2 , C2H5OH , ClCH2CO2H , NH3 • H2O (30%), home-made secondary deionized water. All chemicals were not treated with further purification.Following chemicals were chosen for the experiments: natural flake graphite , KMnOMorphology, composition, microstructure, and surface chemistry of the samples were characterized by JEOL2010 transmission electron microscope , PHI-5400 X-ray photoelectron spectroscopy (XPS), X-ray diffraction (XRD), and Nicolet-Nexus Fourier transform infrared spectroscopy (FT-IR).2SO4:H3PO4 = 9:1 in volume). Then, 18 g of KMnO4 was slowly added. After the temperature was heated to and kept at 50°C, the solution was stirred for 60 h. Finally, the solution was cooled down to room temperature. Deionized water of 300 ml was slowly added for dilution, and 10 ml of H2O2 and 100 ml of HCl were added to remove impurities. Then, a large amount of distilled water was used to make washing and centrifugation, which was repeated for several times until the pH of the solution was adjusted to 5. The brown precipitate was frozen and dried for preservation.Three grams of natural flake graphite was added into 400 ml of ice-mixed acid . A well-dispersed solution was obtained after stirring for 12 h at room temperature. Centrifugal washing was then performed for several times with deionized water to remove the contaminants. Finally, the sample was dried under vacuum at 60°C.A 180 mg of graphene obtained from a thermal treatment on GO  was dissolved into a 20.4 ml methylbenzene. After an ultrasonic treatment for 45 min, the solution was washed with pentane and filtrated to obtain a black solid. Then, 105 mg of the black solid was dispersed into 35.1 ml of deionized water again by adding 52.5 mg of FeSOProcess for preparation is the same as Ge-OH, except that the precursor used in this system is GO.2COOH. Then, ultrasonic reaction was performed for 1.5 h under controlled temperature (approximately 10°C) for another 1 h. Then, the mixture was washed with deionized water and centrifuged for several times until the pH was adjusted to 5. The precipitate was dried under vacuum at 60°C for preservation.A 0.3 g of GO was dispersed into 150 ml deionized water by adding 36 g of NaOH and 30 g of ClCHTwenty-five milliliters of GO  solution (1 mg/ml) was well mixed with 50 ml urea solution (200 mg/ml) to obtain the precursor solution for preparing N-doped graphene. Then, the precursor solution was reacted under hydrothermal treatment for 5 h at 180°C, and black product was obtained. Centrifugal washing was performed for several times with distilled water and ethanol. N-doped graphene was finally obtained by vacuum drying at 60°C. The samples were signed as GO-U, GO-OH-U, GO = O-U, GO-OOH-U, and GO-avg-U.Figure θ = 26° can be indexed as (002) crystal plane of graphite, indicating well-ordered structure of graphite layers with spacing of about 0.34 nm. The characteristic diffraction peak of (002) crystal plane disappears completely after the graphite is fully oxidized, and a new diffraction peak at approximately 11° emerges which can be ascribed to (001) crystal plane of graphite oxide[X-ray diffraction (XRD) patterns of the graphite, pristine GO, and GO-U are displayed in Figure te oxide. After n-1 can be ascribed to stretching vibrations of O-H, C = O, C = C, C-OH, and C-O bonds[-1 are strengthened; whereas, the characteristic peak of C = O disappears. Compared with those of Ge-OH, the characteristic peak intensity of the C = O bond of GO = O sample  evidently increases and the intensity of the characteristic peak for C-OH decreases. Sample GO-OOH exhibits weak characteristic peaks of -COOH, C-OH, and C-O at 1,706, 1,399, and 1,095 cm-1, respectively. Compared to GO and Ge-OH, a red-shift phenomenon is observed clearly in the peak representing C = C bonds of GO = O and GO-OOH samples . This red-shift may be attributed to the abundance of C = O structures in GO = O and GO-OOH; these highly electronegative groups weaken the charge density of carbon atoms in the sample skeleton. Additional fileFigure -O bonds, respectPrevious literature reported that C1s peak of GO can usually be divided into five sub-peaks: C-C  was prepared. This sample has close amount of hydroxyl, carbonyl, and carboxyl functional groups on the graphene surface. Compared with the C/O value of GO-avg, that of GO-OOH is slightly lower while that of carbonyl graphene (GO = O) is slightly higher, which indicates that oxidation of the former increases while that of the latter decreases. By conducting a doping process on the different graphene oxides, graphene with different nitrogen contents and configurations were obtained which can be confirmed by XPS spectra analysis of N-doped graphene converted from GO with different oxygenic groups, as shown in Figure Based on the obtained data, the nitrogen doping and reduction mechanism of graphene oxide to nitrogen-doped graphene with urea under hydrothermal conditions can be proposed:2O and NH3 (1), which acted as a nitrogen source. Under normal pressure and temperatures higher than 300°C, pyridine-like and pyrrolic nitrogen were easy to form on the graphene surface . At the same pressure with temperature over 600°C, graphitic-like nitrogen (N-Q) could be obtained[(a) Decomposition of urea under heating conditions generated Hobtained. Under hobtained.3 to generate amides, including lactam, which was caused by dehydration, as well as imide, and initiated a series of reactions, including tautomeric transformations and removal of small molecules , N-5 and N-6 structures (2 to 4) were formed finally[3 and favored nitrogen doping[(b) Carboxyl groups on the graphene surface provided the main active sites for reaction with nitrogen atoms. This functional group reacted with NH finally. Thus, mn doping. This ma3 are slightly alkaline. Therefore, Ge-OH yielded the lowest nitrogen doping and thought it had a large number of hydroxyl groups. After initial introduction of nitrogen atoms, the original -OH and -OH  generated from C-O-C showed their important contribution in nitrogen doping and subsequent reactions[(c) Hydroxyls were less likely to participate independently in the reactions in this system because a stable p-π conjugated structure was formed between phenolic hydroxyls and the benzene ring. In addieactions,10,20,29eactions.(d) The -OH structure that is adjacent to -COOH preferentially generated N-5 (3), while isolated -COOH and nearby -COOH structures favored the formation of N-6 , which is identical to that for (b).(e) The generation of the N-Q structure includes concerted reactions of various oxygenic groups  (7). N-5 (2) and -OH structures are first obtained in a stepwise manner. Then, dDesigned synthesis of N-doped graphene with different N contents and nitrogen bonding configuration has been successfully obtained under hydrothermal conditions using urea as the nitrogen source. The experimental results indicated that different oxygenic groups  on the graphene oxide surface demonstrated different effects in the nitrogen doping process. The nitrogen doping mechanism depending on the oxygenic groups was discussed. It showed that carboxyl group played the dominant role for graphene oxide in accepting nitrogen atoms by forming N-5 and N-6 doping type, and the maximum nitrogen doping concentration was obtained in the sample. It suggested that hydroxyl did not benefit in the initial nitrogen doping, but the original -OH and C-O-C generated by the -OH played an important role in promoting nitrogen doping and subsequent reaction as long as nitrogen doping was introduced by other groups. The N-Q formation was very complex, requiring synergistic effect of various oxygenic groups. In addition, distribution and position of oxygenic groups were also non-trivial for the nitrogen doping process. We believe that this work may contribute to the understanding of the nitrogen doping mechanism on graphene and promote the development of producing novel graphene-based devices with designed surface functionalization.TEM: transmission electron microscopy; AFM: atomic force microscopy; XRD: X-ray diffraction; XPS: X-ray photoelectron spectroscopy; FT-IR: Fourier transform infrared spectroscopy; GO: graphene oxide; N-5: pyridine-like nitrogen; N-6: pyrrolic-like nitrogen; N-Q: graphitic-like nitrogen.The authors declare that they have no competing interests.YC took the tasks of data collection and analysis and draft writing. BX conducted the experiment of modifying graphene oxide by the selective oxygenic groups. YR carried out the experiment of nitrogen doping of the graphene samples. MY took the task of fabrication of graphene oxide. TX took the task of XRD data measurement. YZ designed the work, proposed growth mechanism of the graphene, and revised the manuscript. YW supervised the work and finalized the manuscript. All authors read and approved the final manuscript.Figure S2. XPS spectra of (a) GO, (b) GO-OH, (c) GO-OOH, (d) GO = O, (e) GO-avg.Typical AFM image of graphene oxide. Click here for file


Antibodies specific for neutralizing epitopes in either Human papillomavirus (HPV) capsid protein L1 or L2 can mediate protection from viral challenge and thus their accurate and sensitive measurement at high throughput is likely informative for monitoring response to prophylactic vaccination. Here we compare measurement of L1 and L2-specific neutralizing antibodies in human sera using the standard Pseudovirion-Based Neutralization Assay (L1-PBNA) with the newer Furin-Cleaved Pseudovirion-Based Neutralization Assay (FC-PBNA), a modification of the L1-PBNA intended to improve sensitivity towards L2-specific neutralizing antibodies without compromising assay of L1-specific responses. For detection of L1-specific neutralizing antibodies in human sera, the FC- PBNA and L1-PBNA assays showed similar sensitivity and a high level of correlation using WHO standard sera (n = 2), and sera from patients vaccinated with Gardasil (n = 30) or an experimental human papillomavirus type 16 (HPV16) L1 VLP vaccine (n = 70). The detection of L1-specific cross-neutralizing antibodies in these sera using pseudovirions of types phylogenetically-related to those targeted by the L1 virus-like particle (VLP) vaccines was also consistent between the two assays. However, for sera from patients (n = 17) vaccinated with an L2-based immunogen (TA-CIN), the FC-PBNA was more sensitive than the L1-PBNA in detecting L2-specific neutralizing antibodies. Further, the neutralizing antibody titers measured with the FC-PBNA correlated with those determined with the L2-PBNA, another modification of the L1-PBNA that spacio-temporally separates primary and secondary receptor engagement, as well as the protective titers measured using passive transfer studies in the murine genital-challenge model. In sum, the FC-PBNA provided sensitive measurement for both L1 VLP and L2-specific neutralizing antibody in human sera. Vaccination with TA-CIN elicits weak cross-protective antibody in a subset of patients, suggesting the need for an adjuvant. The seminal discovery by zur Hausen that certain oncogenic genotypes of Human papillomaviruses (HPV) typified by HPV16 are the etiologic agents of cervical cancer has led to the commercial development of two preventive vaccines, Gardasil and Cervarix An alternative approach to broaden protection is vaccination with the papillomavirus minor capsid protein L2 which induces broadly cross-neutralizing antibodies and protects against experimental challenge with diverse HPV genotypes in animal models in vitro requires specialized culture conditions, and infection does not have a readily discernible phenotype in animals. Hence, HPV pseudovirus production using codon optimized L1 and L2 genes and the encapsidation of a luciferase marker plasmid to facilitate the detection of infection of 293TT cells or upon vaginal challenge of mice have been used to circumvent these limitationsin vitro HPV pseudovirus-based neutralization assay (L1-PBNA) Production of native HPV et al identified spatio-temporal differences during the early events of HPV infection in vivo versus the in vitro infection of 293TT cells resulting in an abbreviated opportunity for furin cleavage of L2, exposure of its neutralization epitopes on the virus surface and thus for L2-mediated neutralization to occur in the L1-PBNA et al utilized an extracellular matrix, exogenous furin and a different target cell line to ensure L2 neutralizing epitope exposure in their L2-PBNA and better emulate infection in vivo. An alternate strategy to improve sensitivity for L2 antisera is the use of a furin pre-cleaved pseudovirion-based neutralization assay (FC-PBNA) in which antibodies prevent the infection of a furin-deficient LoVoT cell line by HPV pseudovirions pre-treated with furin to expose neutralizing epitopes of L2 Day While the L2-PBNA and FC-PBNA have shown better sensitivity towards L2-specific neutralizing antibodies in animal sera, they have not been validated using human serum.Here, we examine the utility of the FC-PBNA for high throughput measurement of both L1 VLP and L2-specific neutralizing antibodies in human sera, and examine it's correlation to the L1- and L2-PBNAs, and protection from vaginal challenge with HPV pseudovirion upon passive transfer of naïve mice with titrated sera of patients vaccinated with L2.Sera of patients vaccinated with Gardasil (n = 30) as part of their routine clinical care were obtained with written informed consent at the University of Alabama, Birmingham and with the prior permission of the University of Alabama, Birmingham Internal Review Board (X081124003). The studies using human sera were done with the prior permission of the Johns Hopkins University Internal Review Board (NA_00043331). Animal studies were carried out in accordance with the recommendations in the Guide for the Care and Use of Laboratory Animals of the National Institutes of Health and with the prior approval of the Animal Care and Use Committee of Johns Hopkins University (MO13M425). Mice were anesthetized with isoflurane inhalation and euthanized by carbon dioxide asphyxiation.http://home.ccr.cancer.gov/Lco/pseudovirusproduction.htm). For HPV fcPsV, the methodology was employed as described in A firefly luciferase expression plasmid was employed as the reporter for both HPV PsV and fcPsV. HPV PsV was used in both neutralization assays and for mouse challenge studies. The HPV fcPsVs were used for FC-PBNA. Standard PsV were generated in 293TT cells as described in the protocol . 2 µg/ml of puromycin was added to 293TTF cells to maintain furin selection and 200 µg/ml of HygromycinB was used to maintain LoVoT selection.Sera of 70/72 patients vaccinated with HPV16 L1 VLP were obtained from a completed phase I clinical study http://home.ccr.cancer.gov/lco/L2neut.htmBriefly, For the L1-PBNA 293TT cells or for the FC-PBNA LoVoT cells were seeded at 15,000 cells/well in a 96 well plate. 24 hours later, using another 96 well plate, serum test samples were either serially diluted two-fold (L2 sera at starting dilution 1∶50) or three-fold  in DMEM culture media, and mixed with HPV PsV (at a dilution previously determined for each batch and type by dilution to ∼100x background). In general, approximately 2 to 5-fold more fcPsV was required for the FC-PBNA (e.g 1∶5000 HPV16 PsV or 1∶1000 HPV16 fcPsV). Mixtures were incubated at 37°C for two hours before being added to 293TT or LoVoT cells. The plates were incubated at 37°C for 72 h. Following this, cells were lysed with 30 µL of Cell Culture Lysis Reagent  for 15 min at room temperature on a rocking platform. The entire lysates were transferred to a 96-well black plate, and luciferase activity was measured by adding 50 µL of luciferin substrate to each well . Assays included a neutralizing serum and L2 antibody to assess batch/assay variability and negative control to determine background. For the L2-PBNA, the assay was carried out as per described in v/v Tween 20 in PBS) before HRP-sheep anti-human IgG diluted 1∶5000 in 1% BSA was added to each well and plates were incubated for 1 h at 37°C. After 3 further washes, 100 µL of ABTS solution, 2,2′Azinobis [3-ethylbenzothiazoline-6-sulfonic acid]  was added to each well for development, and absorbance at 405 nm read using a plate reader Xmark Plus .Maxisorp microtiter 96-well plates  were coated with purified HPV16 full length L2 protein with a 6His tag at 500 ng in 100 µL PBS/well. The plates were incubated overnight at 4°C and then blocked with PBS/1% BSA for 1 h at 37°C. Human serum samples were diluted 1∶50 in PBS/1% BSA were then added to the plates in triplicate for 1 h at 37°C. Following this, plates underwent 3 washes with washing buffer . An equal volume of 3% CMC  in PBS was added to make a total virus challenge volume of 40 µL per mouse. Mice were anesthetized with isoflurane inhalation to effect prior to administration of virus. Half of the challenge dose (20 µL) was injected into the mouse vaginal vault, followed by insertion of a cytobrush cell collector that was turned both clockwise and counter-clockwise 15 times to induce trauma. After removal of the cytobrush, the remaining half of the inoculum was deposited in the vagina. At 72 h after challenge, the mice were anesthetized again and 20 µL of luciferin (7.8 mg/mL) was deposited to in the vaginal vault. Using a Xenogen IVIS 100 imager, bioluminescence was acquired. Signal intensities were further analyzed using Living Image 2.5 software. A mouse plasma volume of 2 mL based on previous estimates was used to estimate the dilutions for 10 µL, 33 µL and 100 µL injected human sera as 1∶200, 1∶60, and 1∶20 respectively.Balb/c mice 6–8 weeks old were purchased from Jackson Laboratories  and were injected subcutaneously with 3 mg of medroxyprogesterone . Three days later, 100 µL of pre-vaccination patient sera or,100 µL, 33 µL or 10 µL of post-vaccinated patient sera was injected intra-peritonally into groups (n = 5) of mice. The following day, each mouse was challenged with 2 µL of HPV58 PsV in 20 µL , the non-linear model Y = Bottom + (Top-Bottom)/(1+10∧((LogEC50-X)*HillSlope)) was fitted to the logWe first tested the WHO international standards for HPV16 (05/134)We next compared the L1-PBNA and FC-PBNA using sera from patients vaccinated with HPV16 L1 VLPs or placebo. The first test involved assessing HPV16 neutralizing antibody titers in sera obtained one month following three intramuscular injections of n = 70/72 volunteers enrolled in a Phase I clinical trial of an experimental HPV16 L1 VLP vaccine or placebo (n = 11) A subset of patients (n = 12) from this study had been previously tested by Pastrana and colleagues when describing the original L1-PBNA which used the secreted alkaline phosphatase reporter system (SEAP L1-PBNA) instead of firefly luciferase The performance of the L1-PBNA and FC-PBNA in detecting HPV16, HPV18, HPV6 neutralizing antibodies elicited by vaccination was next investigated using sera of patients (n = 30) who received Gardasil as part of their routine clinical care . Once agin vitro neutralization response against these non-vaccine related virus types detectable in either assay  patients who were first treated topically with imiquimod at the lesion site for 8 weeks, and then vaccinated three times at monthly intervals with 125 µg of HPV16 E6E7L2 fusion protein (TA-CIN) with no adjuvant in vivo mouse challenge model developed by Roberts and colleagues in vitro neutralizing serum antibody titers detected by the FC-PBNA or L1-PBNA.As the L2 ELISA cannot determine if the antibody response is protective, we next performed passive transfer studies utilizing the in vivo results  were injected intra-peritoneally with 10 µL, 30 µL, or 100 µL of post-vaccinated serum or 100 µL of pre-vaccinated serum. One day after passive transfer of serum, the mice were challenged with HPV58 pseudovirions. As expected, in vivo protection was observed for all five patient sera that were previously positive in both the in vivo challenge and FC-PBNA for L2-neutralizing titers. Importantly, the in vivo protective titers were very similar in titer value and had overlapping confidence intervals with the in vitro neutralizing titers detected by the FC-PBNA  . The resin vitro assays including the HT-PBNASince neutralizing antibodies are the relevant immune correlate and low titers are sufficient for protection, the development of sensitive, robust and high throughput assays for HPV neutralizing antibodies regardless of immunogen is important for clinical development of second generation HPV vaccines, especially those based on L2. As mentioned earlier, although the L1-PBNA developed by Pastrana and colleagues is very sensitive for detection of vaccine type L1-specific neutralizing antibodies, the detection of L2-specific neutralizing antibodies and L1-specific cross-neutralizing antibodies has been problematic. As a result of these limitations, several In this study, we attempted to validate our previously developed FC-PBNA's using human patient serum. To do so, we compared its performance to the L1-PBNA. We first compared sensitivity in detecting both L1-VLP specific and cross-neutralizing titers. With respect to the former, both assays were similarly sensitive for detecting HPV L1-specific neutralizing antibodies from natural infection  or after VLP vaccination . TogetheWhen we previously tested both assays using a single pooled serum from Cervarix-vaccinated mice, the FC-PBNA was more sensitive than the L1-PBNA in detecting HPV31 and HPV45 L1-cross neutralizing antibodies in vitro neutralization assay is of critical importance to evaluate candidate L2 vaccines. The weak and inconsistent L2-specific responses to 125 µg TA-CIN here, were in line with a previous observation in a different cohort of VIN patients vaccinated with a higher dose of TA-CIN (533 µg) which exhibited weaker responses compared to that of healthy volunteersVIN patients vaccinated with TA-CIN produced weak HPV16 L2-specific serum antibody responses detectable by ELISA . Howeverin vitro titers for neutralization measured by the FC-PBNA , Pearson's r = 0.93) . Thus th


Met85), denoted here as Met85-TDP-35. Met85-TDP-35 expressed ectopically in human neuroblastoma cells exhibited reduced solubility, cytoplasmic distribution, and aggregation. Furthermore, Met85-TDP-35 sequestered full-length TDP-43 from the nucleus to form cytoplasmic aggregates. Expression of Met85-TDP-35 in primary motor neurons resulted in the formation of Met85-TDP-35-positive cytoplasmic aggregates and motor neuron death. A neo-epitope antibody specific for Met85-TDP-35 labeled the 35-kDa lower molecular weight species on immunoblots of urea-soluble extracts from ALS-FTLD disease-affected tissues and co-labeled TDP-43-positive inclusions in ALS spinal cord sections, confirming the physiological relevance of this species. These results show that the 35-kDa low molecular weight species in ALS-FTLD can be generated from an abnormal splicing event and use of a downstream initiation codon and may represent a mechanism by which TDP-43 elicits its pathogenicity.The presence of lower molecular weight species comprising the C-terminal region of TAR DNA-binding protein 43 (TDP-43) is a characteristic of TDP-43 proteinopathy in amyotrophic lateral sclerosis (ALS) and frontotemporal lobar degeneration (FTLD). Here, we have identified a novel splice variant of TDP-43 that is upregulated in ALS and generates a 35-kDa N-terminally truncated species through use of an alternate translation initiation codon (ATGThe online version of this article (doi:10.1007/s00401-015-1412-5) contains supplementary material, which is available to authorized users. Supernatants were saved as the low salt fraction, and pellets were homogenized in a high salt Triton X-100 buffer , followed by centrifugation at 16,100×g for 20 min at 4 °C. Supernatants were saved as the high salt Triton X-100 fraction, and the remaining pellets were solubilized in urea buffer dimethylammonio]-1-propanesulfonate (CHAPS), 30 mM Tris–HCl, pH 8.5) and saved as the urea fraction.SH-SY5Y cells grown on 35 mmg for 30 min at 4 °C to float and remove myelin. For immunoblotting, sodium dodecyl sulfate (SDS) sample buffer  was added to 50 μg of each sample fraction, which was heated at 95 °C for 5 min with the exception of the urea fraction to avoid carbamoylation of proteins. Equivalent amounts of protein samples were loaded onto 10 % SDS-PAGE gels and transferred to polyvinyldiflouride (PVDF) membranes, which were blocked in 5 % skim milk dissolved in Tris-buffered saline (TBS) containing 0.2 % Tween-20 for 1 h at RT, followed by overnight incubation at 4 °C with a polyclonal TDP-43 antibody , a polyclonal N-terminal TDP-43 antibody , or a polyclonal C-terminal TDP-43 antibody . Bound antibody was visualized with an ECL Detection System (PerkinElmer) using HRP-conjugated secondary antibodies.For human tissues, gray matter from ALS and control lumbar spinal cord was dissected and weighed, then homogenized at 200 mg/ml using the same biochemical fractionation procedure as for cultured cells but with one exception: after extraction in high salt Triton X-100 buffer, the resulting pellets were homogenized in a high salt sucrose buffer , followed by centrifugation at 16,100×g for 30 min at 4 °C, and the resulting supernatant was saved as the low salt fraction. Pellet was homogenized at 250 mg/ml (starting tissue weight) in high salt buffer  and centrifuged at 180,000×g for 30 min at 4 °C, and the resulting supernatant was saved as the high salt fraction. Pellet was homogenized in 250 mg/ml (starting tissue weight) myelin flotation buffer  and centrifuged at 180,000×g for 30 min at 4 °C, and the floating myelin removed. The resulting supernatant was saved as the myelin flotation fraction. Pellet was homogenized in 50 mg/ml (starting tissue weight) sarkosyl buffer  and incubated with agitation at 22 °C for 30 min, followed by centrifugation at 180,000×g for 30 min at 22 °C. The resulting supernatant was saved as the sarkosyl fraction. Pellet was sonicated with the Ultrasonic Cell Disrupter (Misonix) in 1 g/ml (starting tissue weight) urea buffer  and saved as the urea fraction. All fractions were resuspended with SDS sample buffer and all samples except for the urea fraction were heated at 95 °C for 5 min prior to SDS-PAGE and immunoblot analyses.To extract insoluble proteins from the prefrontal cortex, six ALS cases with pathologically confirmed FTLD-TDP-43, and six neurological normal controls of a similar age group were included. Tissue was dissected from Brodmann’s area 10, weighed, and homogenized at 250 mg/ml in low salt buffer  with a glass Dounce homogenizer (Sigma-Aldrich). Homogenate was centrifuged at 25,000×85-TDP-35 or caspase-3-cleaved TDP-35 (C3-TDP-35) were generated using synthetic peptides corresponding to the neo-epitopes generated by alternative translation initiation at ATGMet85 (M85DETDASSA) or caspase-3 cleavage at DETD89 (A90SSAVKVKR), respectively, linked to keyhole limpet hemocyanin (ProSci Inc.). Both antibodies were diluted 1:1000 for immunoblotting and 1:500 for immunohistochemistry.Rabbit polyclonal antibodies specifically recognizing MetE. coli BL21-AI cells (Life Technologies). Bacterial culture was lysed by sonication on ice with lysis buffer  containing protease inhibitors (Roche Applied Science). Lysate was centrifuged at 15,000 RPM in a Sorvall SS-34 rotor for 30 min at 4 °C. Prior to Ni–NTA agarose bead column (Qiagen) purification of the lysate supernatant, the column was washed with 5 column volumes of 40 mM HEPES–KOH, pH 7.4, 500 mM KCl, 20 mM MgCl2, 10 % glycerol, 20 mM imidazole, 2 mM β-mercaptoethanol, and protease inhibitors. The YFP-TDP-43 fusion protein was eluted from the column with 40 mM HEPES, pH 7.4, 500 mM KCl, 20 mM MgCl2, 250 mM imidazole, 10 % glycerol, 2 mM β-mercaptoethanol, and protease inhibitors. SDS-PAGE analysis of eluent showed that the fusion protein fraction was 95 % pure.Full-length TDP-43 was expressed as a fusion protein with His-tagged Venus YFP at the N-terminus joined by a TEV-cleavable linker. The construct was cloned into the pET-30 vector (EMD Millipore) and expressed in To generate caspase-3-cleaved TDP-43, 2 μg of recombinant human YFP-tagged TDP-43 was incubated with 2 U of active human recombinant caspase-3 (EMD Millipore) in a reaction mix containing 100 mM NaCl, 50 mM HEPES, 10 mM DTT, 1 mM EDTA, 10 % glycerol, and 0.1 % CHAPS, pH 7.4, at 37 °C for 4 h. The reaction was halted by adding 2× SDS sample buffer.85-TDP-35 antibody). Sections were incubated overnight at RT with polyclonal TDP-43 antibody  or Met85-TDP-35 antibody (1:500) diluted in antibody diluent . Bound antibody was visualized using the ImmPRESS Anti-rabbit Ig Peroxidase Kit (Vector Labs) for polyclonal TDP-43 antibody or with the Super Sensitive Polymer HRP Kit (Biogenex) for Met85-TDP-35 antibody labeling. Color development was achieved using NovaRed (Vector Labs) and sections were counterstained with hematoxylin. Images were captured using a QImaging Micropublisher 3.3 RTV digital camera attached to a Leica DM6000 upright microscope and Openlab Imaging software (PerkinElmer).Formalin-fixed paraffin-embedded lumbar spinal cord tissues from ALS and control cases were cut into 4 μm serial sections and then mounted on glass slides. Tissue sections were rehydrated in sequential washes of xylene, xylene:ethanol (1:1), and decreasing concentrations of ethanol, followed by a final wash in PBS. Blocking of endogenous peroxidase activity was achieved by incubation with 3 % (v/v) hydrogen peroxide, and antigens retrieved by incubation in either 10 mM citrate buffer at pH 6.0 for polyclonal TDP-43 antibody (ProteinTech) or Tris–EDTA buffer  and mouse monoclonal TDP-43 antibody  and the secondary antibodies were goat anti-rabbit Alexa Fluor 488 and goat anti-mouse Alexa Fluor 594, respectively. Images were captured with an ORCA ER Hamamatsu digital camera mounted on a Leica DMI600 inverted microscope using Volocity Imaging software (PerkinElmer).For double immunofluorescence labeling, formalin-fixed paraffin-embedded spinal cords were cut into 6 μm sections and subjected to the same processing as above, except that the primary antibodies were rabbit polyclonal Methttp://www.ncbi.nlm.nih.gov/genbank/) yielded two human (accession numbers: BI464234 and BG435185) and fourteen mouse  EST sequences, which revealed the same 91 bp splicing deletion in exon 2 of TARDBP, encoding TDP-43. This splice variant uses a non-canonical splicing pair (UU:AG), thus skipping base pairs 106–196 within exon 2 (denoted here as TDP-43 r.[106_196del]). To investigate the relevance of this splice variant in ALS, we selected twelve cases (11 sALS and one fALS with C9orf72 expansions), all of which exhibited TDP-43- and ubiquitin-positive inclusions upon postmortem examination, and 4 control cases for transcript analysis. RT-PCR was performed using total RNA isolated from lumbar spinal cords, with a forward primer located in exon 2 and a reverse primer in exon 4 of TARDBP  in ALS  compared to controls   of the TDP-43 splice variant transcript showed that if it was transcribed from the first ATGMet1, the 91 bp deletion in exon 2 would cause a frame shift and generate a premature stop codon (TAG) 190–192 bp downstream from ATGMet1, thereby producing a short protein of 63 aa with a pI/MW of 4.32/7.2 kDa. Theoretically, such mRNA transcripts containing premature stop codons would be eliminated through nonsense-mediated mRNA decay  transcript and Met85-TDP-35 were found in the ALS cases with C9orf72 repeat expansions, as well as sALS cases.Immunohistochemical analysis showed labeling of TDP-43 pathology in motor neurons of ALS spinal cord with Met3 itself , 18. As athology , 19, 20.85-TDP-35 would lead to sequestering of full-length TDP-43 from the nucleus to form cytoplasmic aggregates and cause neuronal death, an idea that has been promulgated previously  can generate Met85-TDP-35, there could be additional mechanisms that lead to use of ATGMet85 as an alternate translation start site and these remain to be explored.Based on our observations in culture, expression of Meteviously . As suchMet85, called Met85-TDP-35. A neo-epitope antibody specifically recognizing Met85-TDP-35 labeled the pathologically associated 35-kDa species on immunoblots of ALS-FTLD tissue extracts, and co-labeled TDP-43 pathology. Presence of Met85-TDP-35 was specific to affected motor neurons and could provide a potential mechanism for TDP-43-mediated neurotoxicity in disease.In conclusion, we have identified a splice variant of TDP-43 that is upregulated in ALS and generates a protein species of 35 kDa through alternate translation initiation at ATGSupplementary material 1 (PDF 3471 kb)Supplementary material 2 (PDF 3690 kb)Below is the link to the electronic supplementary material.


For the evaluation of the graphene as well as the nanodiamond quality, Raman spectroscopy used whereas the plasma properties are measured by optical methods. It is found that the diamond nucleation can be influenced by the C2/Hα-ratio in the plasma, while the graphene quality remains mostly unchanged by this parameter. Moreover it is derived from the experimental data that the direct plasma contact with the copper surface is beneficial for the nucleation of the diamond while the growth and quality of the graphene benefits from a larger distance to the plasma. Therefore, this work presents a basis for a method to tailor the deposition of graphene–diamond hybrid films using a MW PE-CVD process or to suppress the diamond deposition entirely if desired.The simultaneous growth of both nanodiamonds and graphene on copper samples is described for the first time. A PE-CVD process is used to synthesize graphene layers and nanodiamond clusters from a hydrogen/methane gas mixture as it is typically done successfully in thermal CVD processes for graphene synthesis. However, the standard thermal CVD process is not without problems since the deposition of graphene is affected by the evaporation of a notable amount of copper caused by the slow temperature increase typical for thermal CVD resulting in a long process time. In sharp contrast, the synthesis of graphene by PE-CVD can circumvent this problem by substantially shortening the process time at holding out the prospect of a lower substrate temperature. The reduced thermal load and the possibility to industrially scale-up the PE-CVD process makes it a very attractive alternative to the thermal CVD process with respect to the graphene production in the future. Nanodiamonds are synthesized in PE-CVD reactors for a long time because these processes offer a high degree of control over the film’s nanostructure and simultaneously providing a significant high deposition rate. To model the co-deposition process, the three relevant macroscopic parameters  are correlated with three relevant process properties (plasma ball size, substrate temperature and C Graphene as well as nanodiamonds, both allotropes of carbon, have obtained significant interest in the last couple of years from scientific as well as from industrial point of view due to their remarkable properties.2-bonded carbon atoms is the reason for the remarkable mechanical, electronic, thermal, optical and chemical properties of graphene. Optical transparence of up to 97 % and electron mobility above 15,000 cm2/Vs have been reported  are fed into a hot-wall reactor at temperatures of around 1000 °C to initialize chemical reactions. Most of these CVD processes use copper substrates and take advantage of the catalytic influences of the copper on the dissociation of hydrogen  and semiconductor technology.There have been some reports of plasma-based methods to decrease the process temperature including the use of microwave plasma CVD to synthesize graphene on nickel foil  films with grain sizes between 5 and 100 nm are deposited from argon-rich CH2/Hα ratio) on the growth process of these two carbon allotropes is investigated. It is aimed to control the deposition of graphene and nanodiamonds independently on each other. Special focus is on process parameters that enable the control of carbon phases formed on the surface to control the allotropes that are synthesized independently and to understand the processes that lead to the respective structures.In this work the influence of the plasma parameters  to atmospheric pressure (1 bar) and above.The co-deposition of Graphene and nanocrystalline diamond structures was performed with a 2.45 GHz IPLAS CYRANNUS®. This spThe reaction chamber is a cylindrical shaped quartz tube with a diameter of 140 mm and a height of 140 mm. The process gas is fed into the process chamber by a gas shower in the top flange and a molybdenum substrate holder of 100 mm diameter is mounted on the bottom flange  similar to those described elsewhere  was introduced in the chamber and a plasma was ignited at a pressure of 5 mbar with a microwave power of 1–2 kW. Plasma cleaning was executed for 20 min to clean the substrate from all remaining organic contaminations and for chemical reduction of the surface. During that time the pressure was raised. Pressure was varied between 20 and 80 mbar for the different experiments. A suitable substrate temperature for the deposition process was measured with an optical pyrometer to be 950–1030 °C. The synthesis of carbon films starts with the introduction of 1.2–36 sccm methane into the process chamber and was performed for 8 min with the methane-hydrogen plasma at constant pressure and substrate temperature. The total gas flux (i.e. hydrogen + methane) was 120 sccm during all coating processes. After that time the reactor was immediately evacuated to inhibit the etching influence of hydrogen on carbon structures at high temperatures. An overview of the process parameters can be found in Table After placing the copper foil in the center of the reaction chamber the reactor was evacuated to 1 × 102 molecules decompose due to the high temperature on the copper surface. Consecutively chemisorption of methane molecules results in carbon species bonded to the top layered copper atoms, which become successively dehydrogenized and eventually building the graphene’s honeycomb pattern. Due to the catalytic influence of copper and the minor solubility of carbon in copper, this leads to the growth of monolayer graphene. The graphene growth process is self-limiting and therefore basically stops after the growth of a monolayer  was coupled to the process chamber via a glass fiber. The entire system was calibrated with a tungsten strip lamp. The influence of the process parameters on the carbon dimer d3Π → a3Π Swan band emission at 516 nm was investigated and a correlation with the film properties  was made. A linear correlation between the intensity of the CRaman spectroscopy is used for the characterization of graphene because it is a popular non-destructive analytical tool used to characterize the structure of the carbon materials  is the combination of phonons with different momenta around K and Γ and therefore provides a measure for the defect density.The D + D′ band (around 2920 cm−1 (measured with 532 nm) for the 2D band is typically assigned to defect free SLG n. The polymer consists of a chain of carbon atoms with alternating single and double bonds between them, each bonded with one hydrogen atom. In this paper Raman spectroscopy is used to quantitatively identify UNCD.The UNCD crystals on the surface can also be characterized by Raman spectroscopy. The main features found in the spectra of UNCD are the already described D- and G-peak due to spAdditionally SEM measurements  are used in this work to identify the crystalline structure of the deposited diamond crystals as well as to measure the nucleation density and the growth rate of the diamond crystals.2/Hα-ratio in the plasma at constant microwave power 1 kW. In Fig. 2/Hα-ratio is depicted as a function of the pressure in a range from 20 to 80 mbar. In addition to that, the methane concentration in the plasma is alternated from 5 to 30 % in steps of 5 %. The measured values are shown as black spots whereas the grey levelled area is a 3D logistic fit. Both increasing pressure and increasing methane concentration lead to a strictly monotonically increasing C2/Hα-ratio.OES is conducted to outline the influence of the macroscopic parameters pressure and methane concentration on the CAs for the deposition of graphene on copper the substrate temperature has been found to have an important influence on the quality of the graphene, temperature measurements of the substrate are conducted as a function of process pressure. The substrate temperature is determined using an optical pyrometer. The general trend in this diagram is that higher pressure leads to higher substrate temperatures that originate form a denser plasma and thus more high-energetic particles interacting with the substrate. However, the two parts of the diagram shown in Fig. The change in slope originates from different heating mechanisms. The higher process pressure also leads to a reduced plasma ball size. As for the higher pressure above approx. 50 mbar the plasma ball is found to lose contact to the substrate beneath. Although the substrate is still heated by the radiation originating from the plasma, direct heat transfer by energetic particles is strongly reduced. This is because the distance between the plasma and the substrate is in the range of mm and energetic particles lose their energy by collisions in the gas, since the mean free path of molecules at a pressure of several tens of mbar is only in the range of a few µm.Repeating the measurements for different methane concentrations  reveals that there is no significant influence of the methane concentration on the pressure–temperature-dependency.2/Hα-ratio on the quality of graphene properly, both substrate-plasma coupling and substrate temperature need to be kept constant.As the size of the plasma ball in MW PE-CVD depends on both the process pressure and the microwave power, these two macroscopic parameters influence the substrate-plasma coupling and therefore the substrate temperature  to the bottom (C2/Hα = 0.045). All the photos are taken with a 100× magnification. An SEM image of the diamond crystals is shown in Fig. The samples are investigated with respect to the quality of the deposited graphene film. Overview photos reveal that during the process black spots are deposited. It is subsequently shown by Raman spectroscopy and SEM images that these black spots are almost ball-shaped nanodiamond clusters. In Fig. 2/Hα-ratio has no decisive influence on the quality of graphene on copper substrates. On all the samples a constant I2D/IG-ratio can be found. Moreover the ID/IG-ratio is constant for all the investigated samples. The symmetric shape of the 2D-peaks is commonly considered to be a hint for monolayer graphene. However, the high intensity of the D-Peak and the shoulder of the G-Peak are strong evidences for graphene structures with high defect density. Quantitatively the slight broadening of the three characteristic peaks in combination with an ID/IG-ratio of about three is indicating an average distance between point defects of about 2 nm in the measured graphene layers  and the two humps located at about 1150 and 1480 cm−1  requires possibly high C2 ratios.The volume V of the diamond clusters and the nucleation density of theses clusters ρtio Fig. . It is s3-bonded carbon is estimated as the product of the volume of one sphere and their area density divided by the concentration of methane in the gas flux  which occurs in all the experiments, however, can be strongly influenced by the C2/Hα-ratio. This finding suggests that the C2/Hα-ratio is a process parameter that allows it to influence the growth of the different carbon allotropes independently. Moreover, the results confirm that the chemical processes that lead to nucleation and growth of graphene and diamond on copper are significantly different and that specific process parameters can be used to synthesize hybrid structures.Summarizing the results of this paragraph, the quality of the synthesized graphene does not significantly depend on the C2/Hα-ratio of merely 0.02. Although based on the results obtained from the previous experiments, an elevated nucleation density on the surface is expected, the nanodiamond particle growth is supposed to be widely inhibited. The results are presented in Fig. Since the interaction of the substrate surface with the energetic particles from the plasma  is believed to lead to point defects in the synthesized graphene layers another experiment has been conducted with a larger distance between the substrate and the plasma to reduce this influence. Therefore, the microwave power is chosen to be 1 kW at a pressure of 65 mbar leading to a substrate temperature of about 990 °C which is close to the previous process temperatures. The methane concentration was chosen to be as low as 1 % providing a CD/IG-ratio of approximately 4.5, a distance of 7 nm between two point defects can be derived. Therefore the average distance of point defects could be reduced by a factor of 5.3 by the larger distance between plasma and substrate.The results ought to be compared to the sample depicted in red in Fig. −1, D′-peak at 1618 cm−1 and 2D-peak at 2661 cm−1; spectrum taken at 1.96 eV laser excitation energy) there is neither considerable strain applied to the sheet, nor sheet doping can be reported.Detailed peak analysis shows that monolayer graphene has been deposited. As there is no peak shift detected (G-peak at 1586 cmIt has been shown that the contact of the sample to the plasma has a negative influence on the quality of the deposited graphene. In contrast to that the direct plasma contact is found to be advantageous in order to synthesize nanodiamonds. This result is in good agreement with earlier experiments explaining the nucleation of diamond by charged or polar precursors in the plasma Buck .In order to obtain additional information about the different growth mechanisms of diamond and graphene and the role the direct plasma contact has on these mechanisms a sample with a substrate geometry as shown in Fig. A cavity is formed in order to enable diffusive transport to the back side of the sample, while inhibiting a direct contact to the plasma. In contrast, the upper side of the sample has direct plasma contact. Equal temperatures at the front side and at the back side of the sample are obtained in this constellation and thus the only difference between the two sides is the plasma-sample interaction.2/Hα-ratio of 0.3. After the deposition the two sides of the sample are investigated separately (results see Fig. The sample is processed at 50 mbar, 1 kW and a methane concentration of 27 % resulting in a substrate temperature of 960 °C and a CWhilst strong diamond nucleation is observed at the front of the substrate, no nanodiamond can be found at its back side (see overview photos in Fig. The most important outcome, however, is that evidence of the different growth mechanisms of diamond and graphene is shown. Direct plasma contact is inevitable for the growth of diamonds. In contrast to that the growth of graphene is based on the temperature activated catalytic influence of copper which does not require direct plasma contact.2/Hα-ratio in the plasma is a process parameter that allows to influence the growth of the different carbon allotropes independently. While the diamond nucleation can be influenced by the C2/Hα-ratio, the graphene quality remains mostly unchanged by this parameter. Moreover, it is derived from the experimental data that the direct plasma contact with the copper surface is beneficial for the nucleation of the diamond clusters while the growth and quality of the graphene benefit from a larger distance to the plasma. Therefore, this work identifies crucial parameters to establish a method to deposit graphene–diamond hybrid films using a MW PE-CVD process, tailor the composition of these films, or entirely oppress the diamond deposition if desired.In this paper the simultaneous growth of nanodiamond clusters and graphene on copper by PE-CVD process is described for the first time. It is found that the C


A scalable technique of chemical vapor deposition (CVD) growth of ultra-thin graphitic film is proposed. Ultra-thin graphitic films grown by a one-step CVD process on catalytic copper substrate have higher crystallinity than pyrolytic carbon grown on a non-catalytic surface and appear to be more robust than a graphene monolayer. The obtained graphitic material, not thicker than 8 nm, survives during the transfer process from a Cu substrate without a template polymer layer, typically used in the graphene transfer process to protect graphene. This makes the transfer process much more simple and cost-effective. Having electrical and optical properties compatible with what was observed for a few layers of CVD graphene, the proposed ultra-thin graphitic film offers new avenues for implementing 2D materials in real-world devices. Graphene, one monolayer of graphite, is one of the most multifunctional materials. Many intrinsic properties of graphene, including high mechanical strength, flexibility, transparency, and good electrical conductivity, make it appealing for a wide variety of applications . These aSynthesis of graphene by an inexpensive and simple technique, enabling the production thereof up to meter scale, is one of the most important breakthroughs towards the graphene industrial revolution , 7. CataDespite CVD technique is recognized as one of the top candidates for graphene mass production, the drawback of CVD graphene is the transfer stage: in the CVD process, graphene is deposited on a transient metal, such as nickel or copper, and thereafter transferred on a final substrate typically through the spin-coating technique. Although the covalent C–C bonds ensure remarkable mechanical properties to graphene monolayer –14, eachIn this sense, for many applications it could be sufficient to use slightly thicker carbon-based material instead of a graphene monolayer. A multilayered graphene could be a good candidate, but it is noteworthy that synthesizing homogeneous multilayered graphene with a desired number of layers is not a trivial task.Pyrolytic carbon or pyrocarbon (PyC) is a disordered carbon material which has been used as thin films for decades , 16. It In this letter, we demonstrate a one-step technique to grow a polycrystalline, graphitic thin film material on a copper substrate. The resultant sub-10-nm-thick graphitic pyrocarbon (GrPyC) film has been investigated using Raman spectroscopy, optical and SEM microscopy, and dc conductivity measurements and has been compared with graphene.The GrPyC film was grown on a Cu foil (99.8 % pure) using the CVD setup described in details elsewhere . The Cu 2 (300 nm). For optical and electric measurements, transferring was done using a conventional routine procedure as for monolayer graphene with polymer support [3 solution and PMMA/GrPyC was rinsed in pure water for 1 hour twice. Next, the sample was deposited on a dielectric substrate, and the PMMA layer was removed by acetone. In Fig. The ultra-thin graphitic pyrocarbon film was transferred from copper to a silica substrate and on silicon coated with thermal SiO support . At firsThe thickness of the GrPyC film was measured by using a Veeco Instruments, Dektak 150 stylus profiler. The transmittance of the GrPyC film was measured by PerkinElmer lambda-18 over a spectral range of 230–800 nm and an empty silica substrate as a reference. The Raman spectrum was measured by inVia Raman microscope using 514-nm excitation wavelength with low excitation power in order to avoid heat-induced effects. Scanning electron microscopy imaging was done with SEM LEO 1550 Gemini. The sheet resistivity was measured using a four-point probe technique with 3-mm probe spacing.The optical transmittance spectrum in Fig. −1 is very narrow  and resembles to that of measured from graphite rather than that of found from PyC, where the width is typically more than 100 cm−1 [The Raman spectrum of GrPyC film in Fig. 100 cm−1 , 28. Suc−1, one can observe a strong 2D peak. In the amorphous PyC, the 2D peak is not found, but it is usually related to a more crystalline graphitic carbon. However, on the right-hand side, one can observe small shoulder in 2D (at 2780 cm−1). This wide peak is typical for Raman spectra of PyC, which indicates the presence of an amorphous PyC-like material [At 2710 cmmaterial .−1 is related to the disordering in the graphitic lattice. D mode is strong having intensity of about half in comparison to the G peak. This shows that the carbon film not only consists of highly crystalline graphitic structure but also contains more defective sp2 carbon, i.e., some pairs of 5- and 7-membered rings exist in the material to give a contribution to the D band. Moreover, the D´ mode at 1550 cm−1 indicates that the film contains not only disordered sp2 carbon but also small amounts of sp3 hydrogenated carbon [The presence of D peak at 1350 cmd carbon , 28.In Fig. −1 for a monolayer graphene and 1.15 kΩ−1 for GrPyC film. Although, ultra-thin graphitic film is more than 20 times thicker than a single-layer graphene, the sheet resistance of the film is only about half of that of graphene. This can be explained by the amorphous fraction of the GrPyC material, which is supposed to be less conducting in comparison with graphite due to a strong electron scattering.The measured sheet resistance was 2.08 kΩA graphene monolayer is conventionally grown on a copper substrate with low hydrocarbon concentration. Because of the low carbon solubility, and of the limited amount of reactant hydrocarbons, the CVD process typically results only in a monolayer graphene , which cThe mechanism behind this is the following: first, a graphene monolayer grows on the copper substrate, taking benefit from the catalytic effect to achieve a very high crystalline order. Then, deposition continues over this first layer, but without the catalytic effect of copper, which is now masked by the first graphene layer; consequently, more defects are introduced and the crystalline size of the graphenic domains diminishes, as reported in “templated growth” experiments .The quality of a transparent conductive film can be estimated by the electrical to optical conductivity ratio . The calDespite the fact that GrPyC does not have as high a quality factor as graphene does, there is a significant benefit, which gives GrPyC an advantage over graphene. A graphene monolayer without an external polymer layer does not survive during the transfer process. This will hamper the use of graphene because the polymer often leaves some remains on the graphene surface after the transfer. Moreover, in an industrial scale, the extra polymer layer increases the cost of the resultant material.In our experiment, we were able to transfer a 1-cm scale GrPyC layer without spin-coated polymer; see Fig. In summary, the CVD-grown nanometrically thin graphitic film on a copper substrate was observed to consist of graphitic carbon with higher crystallinity than amorphous PyC. The sheet resistance of the GrPyC film was lower than that of graphene, while still preserving a rather high transparency. The obtained sub-10-nm-thick GrPyC film demonstrates outstanding robustness: the film survives during the transfer process from a Cu substrate on a microstructure without a template polymer layer. We believe that a significant reinforcement effect of PyC on graphene multi-layers offers new avenues for implementing graphene and graphene-like materials in real-world devices.