A first-order model for graphene CVD on transition metal catalysts that combines kinetic and thermodynamic considerations is developed and experimentally verified. The chemical vapour deposition (CVD) of graphene on three polycrystalline transition metal catalysts, Co, Ni and Cu, is systematically compared and a first-order growth model is proposed which can serve as a reference to optimize graphene growth on any elemental or alloy catalyst system. Simple thermodynamic considerations of carbon solubility are insufficient to capture even basic growth behaviour on these most commonly used catalyst materials, and it is shown that kinetic aspects such as carbon permeation have to be taken into account. Key CVD process parameters are discussed in this context and the results are anticipated to be highly useful for the design of future strategies for integrated graphene manufacture. The question of which catalysts can be used is not only of fundamental importance but also a key issue for integrated graphene manufacture. Roll-to-roll approaches demand reusability of the catalyst, low cost and ease of graphene removal,8 whereas approaches without graphene transfer where the catalyst is part of the device structure or removed while the graphene stays on the target substrate demand graphene growth on a wide range of different catalyst materials at suitably benign CVD conditions.10 In spite of numerous surface science studies on the atomic structure of graphene layers on different single-crystal metal surfaces,12 the conditions and challenges for scalable graphene CVD are notably distinct and with a lack of basic understanding of graphene formation, catalyst selection criteria are yet to be well established.Chemical vapour deposition (CVD) has emerged as the dominant method to synthesise large single-crystalline domains and continuous films of “electronic-grade” graphene and other two-dimensional (2D) materials.14 to avoid graphene layer formation by carbon precipitation on cooling.18 Uniform growth of single layer graphene (SLG) on polycrystalline catalysts with significantly higher carbon solubility, including Ni19 and Pt,20 can, however, be routinely demonstrated,22 and it has been shown that single and few-layer graphene growth in particular for lower process temperatures occurs predominantly during hydrocarbon exposure at temperature and not during cooling.24 While simple thermodynamic considerations of carbon solubility are insufficient to adequately describe even basic growth behaviour on these most commonly used catalyst materials, we show that kinetic aspects such as carbon permeation have to be taken into account. Our simple model allows us to highlight key CVD parameters including the catalyst thickness alongside the growth temperature, carbon precursor pressure, exposure time and cooling rate. This captures well the behaviour apparent in our graphene CVD calibrations on Co, Ni and Cu and we expect our results to be highly useful for the design of future strategies for integrated graphene manufacture.Here, we systematically compare graphene CVD on three polycrystalline transition metal catalysts, namely Co, Ni and Cu, and propose a first-order growth model that can serve as a reference to optimize growth on any elemental or alloy catalyst system. A rationale widely adopted in current literature is that a catalyst material with low carbon solubility, such as Cu, is necessary for single layer growthvia continued hydrocarbon exposure and dissociation at the catalyst surface at constant temperature, which we refer to as isothermal growth, or it can also be crossed vertically at a given carbon concentration via catalyst cooling and the reduction in carbon solubility, which we refer to as precipitation on cooling. For a basic CVD process, consisting of heating up and pre-treatment of a catalyst , exposure to a hydrocarbon at constant temperature and cooling down in an inert atmosphere, graphene formation typically proceeds via isothermal growth but additional growth may also occur during cooling. Hence it is often argued that for the CVD of SLG, a catalyst with low carbon solubility is essential, and that for high carbon solubility metals, additional layers grow by precipitation upon cooling leading to multilayer formation.18JI, and that related to carbon diffusion into the catalyst, JD, with the difference in fluxes, JG, feeding the growing graphene layer  pressures lead to the direct nucleation of multilayer graphene due to a high supersaturation developing at the catalyst surface. Low precursor  pressures relative to the permeability  of carbon in the catalyst will lead to an effective filling of the catalyst bulk with carbon, which in turn can precipitate depending on cooling rate (see below). The thickness of the catalyst dictates the size of the potential carbon reservoir and hence plays an important role in this kinetic model. A thick catalyst will allow continued carbon diffusion into the bulk, JD, which acts as mediating sink for carbon via the flux balance between JI, JD and JG at the catalyst surface and hence provides robust conditions for isothermal SLG growth. The bulk of thinner catalysts will quickly saturate throughout with carbon, rapidly lowering JD and leading to FLG formation and inhomogeneous growth as JI ≫ JD. Exposure time to the carbon precursor is another important parameter. For SLG CVD conditions, too short exposures lead to isolated graphene domains, whereas too long exposure leads to the isothermal formation of additional layers even after complete SLG coverage has been achieved. In this context, catalytic graphene CVD is often described as inherently being self-limited, since coverage of the catalyst with SLG will lower JI. However, this only refers to conditions of low chemical potential i.e. low precursor pressures where a window of exposure time exists over which SLG can be uniformly stabilised. For typical catalyst metals it can be shown (see below), that at higher precursor pressures and/or longer growth times additional graphene layers nucleate at the interface between the catalyst and the initial graphene layer fed through intrinsic defects (including grain boundaries) in the initial SLG. This presents a possible pathway for the controlled CVD of bi- and tri-layer graphene films. The effect of catalyst cooling will depend on the amount of carbon in the catalyst bulk. For slow cooling rates, the direction of JD will reverse and assuming sufficient carbon in the catalyst bulk additional graphene layers can form. For fast cooling rates, however, JD will be suppressed and the carbon will remain in the catalyst bulk, i.e. a high cooling rate helps prevent the precipitation of additional layers upon cooling.The insets of 25 and Ni (∼0.19 atom% at 600 °C),26 to a common low carbon solubility catalyst, Cu (0.0007–0.0280 atom% at 1000 °C).28 For the latter it should be noted that while the carbon solubility is relatively low , there is little data available on the diffusivity of carbon in Cu,27 meaning that significant carbon permeation into thicker Cu may be possible, and thus the role of the catalyst bulk still needs to be considered. This is also relevant to growth strategies that target graphene growth at the back interface of the catalyst, i.e. growth at the catalyst-substrate interface or growth based on catalyst foil pockets.32The growth model outlined in 2-supported sputter-deposited films of Co, Cu, and Ni as catalysts, and adopt a relatively simple CVD process24 in which these are heated to growth temperature in a H2-containing atmosphere, exposed to a hydrocarbon precursor , and then cooled to room temperature directly following removal of the precursor . JD ≈ 0, JG ≈ JI) and thus the supersaturation necessary to nucleate the new layer is rapidly reached. In contrast, the red line shows graphene growth evolution for unsaturated catalysts. In this case, the catalyst bulk provides a sink into which carbon arriving at the catalyst surface can diffuse, mediating the SLG formation at the surface (JG = JI – JD). The unsaturated catalyst thus provides a broader processing window, under which SLG can be stabilized before the nucleation of the second layer.We investigate here the growth of graphene on commercially available polycrystalline foils and SiO2(300 nm)/Si show contrast indicative of uniform monolayer graphene coverage,34 and the uniformity confirms SLG formation across large areas on each of the catalysts. This is further confirmed by Raman spectra of the graphene grown on the Ni, Co and Cu foils (–1 and I2D/IG ratio >2), with 2D peaks well fitted with single Lorentzians all with fwhm of ∼35 cm–1, and with I2D/IG ratios of 2.8, 3.7 and 3.0 respectively. The ID/IG ratios, which relate to defects within the graphene lattice, are all low with values of Cu (7%), Co (8%) and Ni (10%) and thus indicative of high graphitic quality. We note that the growth conditions under which this uniform SLG is achieved vary greatly for the different catalysts in terms of the temperature, precursor pressure, and exposure time which will be discussed in more detail below.To demonstrate these two scenarios, we grew graphene on polycrystalline films and foils of Co, Cu and Ni following the first-order framework we have developed. We first focus on graphene growth on foils. Cu foils  which sh2(300 nm)/Si, optical micrographs of the graphene grown on Co  or an increase in precursor partial pressure (step 3) leads to the nucleation and growth of additional layers beneath the existing SLG37 resulting in FLG as shown in 38Despite the demonstration here that SLG can be stabilized across different transition metal catalysts with a broad range of carbon solubilities, we note that FLG can also be produced on these same catalysts under different conditions, even on Cu for which a self-limited growth behaviour is often suggested in the literature.–6 mbar for 15 min followed by an increase in pressure to ∼10–5 mbar for 5 min. On varying the exposure temperatures from 400 °C to 800 °C , we observe a general improvement in graphitic quality with increasing exposure temperature. At higher temperature however, the improvement in graphene quality is accompanied by the formation of additional graphene layers with inhomogeneous FLG coverage observed , as well as by diffusion into or out of the catalyst bulk.41 Typical graphene CVD processes adopt a pre-treatment phase to reduce carbon contamination within the catalyst and thus provide a more defined starting point prior to growth. The gas-phase supply/removal of carbon is complex, affected by parameters such as temperature, catalyst activity, precursor/etchant chemistry, pressure and the boundary layer that can be present for growth conditions with higher total pressures.49 Nevertheless, the growth atmosphere is typically adjusted to deliver a net flux of carbon to the catalyst surface, with graphene growth at the catalyst surface (JG) fed by the balance between this gas-phase supply (JI) and diffusion into the catalyst bulk (JD). Although different hydrocarbon precursors are selected here for each catalyst , this is to ensure sufficient catalytic dissociation at the chosen growth temperature whilst avoiding undesired pyrolytic dissociation, and is not observed to significantly alter the growth evolution otherwise. For graphene grown on Cu, CH4 is mixed with Ar and H2 during the exposure period. An inert buffer gas (e.g. Ar) is used in conjunction with CH4 because it allows the dilution of carbon required without the need for stringent pumping . H2 is expected to participate more actively in reactions occurring at the catalyst–gas interface and it has been suggested to play a key role in limiting the formation of copper oxide at the catalyst surface,50 helping to activate surface bound carbon needed for growth51 and promoting the desorption/etching of small active carbon species.51 The detailed role of the CH4/H2 balance in the growth reactions has been discussed in previous reports, on the basis of adsorption modelling52 and thermodynamic analysis.53On the basis of our results and existing literature, we now rationalize the observed growth behaviour and develop a first-order model for graphene CVD on catalyst surfaces that can be applied generally to transition metal catalysts. We focus on isothermal growth, as despite the non-negligible carbon solubility of Ni and Co, our extensive JI = JD, and the local concentration of carbon at the surface increases until it reaches the catalyst's solubility limit and a local supersaturation (Δc) develops. Graphene islands subsequently nucleate and grow with continuing exposure (JG = JI – JD) until they eventually impinge on one another to form a complete graphene layer with grain boundaries and defects where the graphene grains merge.36 Despite the graphene coverage, carbon continues to be supplied to the catalyst surface through leakage pathways such as defects and grain boundaries, causing the carbon concentration at the catalyst surface to again increase until additional graphene layers nucleate , the catalyst thickness (l) and its permeability (P). The relation between these parameters is explained in detail by Weatherup et al.19 and it provides a lower bound for the thickness required to avoid catalyst saturation prior to complete SLG coverage given by:Whether or not a catalyst becomes saturated with carbon throughout its thickness prior to complete SLG coverage depends on the rate at which carbon is delivered to the catalyst  can also account for the different behaviour of Cu in comparison to the Ni and Co, when the catalyst thickness is reduced (12 atoms per m per s and Co at 700 °C ∼2 × 1013 atoms per m per s),54 based on the values of diffusivity available in literature,57 and indeed both catalysts show a similar transition from uniform SLG formation to inhomogeneous FLG formation as the catalyst thickness is reduced. There is little data available on the diffusivity of carbon in Cu,27 and reported values of carbon solubility at 1000 °C vary between 0.0007 atom% and 0.0280 atom%,28 however the lack of the transition from uniform SLG formation to inhomogeneous FLG formation indicates Cu has a significantly lower carbon permeability under these conditions. We therefore suggest that the broad processing window for SLG formation on Cu is related to this low permeability , and is not solely the result of the low carbon solubility of Cu. We note that catalyst alloying provides new opportunities for rational catalyst design by allowing properties such as permeability to be tuned to obtain a desired outcome by matching the catalyst with the growth process.This inequality indicates that a low value of  islands . For thi reduced . Whilst 60 The increase in graphene domain sizes at higher growth temperatures, and improved graphitic quality, is thus attributed to the increase in surface diffusivity with temperature.41 We further note that increased catalytic dissociation and increased probability of defect healing may also contribute to this improvement in quality observed for higher growth temperatures.61 This might suggest that higher growth temperatures are in general preferable, however this must be balanced against the increased likelihood of saturating the catalyst throughout with carbon due to the increase in permeability with temperature, as revealed by the inhomogeneous FLG growth observed at 800 °C on Co  without compromising quality.19 Further reductions in growth temperature to those compatible with back-end CMOS integration (≤450 °C) can be achieved with Au–Ni alloy catalysts, whilst still maintaining reasonable graphene quality.62The improvement in graphene quality observed with increasing growth temperature Fig. S1 can also Fig. S1i. The pot31 as schematically indicated in 63The catalysts considered herein , all present simple bulk phase diagrams for the growth temperatures used, consisting of only graphite and a metal-carbon solid-solution, without the involvement of other bulk intermediate phasesin situ observations of isothermal growth, nor by the extent of FLG formation we observe. Instead, our results indicate a distinction in growth behaviour based on whether a catalyst becomes saturated with carbon throughout its thickness during the growth process. For the conditions at which we obtain high-quality and uniform SLG on Co, Ni, and Cu, we remain in a regime in which the catalyst is not filled with carbon throughout its thickness. The catalyst bulk thus provides a sink into which carbon arriving at the catalyst surface can diffuse, mediating the SLG formation at the surface. For conditions at which the Cu, Ni and Co catalysts become saturated with carbon throughout, the undesired formation of inhomogeneous FLG is instead observed, although the low permeability of Cu leads to a particularly broad window of processing conditions over which SLG can be stabilized.In summary, we have developed a first-order model for graphene growth on transition metal catalysts with which we rationalise our systematic CVD calibrations for high-quality uniform SLG on Co, Ni, and Cu. We thereby identify key CVD process parameters  that must be adjusted to achieve the desired outcome based on consideration of the catalyst properties  and the kinetics of growth. SLG and FLG formation on all of these catalysts occurs predominantly at temperature during precursor exposure  for the process conditions adopted herein, rather than by precipitation on cooling. The simple distinction previously made in literature between catalysts with low carbon solubility, where surface segregation/adsorption of SLG is proposed to dominate, and higher carbon solubility catalysts, where FLG formation by precipitation on cooling is assumed, is therefore not supported either by our 66 and to provide a framework for the rational design of catalysts and processes for achieving graphene with properties tailored to specific applications. We expect this to be important in the design of growth strategies to obtain a desired growth outcome on a specific catalyst material, and/or where constraints are placed on the process conditions that can be used, e.g. direct integration into device structures.10The understanding developed herein provides important insights into the CVD of graphene on transition metal catalysts. Given the general nature of these insights, we expect them to be relevant to a range of different catalyst materials including alloys,2 and commercially available polycrystalline 25 μm Ni(99.99% purity), Cu(99.999% purity) and Co(99.95% purity) foils are studied.Graphene is synthesized by chemical vapour deposition in a custom-built cold-wall reactor for Co and Ni catalysts whilst a commercially available Aixtron BM Pro (4 inch) machine is used for graphene growth on Cu. Polycrystalline sputter-deposited films  on SiO2 (1 mbar) heating at ∼300 °C min–1 to the growth temperature. The chamber is then quickly pumped down and once the base pressure ∼10–6 mbar is reached, samples are exposed to a hydrocarbon precursor [C2H2 with pressures in the range 10–6–10–3 mbar for 5 s to 180 min] and subsequently cooled down to room temperature in vacuum at ∼100 °C min–1.For graphene grown on Co and Ni, samples are annealed for 15 min in H2/Ar (50 sccm/200 sccm) heating at ∼100 °C min–1 to the growth temperature. CH4 diluted 0.1% in Ar is then introduced to the chamber for 180 min promoting growth under a CH4/H2/Ar (12 sccm/26 sccm/250 sccm) atmosphere and finally cooled down to room temperature in Ar.For graphene grown on Cu, samples are annealed for 30 min at 250 mbar in a mixture of Hex situ using scanning electron microscopy . When imaging graphene with secondary electrons (SE), it generally appears darker than the catalyst surface due to the low generation of SE in graphene.67 The greater the number of graphene layers the darker they appear. Electron channelling contrast  can also be seen from the variations in contrast within the graphene regions in SEM images. Optical microscopy and Raman spectroscopy  are performed after graphene is transferred to SiO2(300 nm)/Si substrates and corroborate the SEM results.As grown graphene is characterised 3 or 0.5 M (NH4)2S2O8 aqueous solutions are used as etchants and for graphene grown on Co 10 M HCl acid is used. Graphene synthesized on Ni catalysts is transferred using an electrolysis-based bubbling technique in an aqueous NaOH (1 M) solution.68 The graphene/polymer film is rinsed in DI water and then transferred to SiO2(300 nm)/Si substrate. The polymer support is subsequently removed by immersion in a suitable solvent [acetone for PMMA and ethyl acetate for PS] followed by a bath in IPA and drying with an N2 flow.Graphene transfer is performed by depositing polymer supports [polymethylmethacrylate (PMMA) or polystyrene (PS)] onto the graphene/metal-catalyst sample. Wet etching techniques are employed to remove the metal. For graphene grown on Cu, 0.5 M FeCl


Reststrahlen band in the MIR. This allows us to present IR reflectance data exhibiting a continuous evolution from the regime of intraband to interband charge carrier transitions. A consistent and simultaneous analysis of the contributions from both transitions to the optical response yields precise information on the carrier dynamics and the number of layers. The properties of the graphene layers derived from IR reflection spectroscopy are corroborated by other techniques . Moreover, we also present MIR microscopy mapping, showing that spatially-resolved information can be gathered, giving indications on the sample homogeneity. Our work paves the way for a still scarcely explored field of epitaxial graphene-based THz and MIR optical devices.Graphene has emerged as a promising material for infrared (IR) photodetectors and plasmonics. In this context, wafer scale epitaxial graphene on SiC is of great interest in a variety of applications in optics and nanoelectronics. Here we present IR reflectance spectroscopy of graphene grown epitaxially on the C-face of 6H-SiC over a broad optical range, from terahertz (THz) to mid-infrared (MIR). Contrary to the transmittance, reflectance measurements are not hampered by the transmission window of the substrate, and in particular by the SiC  Versatile is probably the most appropriate word to qualify the properties of graphene, this one-atom-thick crystal made of carbon121435−1) and around 200 meV (1613 cm−1), due to the phonon-related absorption bands in the SiC substrate10121415161767891011As a prerequisite for harnessing the peculiar properties of graphene in such applications, it is mandatory to devise accurate methods to perform IR spectroscopy of graphene on different types of substrates of interest. For epitaxial graphene, i.e. graphene obtained by decomposition of SiC at high temperature878900 meV 16 cm−1, du−1, associated with the longitudinal optical (LO) phonon in SiC, and vertically shifted for clarity. The Raman scattering spectra confirm that graphene was successfully grown in all samples, as shown by the presence of the characteristic graphene bands D, G, and G’ located around ~1362 cm−1, ~1581–1596 cm−1 and 2704–2714 cm−1, respectively202122−1, less pronounced for the thicker MEG samples. In the MEG samples, the G and G’ bands are more intense, indicating a larger number of layers. In all samples (except samples #6 and #7 - see below), the shape of the G’ band is well fitted by a single Lorentzian curve, as shown in 82425−1 and 28–52 cm−1, in good agreement with previous studies202122et al. for a thick graphene stack with 70–90 layersD’ band appears (~1623 cm−1) in the spectrum of sample #6, which indicates a larger density of defectsD band varies among the samples, reflecting different disorder densities in the epitaxial graphene layers. Moreover, within the same sample, the relative intensity of the D, G and G’ bands and the FWHM of the G’ band varies (not shown), an indication of local variations in the number of layers and/or doping level and mechanical constraints. Furthermore, some dark inclusions are present in the MEG samples #6 and #7. Two representative spectra measured at the location of these inclusions are shown in the inset of D peak is more prominent, together with a D’ band appearing as a shoulder around 1623 cm−1. By comparing the intensity of the graphene features with the LO phonon of the SiC substrate we can infer that these inclusions are a stack of multilayers, thicker in sample #6 than in sample #7. Nevertheless, the G’ peak is almost single Lorentzian, showing no clear signature of graphitic residues as in earlier studies21The micro-Raman spectroscopy results on FEG and MEG samples are summarized in pears ~16 cm−1 in σ1L = σintra + σinter, which includes both intraband and interband processes92728EF, carrier scattering time τ and thermal broadening Γ) on graphene optical conductivity in the THz-MIR range can be visualized in σ(ω) and of both σintra(ω) and σinter(ω) are shown for varying EF. The computed curves are normalized by σ0 = e2/4ħ. The interband threshold is clearly observed at ħω ~ 2EF, and the total conductivity σ(ω) exhibits significant variations over the full THz and MIR ranges, except above the interband threshold, where σ(ω) ~ σ0. The inset shows σ(ω) in the THz range, dominated by the intraband Drude-like absorption. In σ(ω) was computed for a fixed Fermi level (EF = 100 meV) while varying τ and Γ, respectively. Increasing τ impacts σ(ω) in the THz-MIR range, as shown in et al.i.e. a value close to the thermal energy kBT at room temperature.The optical properties of monolayer graphene can be fully described by its complex dynamical in-plane optical conductivity σinter and σintra can have significant contributions over the full spectral range.Note that in most previous studies on the optical response of epitaxial graphene the intra- and interband contributions were treated separately, mostly due to the difficulty to gather data over the full spectral range, e.g. due to the substrate absorption in particular ranges89σ1L(ω)The optical reflectance of graphene layers on a SiC substrate was computed in the framework of the thin film approximation, as shown schematically in the inset of EF, N, τ . In the MIR region, the optical response depends both on intraband and interband transitions, in particular in the Reststrahlen band. For a scattering time τ = 20 fs , the reflectance presents a different trend, as detailed in the upper right inset of EF and considering the Reststrahlen region, R first increases (for EF < 150 meV) and then decreases (for EF > 150 meV). This is due to the intraband transition in the low frequency range, and the interband threshold at higher frequencies: below 2EF, the main contribution to the optical conductivity of graphene arises from the carriers absorption. The latter effect is better visualized in R gradually decreases in the MIR range (in the Reststrahlen region), when increasing EF up to EF ~ 150 meV and rises again when EF increases above this value. The inset of R in the vicinity of the interband threshold (~2EF) when EF is varied in the range 25–400 meV. The effect of increasing τ is shown in Reststrahlen region and in the low frequency limit of the THz range due to the intraband transitions.To investigate on the impact of graphene on the overall optical response, we have computed the reflectance of air/SiC and air/graphene/SiC for varying EF, N, τ . In eachR for increasing N. From monolayer (N = 1) to FEG/MEG, R increases in the THz and MIR regions, with an opposite trend in the Reststrahlen band. The optical contrast increases with increasing N, with respect to the SiC substrate, leading to remarkable changes within this region. The latter feature alone can be used to assess the formation of graphene and evaluate the number of graphene layers. As reported recentlyN, analogous to transmission measurements. For highly doped graphene, this should be performed in the NIR region due to the interband threshold , although the contrast increases with the number of layers. Therefore slight variations in EF, N and τ will result in a unique optical spectrum that can be distinguished from those obtained using different sets of parameters, as will be demonstrated below.R changes markedly with varying N. Moreover, the shape of the curves in the THz range is qualitatively different for each sample, an indication that the carriers properties are distinct. The analysis of the experimental data was accomplished by fitting the reflectance model to the given spectra over the full THz-MIR ranges. We present the obtained fitting curves in N from XPS for samples #6 and #7. Nevertheless, it corroborates with the fits to the IR reflectance data: we indeed find N > 35.To compare with the theoretical predictions, we present in 173233Reststrahlen band are closely related to the properties and number of graphene layers. This can be demonstrated by examining local IR reflectance data from a FEG sample with a heterogeneous number of layers N. Due to its trapezoidal shape, sample #3 was subject to an intentional temperature gradient during graphene growth. We could therefore anticipate that this would give rise to a gradient in N along the sample length. N. We now show that this gradient can be measured thanks to a mapping of reflectance spectra in the MIR range, a rapid and non-invasive method. The number of graphene layers is usually estimated from data in wavenumber range above 3200 cm−1 Reststrahlen band, can be used to estimate N.As shown above, the remarkable changes in the Reststrahlen band: the reflectance is decreasing when going from the left- to the right-hand side of the sample. Fitting the data to the reflectance model yields the spectra presented in N is gradually varying from 2.2 to 14 layers between the two ends of the measured region (see Methods and 2) performed on the same sample, as shown in 2) to approach the experimental conditions of the XPS measurements (1 mm2 spot). Note that for the spectral region analyzed here and our experimental setup configuration, the spatial resolution could be improved to 25–30 μm.IR spectra taken at the positions indicated by the numbered colored squares on EF can be estimated from the threshold in the MIR region. The same kind of analysis can be performed on CVD graphene transferred to other opaque or transparent substrates e.g. Si/SiO2 or quartz. Note that, in our case, the detector used in the IR mapping has a higher sensitivity than the one used in the classical IR spectroscopy. We can thus observe that when N increases, two features appear: a peak at 1584 cm−1, and a broad band centered around 3200 cm−1 (inset of Eu1 (not IR active in monolayer graphene)γ1 ~ 0.37 eV (~2984 cm−1) in bilayer graphene81531Besides being a rapid and non-invasive tool to differentiate between epitaxially grown FEG or MEG samples, it also offers perspectives for mapping the homogeneity of the electronic properties across graphene samples. inset of  The firsEF and τ was obtained through magnetotransport experiments on the FEG sample #1. The Hall coefficient was measured in a Hall bar geometry at low temperature (4.2 K) to avoid parallel conduction through the substrate37EF = 172 meV (with n-type charge carriers), i.e. in excellent agreement with EF = 170 ± 5 meV found from the analysis of IR spectra. In addition, measurement of the sheet resistance at zero magnetic field yielded τ = 12.5 ± 3 fs using the Einstein relationIR = 16 ± 2 fs), considering the differences in experimental conditions (the electron transport experiment probed a few micrometers-squared area).For the sake of comparing the results obtained from the IR analysis, an independent measurement of R , where 1 THz = 33.3 cm−1 = 300 μm = 4.1 meV. The IR beam aperture diameter on the sample was set between ~1.25 and ~2.5 mm2, depending on sample dimensions, and optimized to obtain the best signal-to-noise ratio. All measurements were performed at near-normal incidence (θ ~ 8°). A bare Al mirror was used as reference. In addition, IR spectroscopy mapping was employed to investigate sample homogeneity. In this case, the measurements were performed under ambient conditions using an IR microscope (Bruker Hyperion 3000) coupled to a FTIR spectrometer (Bruker Vertex 80v) under vacuum (~1 mBar). A liquid nitrogen-cooled mercury cadmium telluride (MCT) detector, a KBr beamsplliter and a globar source were used, covering the MIR and NIR regions, between 400 and 7000 cm−1 (12–210 THz or 25–1.4 μm). The IR beam was focused on the sample by a Cassegrain 15x IR objective (NA = 0.4) with a square beam-aperture of 170 × 170 μm2. An Au-coated mirror was used as reference. In all measurements the spectral resolution was set to 4 cm−1. The systematic error in the relative value of the reflectance (for R = 0.985) is around 1% and 2–3% for the IR spectroscopy and microspectroscopy, respectively. Micro-Raman measurements were performed in a LabRam HR spectrometer (Jobin Yvon - Horiba), in a confocal back-scattering setup with a 514 nm (2.41 eV) laser beam focused by a 100× optical objective (NA = 0.95). A high-resolution diffraction grating of 1800 grooves/mm was used. The laser power was kept below 1 mW to avoid damaging of the graphene layers.IR spectroscopy was performed from FIR or THz up to the MIR range in a Fourier-transform infrared spectrometer (FTIR) Bruker IFS 113v under vacuum (~1 mBar) equipped with a liquid helium-cooled Si bolometer, mylar  beamsplitters, a deuterated-triglycine sulfate (DTGS) detector, a KBr beamsplitter, and a silicon carbide light source (globar). The measurements covered the spectral range between 35 and 4500 cmα line. The number of graphene layers was estimated by the integrated intensity ratios of the C1s core level for graphene (~284.5 eV) and SiC (~282.7 eV), as described in Biederman et al.Angle-resolved X-ray photoelectron spectroscopy (XPS) was performed using an Escalab 250Xi instrument from Thermo Scientific and the Al KR between the two media separated by the graphene layer is given by:30σ1L is the conductivity of monolayer graphene, c is the speed of light in vacuum, ε0, ε1 and ε2 are, respectively, the permitivity of vacuum, and the relative dielectric functions of vacuum (ε1 = 1) and of the SiC substrate. We have included the number of graphene layers N in the model by considering that rotational disorder between layers is large enough so there is no Bernal (AB) stacking811354142811354141σFEG/MEG = Nσ1L.The optical response of SiC is described by a Drude-Lorentz model:4445ε∞ is the high-frequency dielectric constant. The first term in the above equation is the well-known four parameters damped oscillator model, where ωLO(ωTO) and ΓLO(ΓTO) denotes respectively the frequency and damping rate of longitudinal (LO) and transverse (TO) optical phonons. The second term describes the Drude absorption due to the free carriers in the n-doped 6H:SiC. All the main features of the reflectance spectrum of SiC were well reproduced using the following fixed parameters: ε∞ = 6.55, ωTO (ωLO) = 797.7 (970.7) cm−1, ΓTO(ΓLO) = 2.3 (5.6) cm−1, ωP = 103 cm−1 and ΓP = 90  cm−1, in good agreement with previous reports323334454749n and average mobility μ are respectively given by n = (N/π)(EF/ħvF)2 and μ = evFτ/ħ(πn)1/2 813F is the Fermi velocity in graphene (~106 m/s). For epitaxial graphene, it was demonstrated that only the first few graphene layers, closer to the substrate, and the outermost layer in MEG are doped, whereas the remaining ones are quasi-neutralsIn the case of two dielectric media 1 (vacuum) and 2 (SiC substrate), as shown schematically in the inset of How to cite this article: Santos, C. N. et al. Terahertz and mid-infrared reflectance of epitaxial graphene. Sci. Rep.6, 24301; doi: 10.1038/srep24301 (2016).


Selective filtration of gas, water, and liquid or gaseous oil is essential to prevent possible environmental pollution and machine/facility malfunction in oil-based industries. Novel materials and structures able to selectively and efficiently filter liquid and vapor in various types of solutions are therefore in continuous demand. Here, we investigate selective gas-water-oil filtration using three-dimensional graphene structures. The proposed approach is based on the adjustable wettability of three-dimensional graphene foams. Three such structures are developed in this study; the first allows gas, oil, and water to pass, the second blocks water only, and the third is exclusively permeable to gas. In addition, the ability of three-dimensional graphene structures with a self-assembled monolayer to selectively filter oil is demonstrated. This methodology has numerous potential practical applications as gas, water, and/or oil filtration is an essential component of many industries. Water-oil separation is a key step in a number of processes, namely oil drilling in exhausted oil fields, tertiary oil extraction, oil spill cleanups, and oil-bearing industrial wastewater purification. In this context, many studies have recently focused on the separation or absorption of water-oil123The volume capacity of existing 2D meshes is however limited owing to the weight of larger mesh substrates. Two-dimensional meshes are furthermore susceptible to damage to any part of the mesh, do not readily filter vapor oil, and separate oil less efficiently than three-dimensional (3D) meshes. Indeed, large volumes are readily filtered through thicker 3D meshes without increasing the size of the substrate. Moreover, in terms of reliability, surface damage does not compromise the filtering ability of 3D meshes because of the underlying connected structure. Finally, the larger reaction surface presented by 3D meshes traps significantly larger volumes of oil vapor. Thus, in order to overcome many drawbacks of 2D filters, 3D structures for water-oil separation were investigated5673O4 supercapacitor electrodes have been prepared by creating 3D graphene networks through chemical vapor deposition14Graphene has superior mechanical, thermal, and electrical characteristics than inorganic materials and its antibiotic properties inhibit the formation of bacteria101112via surface treatments .In this study, 3D graphene structures were employed to selectively filter gas, water, and oil. To that end, three forms of graphene foams were prepared by adjusting their hydrophobicity 2-plasma treatment, the hydrophobic pristine graphene foam, and the omniphobic foam with self-assembled trichlorosilane (HDF-S). As synthesized, 3D structural graphene is hydrophobic but the O2-plasma treatment generates defects on the graphene surface, rendering it omniphilic. For the omniphobic structure, the hydroxyl groups (−OH) on the graphene surface and the active group (trichlorosilane–SiCl3) of HDF-S form covalent bonds from which a strong siloxane network is built up. These three different types of graphene foam allow selective filtering of gas and/or liquids.A schematic diagram and a photograph of a piece of 3D graphene foam 25.4 mm in diameter and 2 mm thick are shown in 2-plasma-treated, pristine, and with self-assembled HDF-S) to water, gasoline, kerosene, and olive oil. The average and standard deviation of five measurements performed for each sample are shown. The contact angle of all the liquids on the O2-plasma-treated graphene foam is 0°, which indicates omniphilic behavior. The contact angles of water on pristine graphene is 107.9 ± 1.1°, which contrasts with the 0° measured for gasoline, kerosene, and olive oil. The (hydrophobic) pristine graphene therefore blocks water only. On the other hand, the 3D graphene covered with self-assembled HDF-S is not only superhydrophobic , but also oleophobic, with gasoline, kerosene, and olive oil contact angles of 78.1 ± 1.4°, 105.6 ± 3.8°, and 121.7 ± 1.7°, respectively. It is thereby constitutes an omniphobic filter.i.e. ~0.22 N/m) for our systems. Based on this criterion, sp2 carbon is the important factor governing the filtration of gasoline, kerosene, and olive oil.The adsorption strength between each solvent and the substrate was investigated from a molecular perspective to elucidate different filtration behavior. First of all, the theoretical balance between surface tension and the gravitational force was calculated in a model pore system to estimate the maximum height of solvent that can be maintained by surface tension alone on a single pore see , suggestIn summary, 3D graphene structures were fabricated whose wettability was controlled using a self-assembled monolayer. The hydrophobicity of the 3D graphene structures is substantially increased by the self-assembled HDF-S surface layer, with the water contact angle increasing from ~107.9° to ~143.2°. Raman analysis shows that the graphene in these structures is uniformly formed and X-ray photoelectron spectroscopy (XPS) measurements demonstrate that HDF-S self-assembles as desired on the graphene. The static apparent contact angle was measured on the graphene foams for liquids with different values of surface tensions, namely water, gasoline, kerosene, and olive oil. Water is selectively filtered by pristine (hydrophobic) graphene, while the HDF-S functionalized (omniphobic) graphene filter blocks all solvents. The porous structure of both filters (with 150–200 μm diameter pores) means, however, that they are both permeable to gas. Molecular dynamics simulations show that this selective filtration stems from the effects of different functional groups grafted onto the graphene. Indeed, while a high-percentage coverage of −OH and −COOH blocks kerosene, water is not filtered out with low-percentage epoxide coverage. This suggests a topic for further investigations, that the surface functionalization of graphene can be tailored to solve particular separation problems involving oil and water mixtures. This approach has numerous potential industrial applications. This study also provides a basis for future advances in gas-oil-water separation methods.2 gas (100 sccm); the substrate temperature was then raised from room temperature to 1020 °C and kept at this temperature, first for 90 min under H2 gas and then for further 30 min under CH4 gas (20 sccm). The CH4 gas flow was then interrupted and the CVD chamber was cooled at a rate of 150 °C/min in H2 environment.Ni foam with a 3D porous structure (pore size: 150–200 μm (120 pores per 25.4 mm), diameter: 25.4 mm, thickness: 2 mm, purity: 99.9%) was used as a catalyst and supporting frame for the fabrication of 3D structural graphene foams. The Ni foam was soaked in acetone and ulrasonicated for 3 min for cleaning and to remove organic contaminants from its surface. The clean foam was then annealed at 600 °C under flowing Ar (100 sccm) for 1 h. To grow graphene, the annealed Ni foam was placed in a chemical vapor deposition (CVD) chamber under H2-plasma treatment . Graphene is hydrophobic as grown. In order to render it omniphobic, HDF-S (Gelest) was used to produce a trichlorosilane monolayer that self-assembles. The fabricated 3D graphene foam was immersed in a mixed solution of anhydrous toluene (99.8%) and silane precursor (3.0 mM), and trichlorosilane was left to self-assemble on the graphene surface in N2 ambient for 1 h. The graphene substrate covered with self-assembled HDF-S was then rinsed three times in anhydrous toluene and twice in isopropyl alcohol, before being left to dry for 1 h under ambient conditions.Omniphilic graphene was obtained by OThe surface topology of the foams was observed by FE-SEM  and the graphene was characterized by Raman spectroscopy. The presence of HDF-S and the surface chemical state of the graphene were verified by X-ray photoelectron spectroscopy ). Contact angles  were measured by placing droplets (5 μL) of water, gasoline, kerosene, and olive oil on the surface of the samples under ambient conditions. For the gas penetration tests, the graphene foams were placed above dry ice in a closed bottle. To reveal the penetration of the different solvents, a water-soluble dye  or an oil-soluble dye  was mixed into appropriate samples.NVT  ensemble at room temperature with 1 fs time steps. The systems were left to relax for 200 ps ahead of the 1 ns production runs used to calculate the adhesion energies. The values reported are averages from two independent simulations and the COMPASS force field was used throughout1718Simulations were performed with four different solvents, namely water, gasoline , kerosene (consisting of n-hexadecane and n-eicosane), and olive oil  .


We investigated single-layer graphene on SiC(0001) by atomic force and tunneling current microscopy, to separate the topographic and electronic contributions from the overall landscape. The analysis revealed that the roughness evaluated from the atomic force maps is very low, in accord with theoretical simulations. We also observed that characteristic electron scattering effects on graphene edges and defects are not accompanied by any out-of-plane relaxations of carbon atoms. Graphene epitaxially grown on a substrate differs in many aspects from free-standing graphene or graphene exfoliated onto insulating surfaces. The influence of the substrate hinders applications of epitaxial graphene in the nanoelectronics . The two× 6 quasi-periodic corrugation (q-6 hereafter) [Understanding the interplay between atomic and electronic structure of graphene grown on SiC substrate is a crucial factor for device construction. In this respect structural properties of graphene were extensively studied by atomic force microscopy (AFM) –9 and screafter) ,13–15. Ireafter) –18. Howereafter) –21.Until now, the morphology of SLG on SiC(0001) has been studied by STM and AFM separately. Therefore, an investigation using the combined STM–AFM technique, that has simultaneous access to the electronic and topographic channels, is needed.×R30°  modulations, characteristic for the areas of graphene edges and point defects. Our results complement the previous findings about the origin of scattering at the defects and also serve for a comparison with properties of different kinds of epitaxial graphene.For the first time we bring experimental data, that can distinguish the topographic landscape from the local electronic structure of SLG on a 6H-SiC(0001) substrate. At room temperature we employed a combined STM and dynamic atomic force microscopy (dAFM) based on the Q-plus sensor working under UHV conditions. We observe the −10 mbar. N-doped Si-face 6H-SiC(0001) wafers purchased from CREE Inc., were cut into 3 mm × 10 mm stripes and mounted onto sample holders constructed for a direct-current sample heating. After inserting into the UHV, the samples were degassed at 600 °C and annealed to 800–850 °C under a Si flux for 30 min, which resulted in a perfect Si type 3 × 3 reconstruction and typical terrace widths above 50 nm. As the source of the Si flux we used a Si wafer, which was annealed up to 1200 °C by passing the direct current. The deposition rate was 1 mL per 5 min. The temperature of the sample was measured using an optical pyrometer with spectral emissivity set to 0.6, operating at 1.6 μm wavelength, focused onto a spot on the sample with a diameter <2 mm. Buffer layer and graphene growth were achieved by further annealing of the sample up to 1150 °C, repeated in 10 min increments, until the desired coverage of graphene was achieved (about 2/3 of the surface) [−1 [f), is the possibility of doing measurements in the region of a negative frequency shift gradient (repulsive regime), even at room temperature, without enhanced risk of losing the tip apex. This approach is specifically chosen for the conditions when the graphene contrast provided by the Δf is not giving atomic resolution in the attractive regime. Kelvin probe force measurements (KPFM) were also done in the constant-height mode, with a slow feedback between the measurement points, to compensate the tilt of the sample. KPFM parabola was measured by sweeping the bias voltage and measuring the Δf value [The experiments were performed in an ultra-high vacuum (UHV) environment with a base pressure not exceeding 1 × 10 surface) –14 and tace) [−1 . The conace) [−1 –23. The Δf value .ss* basis set was used for the H atoms, a sp basis set for the Si and C atoms. The cutoff radii of the pseudo-atomic basis functions were as follows: R = 4.5 a.u., R = 4.5 a.u., R = 4.5 a.u, R = 4.5 a.u., R = 6.0 a.u., R = 6.00 a.u. Local density approximation is used for the exchange correlation functional. Our atomic models of the SLG-SiC(0001) surface consisted of a slab of 1648 atoms, 10 atomic layers thick (including the buffer and SLG layer) with an additional passivating hydrogen layer on the underside. The lateral size of the supercell was 6× 6. The calculations were restricted to the Γ point of the first surface Brillouin zone. The bottom Si atomic layer as well as the hydrogen layer were kept fixed during the geometry optimization while all other atoms were allowed to relax freely into their equilibrium positions. The criterion for terminating the relaxation was that maximal forces on free atoms had to be below 0.05 eV/Å and the change of total energy between subsequent iterations had to be smaller than 10−4 eV per unit cell.In order to estimate the SLG corrugation we carried out large scale total energy density functional theory (DFT) calculations. We used local-orbital FIREBALL code –26. FIRE) modulation corresponding to the 6× 6R30° quasiperiodic structure detected by low-energy electron diffraction [ periodicity. This characteristic feature has been observed before, not only near to the graphene boundaries, but also around single-point defects in graphene and graphite [fraction . Figure graphite ,28–29 angraphite . We alsographite . In contf, time-averaged current (<It>) are measured vs the tip–sample relative distance (Z) during an approach and retraction of the tip. Voltage is kept at 100 mV. The <It> grows exponentially with Z and the Δf has a minimum near to −80 Hz. The value of Z = 0 corresponds to the tip–sample separation at <It> setpoint of 0.46 nA. In the region of the set point, the Δf is rising at a rate of (0.98 ± 0.04) Hz/pm, as determined by the least-square fitting. Force (F) is calculated from the Δf signal, using the approach proposed by Sader et al. [The set of curves in r et al. . Maximum structure of graphene a Kelvin-probe measurement [ modulation) in the Δf image, at 0.5 nm from the setpoint, further from the sample. In To identify any charge transfer effects within the q-6surement  has beenf and <It> maps are presented for the same area of the sample, taken in a constant height regime with a slow feedback setpoint at 0.46 nA, producing a tip–sample separation close to the Z = 0 in It> map shows an image resembling the STM topography in  periodicity. The Δf signal shows the honeycomb structure and the q-6 modulation as well. In this case, the slow feedback control of the tip–sample distance allows us to stay in the repulsive regime, achieving atomic resolution. This would be otherwise possible only at very stable conditions, e.g., at low temperatures. The lower part of the <It> image in  ripples on graphene visible in the <It> channel. In contrast, the Δf image shows only a flat honeycomb structure of graphene.In  modulation of graphene in STM has, apart from a significant electronic contribution, a topographic component. It is a question how strong is the topographic effect and whether it can be evaluated from the semi-constant height data. Neglecting the effect of the slow feedback and by using the Δf(Z) curve from Z height modulation of graphene from the Δf images in f region, which has a of 1.0 Hz/pm, we use the Δf roughness corresponding to the q-6 modulation (5.7 Hz RMS), to obtain an average value of about 6 pm RMS, corresponding to 16 pm corrugation (peak-to-peak). This is significantly lower than the roughness obtained by standard topographic mode of STM, pointing out the dominant electronic contribution of the buffer layer to the STM topography of SLG. In our fully relaxed theoretical model of SLG on the buffer layer , which corresponds well to the experimental estimation and is comparable to a similar calculation in [The experimental results reveal that the q-6er layer , which iation in . The valation in –39. pattern . We can conclude that variation at this scale detected by previous STM measurements is caused mainly by electronic contributions arising from the buffer layer. In the case of the characteristic  rippling of graphene, which occurs due to presence of defects or boundaries, any relaxations of carbon atoms perpendicular to the surface can be excluded. These results are consistent with recent findings and show the single-layer graphene on SiC(0001) as a morphologically very flat substrate.In summary, we successfully probed the epitaxial graphene on SiC(0001) by a combined STM and dAFM technique, gathering extended information on its atomic structure near defects. By atomic force microscopy, we detected a topographical corrugation of the q-6The research was funded by GACR grant no.14-02079S, GACR grant no. EXCELENCE 14-374527G.


This paper identifies trends within and relationships between the amount of participation and the quality of contributions in three crowdsourced surveys. Participants were asked to perform a collective problem solving task that lacked any explicit incentive: they were instructed not only to respond to survey questions but also to pose new questions that they thought might-if responded to by others-predict an outcome variable of interest to them. While the three surveys had very different outcome variables, target audiences, methods of advertisement, and lengths of deployment, we found very similar patterns of collective behavior. In particular, we found that: the rate at which participants submitted new survey questions followed a heavy-tailed distribution; the distribution in the types of questions posed was similar; and many users posed non-obvious yet predictive questions. By analyzing responses to questions that contained a built-in range of valid response we found that less than 0.2% of responses lay outside of those ranges, indicating that most participants tend to respond honestly to surveys of this form, even without explicit incentives for honesty. While we did not find a significant relationship between the quantity of participation and the quality of contribution for both response submissions and question submissions, we did find several other more nuanced participant behavior patterns, which did correlate with contribution in one of the three surveys. We conclude that there exists an optimal time for users to pose questions early on in their participation, but only after they have submitted a few responses to other questions. This suggests that future crowdsourced surveys may attract more predictive questions by prompting users to pose new questions at specific times during their participation and limiting question submission at non-optimal times. Crowdsourcing  holds thet al.. Tab. Tab29].R2 values). We hypothesized that these poor results were due to the one participant who submitted 318 of the 591 questions to the survey. Due to one individual submitting so many questions, many of the eight variables described above had nearly identical values for all 318 observations in our dataset. For this reason, this participant was removed from the dataset and only the remaining 273 questions were considered.However, we did eventually find useful insight from the EnergyMinder dataset. Initially, we did not find consistent trends in the modeling results  In addition, the finding that some users submit many more questions than others without a measurable increase in quality suggests that it may be beneficial to limit question submission once a user has already participated a great deal.We obtained permission to conduct human subjects research from the Institutional Review Boards at the University of Vermont. Our IRB proposal was entitled “Wikinnovation: Automating and Democratizing Online Survey Construction.” Participants granted written consent by clicking on a button to proceed to the online survey after watching a video outlining the purpose of the crowdsourced surveys.https://github.com/BobFromUVM/crowdsurvey2014. All images used in this paper are also accessible from the above URL under the MIT license.The data obtained from our three crowdsourced surveys is freely accessible under the MIT license from S1 File(PDF)Click here for additional data file.


Hepatocellular carcinoma (HCC) accounts for approximately 80% of liver cancers, the third most frequent cause of cancer mortality. The most prevalent risk factors for HCC are infections by hepatitis B or hepatitis C virus. Findings suggest that hepatitis virus-related HCC might be a cancer in which LINE-1 retrotransposon, often termed L1, activity plays a potential role. Firstly, hepatitis viruses can suppress host defense factors that also control L1 mobilization. Secondly, many recent studies also have indicated that hypomethylation of L1 affects the prognosis of HCC patients. Thirdly, endogenous L1 retrotransposition was demonstrated to activate oncogenic pathways in HCC. Fourthly, several L1 chimeric transcripts with host or viral genes are found in hepatitis virus-related HCC. Such lines of evidence suggest a linkage between L1 retrotransposons and hepatitis virus-related HCC. Here, I briefly summarize current understandings of the association between hepatitis virus-related HCC and L1. Then, I discuss potential mechanisms of how hepatitis viruses drive the development of HCC via L1 retrotransposons. An increased understanding of the contribution of L1 to hepatitis virus-related HCC may provide unique insights related to the development of novel therapeutics for this disease. Hepadnaviridae family, which has a relaxed circular DNA (rcDNA) as a viral genome , accounts for 9% of all cancer deaths worldwide  retrotransposons are genetic elements that constitute approximately 17% of the human genome : preC/C (core and Hepatitis B e-Antigen [HBeAg]), P , preS/S (Hepatitis B surface Antigen [HBsAg]) and X  genes Figure . In the Type I interferons (IFNs) play a major role in anti-viral immunity  and non-structural  proteins Figure . The repAlthough it is thought that non-retroviral RNA viruses are not integrated into host genomic DNA, we and others have demonstrated that they do become integrated into the host genome via host retrotransposon machineries (Geuking et al., Most HCV-infected patients develop a chronic infection, suggesting that HCV has developed successful strategies to evade host immune responses (Gale and Foy, The HCV core protein has an oncogenic potential (Moriya et al., de novo insertions identified in cancer tissue occurred at genes commonly mutated in cancer (Lee et al., de novo insertions can affect the expression of nearby genes and the genes in which they inserted (Lee et al., de novo insertion (Ding et al., Although a definitive role for L1 activity in contributing to HCC etiology has not been established thus far, investigating a possible link between L1 activation and the development of HCC would be of considerable interest for a number of reasons Figure . FirstlyTaken together, I conclude two potential roles for L1 elements in the development of hepatitis virus-related HCC. The first relates to a chimeric transcript specific to HBV-related HCC, HBx-L1, which can be detected in more than 20% of HBV-related HCC and correlates with a poor outcome (Lau et al., Presented lines of evidence suggest potential links between hepatitis virus infection and L1 retrotransposon activity. Especially, L1 hypomethylation or some L1 chimeric transcripts are associated with poor prognosis of HCC, suggesting that it can have a significant effect on HCC phenotypes and supporting the idea that HCC is a cancer in which L1 plays a role. However, knowledge of how L1 activation by chronic hepatitis virus infection enhances the development of HCC is still limiting. Further accumulation of examples of recurrent L1 insertion sites in the host genome or recurrent chimeric transcripts specific to hepatitis virus-related HCC will be promising ways to understand L1 involvement in HCC etiology. Single cell analyses of L1 retrotransposition events and expression in tumor cells and surrounding normal cells may enhance these processes. Understanding the potential roles of L1 in HCC may open avenues to developing novel therapeutics, such as RNA interference against HCC-specific L1 chimeric transcripts.TH wrote the manuscript and approved it for publication.The author declares that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.


In this context, the control of the charge transfer between the constituting layers is of crucial importance. Compared to bulk or 0D system, 2D materials are characterized by a large exciton binding energy (0.1–1 eV) which considerably affects the magnitude of the charge transfer. Here we investigate a model system made from colloidal 2D CdSe nanoplatelets and epitaxial graphene in a phototransistor configuration. We demonstrate that using a heterostructured layered material, we can tune the magnitude and the direction (i.e. electron or hole) of the charge transfer. We further evidence that graphene functionalization by nanocrystals only leads to a limited change in the magnitude of the 1/f noise. These results draw some new directions to design van der Waals heterostructures with enhanced optoelectronic properties.Two dimensional layered ( One of the promising next steps is their integration in optoelectronic devices such as photodetectors or photovoltaic cells128 and responsivity of 107 A.W−1 can be achieved−14 W.cm−2), thanks to the high mobility of the graphene channel and the long carrier lifetime of the photogenerated carriers in the CQD layer.As an alternative, van der Waals heterostructures can be used to decouple the photogeneration from the transport. Such devices can be built by coupling CQD with efficient transport layer such as graphene1011132 or WS2, we use colloidally grown CdSe nanoplatelets (NPL). In addition to their 2D character2527In spite of all these progresses, the performances of photoconductors are limited by several parameters. Some of them have already been addressed212223Colloidal metal-chalcogenides nanoplatelets (NPL) such as CdSe offer excellent optical properties originating from the control of their thickness at the atomic level2914 cm−2 which will be of utmost interest to evidence a strong charge transfer. We use a fast sweep rate (50–100 mV.s−1) only allowing the gating of graphene why leaving the charging state of the NPL unaffectedAs a probe for the charge transfer, we choose to use an electrolytic transistor configuration33n doping of the channel where C is the surface capacitance of the electrolyte estimated to be 2 illumination, the negative Dirac point shift corresponds to a Fermi level change from −4.5 eV to −4.3 eV, that is to say that about 1012 electrons.cm−2 are transferred from the NPL to the graphene channel. As a comparison, under the same illumination condition, 109 charges.cm−2 are generated in a film of NPL  in the conduction band  of the NPL and the graphene, λ the reorganization energy and Ea the activation energy. The latter can be writtenwhere EExc being the exciton binding energy and ΔE = sign(charge)×,NPL) the difference between graphene Fermi level and the conduction  band of NPL. Faster charge transfer consequently requests to combine a weaker exciton binding energy with almost resonant coupling of one type of carrier to graphene. The second band of the semiconductor has to stay out of resonance to avoid charge recombination within the graphene. As proposed for quantum well structures, we demonstrate that the use of a heterostructure is well suited to achieve charge displacement within the semiconductorCdSe NPLs with their 2D aspect and limited static dielectric (ε = 10) constant55727EExc and the oscillator strengthfthanks to the relation 65859To quantify the reduction of the exciton binding energy, we use time resolved photoluminescence (PL) spectroscopy, see ε0 the vacuum permittivity, m0 the rest mass of the electron, n the optical index of the medium. The oscillator strength can be estimated aswith Ep being the Kane parameter and Ψe/h the electron or hole wavefunction. Definingas the overlap integral, the radiative lifetime is now given byEExc value is twice smaller than for core only NPL  noise nor by the shot noise because graphene possesses 1/f excess noise at frequencies below a few kHzV between the middle of the two branches are the averaged noise coming from the 4 graphene channels, i.e. VI and presented in equation 14where R = 8kΩ. The value of the power exponent is γ = 1.05 ± 0.05 as expected for 1/f noise without generation-recombination bulges. The magnitude β of the noise is plotted versus the applied voltage bias in the VDS = VI/2 is the voltage applied on each graphene channel. This gives an exponent value c = 1.9 ± 0.1 and a coefficient A0 = 6.5 ± .5×10−9, consistent with previous report of 1/f noise in grapheneN, we can calculate the α parameter characterizing the intrinsic noise of a sample:is performed on each measurement. The obtained thermal noise 2 and the carrier density of bare graphene on SiC being 8.1012 cm−2  choose the availability of one or the other type of carrier toward charge transfer. Applied to a photodetector made from CdSe-based nanoplatelets decorated on a graphene channel, CdS-shell covered NPL induce a more efficient n-type photogating of the graphene while CdTe-crown surrounded NPL produce a p-type photogating. We expect this strategy to be profitable to other 2D materialsHeterostructures grown at the nanocrystal scale is a practical strategy to 1) lower the exciton binding energy by delocalizing one type of carrier How to cite this article: Robin, A. et al. Engineering the Charge Transfer in all 2D Graphene-Nanoplatelets Heterostructure Photodetectors. Sci. Rep. 6, 24909; doi: 10.1038/srep24909 (2016).


Exfoliation of graphite is a promising approach for large-scale production of graphene. Oxidation of graphite effectively facilitates the exfoliation process, yet necessitates several lengthy washing and reduction processes to convert the exfoliated graphite oxide  to reduced graphene oxide (RGO). Although filtration, centrifugation and dialysis have been frequently used in the washing stage, none of them is favorable for large-scale production. Here, we report the synthesis of RGO by sonication-assisted oxidation of graphite in a solution of potassium permanganate and concentrated sulfuric acid followed by reduction with ascorbic acid prior to any washing processes. GO loses its hydrophilicity during the reduction stage which facilitates the washing step and reduces the time required for production of RGO. Furthermore, simultaneous oxidation and exfoliation significantly enhance the yield of few-layer GO. We hope this one-pot and fully-scalable protocol paves the road toward out of lab applications of graphene. The major differences between them are the yield and the defect content of their products. Although the use of less-defective graphene is essential in some applications (e.g. electronic devices), in some other cases, the defects and imperfections of graphene are less important or even desirable9Due to the lack of a reliable method to supply the large demand for graphene, we have developed a fast and fully-scalable method for mass production of RGO. The consecutive stages of RGO production are shown in The main intent of this article is to propose a general approach for large-scale production of RGO so that the main idea, reduction of GO prior to washing, can be extended to most of oxidation-based synthesis methods of RGO. Ascorbic acid as a green alternative for conventional reducing agents of GO (e.g. hydrazine hydrate) is an inexpensive and abundant substance which can reduce GO to an acceptable extent{1100}/I{2110} > 1). Diffracted intensity profile along the passing line from  was characterized by various chemical and physical methods. A sample was also taken from the oxidized graphite (AGO) prior to the reduction step with a great exertion for washing [see Methods]. The morphology and microstructure of ARGO sheets were investigated using field emission scanning electron microscopy (FE-SEM) and transmission electron microscopy (TEM) techniques. Well-exfoliated but crumpled and aggregated RGO sheets are observable in the SEM image , which i1522X-ray diffraction (XRD) patterns of pristine graphite, AGO, and ARGO are represented in 2O, CO2, and CO gases. The weight loss with a slower rate at >300 °C is assigned to the removal of more stable OFGsThermogravimetric analysis TGA,  demonstr−1 (G band) which is due to the first-order scattering of the E2g mode and two weaker peaks located at 1340 cm−1 and 2687 cm−1  originated from the second-order double resonant process between non-equivalent K points in the Brillouin zone of graphene−1 in the Raman spectrum of AGO. The blue shift of the G band has been attributed to the activation and merging of the Raman-inactive D′ band with the G band, the exfoliation of graphite to a single graphene sheet, and the resonation of isolated double bonds at higher frequencies than the G band of graphite−1 was observed that can be ascribed to the reduction of AGO. Such a shift back has been reported in both chemical and thermal reductions of GO2D/IG), which is an estimation of the disorder level in graphene, exhibits a non-monotonic behavior. In a low concentration of defects, the introduction of new defects creates more elastic scattering, resulting in an increase of ID/IG. Further addition of defects leads to a more amorphous structure, attenuating all Raman peaks and decreasing ID/IG . The peak at 1622 cm-1 is attributed to the C = C stretching vibration. Other prominent signals in the AGO’s spectrum which have been attenuated in the ARGO’s spectrum are 1399 cm-1, 1240 cm−1, and 1071 cm-1 originating from the O–H deformation, the C–O (epoxy) stretching vibration, and the C–O  stretching, respectivelySince the attachment of the OFGs to the graphite layers during the oxidation dramatically affects the structure and properties of graphene, their removal is a great concern for a successful RGO production. Although numerous models for the chemical structure of graphene oxide have been proposed to date, there is still a considerable vagueness. Nevertheless, it is widely accepted that the hydroxyl and epoxide groups attached to the basal plane of the graphene and carboxyl and carbonyl groups located at the edges are the dominant functional groups21281S peaks were deconvoluted using Gaussian-Lorentzian functions  is a widely accepted criterion for evaluation of the reduction process. According to the XPS survey spectra , the C/Ounctions . The ass spectra . The cheAs an important characteristic of the synthesis method, the powder yield was determined by performing a complete process, starting with 5 g graphite and obtaining ~6.37 g ARGO powder, which indicated an extremely high powder yield of ~127%. The yield of higher than 100% is attributed to the remaining of some functional groups as described earlier. In order to evaluate the purity of the final product, the metallic ion contents of the ARGO, especially Mn, K, and Na ions which had been introduced to the system during the synthesis procedure were determined using inductively coupled plasma optical emission spectrometry (ICP-OES) and the results are given in In summary, the oxidation of graphite significantly increases its interlayer spacing and effectively facilitates the exfoliation process, but necessitates various post-treatments to remove the chemical oxidants which extremely prolongs the production time. While the use of filtration and centrifuge systems in typical oxidation-based methods are indispensable, they limit the production yield to few grams a day! Although the main idea behind the present method is simple, it removes a huge obstacle from conventional synthesis methods of graphene. GO loses its hydrophilicity which makes the whole process up-scalable and one-pot oxidation-exfoliation-reduction process, considerably reduces the overall production time from several days to few hours, hence kilogram or even ton scale production of RGO is simply attainable just by increasing the capacity of the reaction vessels. The SEM, HRTEM and SAED analyses clearly demonstrate that this method can produce crystalline two-dimensional graphene platelets. The FT-IR, Raman, TGA and XPS analyses verify the successful reduction of GO even in the presence of oxidative agents and the ICP-OES analysis confirms the high purity of the final product. The characteristics of the synthesized product are comparable with other high quality chemically-synthesized RGOs.In a typical procedure, 1 g of graphite flakes  was added to 50 mL concentrated sulfuric acid  while stirring in an ice-water bath. 3 g potassium permanganate  was gradually added by maintaining the temperature under 10 °C. Then, the suspension was stirred at room temperature for 25 min followed by 5 min sonication in an ultrasonic bath . After repeating the stirring-sonication process for 12 times, the reaction was quenched by the addition of 200 mL distilled water. An extra 2h ultrasonic treatment was carried out before dividing the suspension into two equal parts; one washed to obtain AGO (described later) and the other was further processed for preparation of ARGO. After adjusting the pH at ~6 by the addition of 1M sodium hydroxide  solution, the suspension was further sonicated for 1 h. 10 g L-ascorbic acid  was dissolved in 100 mL distilled water and then was slowly added to the exfoliated graphite oxide suspension at room temperature. The reduction was performed at 95 °C for 1 h. The resultant black precipitates were simply filtered by cellulose filter paper and further were washed with a 1M hydrochloric acid solution  and distilled water to neutral pH. Finally, the filtrate was freeze-dried to obtain ARGO powder.The washing of AGO was adapted from a modified Hummers methodα radiation (λ = 1.5406 Å). The investigation of thermal properties was carried out using TGA  under air with a ramp rate of 1 °C min−1. Raman spectroscopy measurements were taken using a Renishaw InVia Reflex Raman microscopy system with a 532 nm laser. FT-IR spectra were collected using a Bruker Tensor 27 instrument. The XPS spectra were recorded using a VG Scientific-ESCA Lab 250 XPS spectrometer with a monochromatic Al Kα radiation source (1486.8 eV). The Spectral Data Processor  software was employed for curve fittings and atomic percent calculations of XPS spectra. The ICP-OES tests were performed using a Varian Vista-Pro Axial. The ICP-OES samples were digested in 2 mL 35% H2O2 and 4 mL 65% HNO3 using a microwave digestion system .SEM images were taken by a Leo Supra 35VP field emission scanning electron microscope. A JEOL JEM 2100F transmission electron microscope was used for TEM imaging, acquiring SAED patterns, and HR-TEM analysis. XRD patterns were recorded with a Bruker AXS advance powder diffractometer equipped with a Siemens X-ray gun, using Cu KS.A. and H.A. conceived and designed the research. S.A. and H.S.K. carried out the experiments. S.A. and H.A. interpreted the results and wrote the manuscript. All authors reviewed and approved the final version of the manuscript.How to cite this article: Abdolhosseinzadeh, S. et al. Fast and fully-scalable synthesis of reduced graphene oxide. Sci. Rep.5, 10160; doi: 10.1038/srep10160 (2015).


Therapeutic options for patients with amyotrophic lateral sclerosis (ALS) are currently limited. However, recent studies show that almost all cases of ALS, as well as tau-negative frontotemporal dementia (FTD), share a common neuropathology characterized by the deposition of TAR-DNA binding protein (TDP)-43-positive protein inclusions, offering an attractive target for the design and testing of novel therapeutics. Here we demonstrate how diverse environmental stressors linked to stress granule formation, as well as mutations in genes encoding RNA processing proteins and protein degradation adaptors, initiate ALS pathogenesis via TDP-43. We review the progressive development of TDP-43 proteinopathy from cytoplasmic mislocalization and misfolding through to macroaggregation and the addition of phosphate and ubiquitin moieties. Drawing from cellular and animal studies, we explore the feasibility of therapeutics that act at each point in pathogenesis, from mitigating genetic risk using antisense oligonucleotides to modulating TDP-43 proteinopathy itself using small molecule activators of autophagy, the ubiquitin-proteasome system, or the chaperone network. We present the case that preventing the misfolding of TDP-43 and/or enhancing its clearance represents the most important target for effectively treating ALS and frontotemporal dementia.The online version of this article (doi:10.1007/s13311-015-0338-x) contains supplementary material, which is available to authorized users. Amyotrophic lateral sclerosis (ALS) is the most common adult-onset motor neuron disease, and is characterized by the progressive loss of upper and lower motor neurons from the spinal cord, brain stem, and motor cortex, leading to muscle weakness and eventual respiratory failure. Approximately 5–10 % of ALS cases are familial with the remaining 90 % being sporadic, indicating that both genetic and environmental factors contribute to risk. Despite this diverse etiology of disease, 97 % of patients display a common phenotype in disease-affected tissues, namely the deposition of the TAR-DNA binding protein (TDP)-43 . DepositTARDBP, TDP-43 is a ubiquitously expressed DNA-/RNA-binding protein [Encoded by  protein , and a g protein , 7. TDP- protein . In the  protein , 9. TDP- protein , 11. In  protein , 12–15, Although mostly nuclear, up to ~30 % of TDP-43 protein can be found in the cytoplasm , with nuSOD1 or FUS [TDP-43 protein was identified as a major component of the ubiquitinated neuronal cytoplasmic inclusions deposited in cortical neurons in FTD and in spinal motor neurons in ALS . TDP-43-1 or FUS –28. Neur1 or FUS . This tr1 or FUS , which sHowever, a key question in ALS research is which of these features of TDP-43 proteinopathy are required for the development of disease and thus represent therapeutic targets. Is ALS pathogenesis linked to the loss of wild-type TDP-43 function through protein misfolding and failure to interact with binding partners, or is it linked to a gain of toxic function of the aforementioned TDP-43 aggregates, which are the hallmark of TDP-43 proteinopathy? A number of studies have examined the roles of TDP-43 gain or loss of function in disease. Overexpression of wild-type TDP-43 recapitulates TDP-43 proteinopathy and disease phenotypes in a range of animal models –33, suppWe have introduced the concept that TDP-43 is the convergence point for a range of upstream risk factors for ALS. Here we briefly review these genetic and environmental risk factors for the development of ALS, and for the development of TDP proteinopathy, and explore the potential for targeting therapeutics towards these diverse risk factors.Currently, genetic causes are known for approximately 15 % of all ALS cases; accounting for 11 % of sporadic ALS and 68 % of familial ALS . Many ofTARDBP are a rare cause of ALS [TARDBP mutations have been identified in both familial and sporadic ALS, most clustering in the region encoding the C-terminus, and accounting for approximately 1–2 % of total cases [TARDBP-linked ALS patient tissue is characterized by cytoplasmic accumulation as aggregated and insoluble deposits [Mutations in e of ALS . To dateal cases –58. Likedeposits , 55, nucdeposits , and C-tTARDBP-linked disease. Loss of RNA processing and axonal transport function of mutant TDP-43 has been suggested to precede other features of pathology [TARDBP mutations cause loss of function [TARDBP mutations cause ALS provides robust evidence that altered TDP-43 structure  and the resultant loss and gain of function is not simply a cellular response to disease but is pathogenic.Cellular and transgenic animal models indicate that both overt proteinopathy, as well as preproteinopathic changes in TDP-43 may be at play in athology , 19, 59,function , 39. Howfunction , 46, 60,function , 61, 62,function , 63–65, function –69. ThesTARDBP mutations linked to disease.We previously demonstrated that a mutant TDP-43 allele (M337V) can be selectively silenced using small interfering RNA , and, inSQSTM1 [VCP [UBQLN2 [OPTN [ALS-linked mutations in SQSTM1 , VCP [73 [UBQLN2 , and OPTN2 [OPTN  are rareN2 [OPTN –80. TDP-N2 [OPTN –87.Notably, VCP and p62 are required for the formation of “aggresomes” , 88, 89,VCP mutation is the nuclear accumulation of TDP-43 rather than VCP protein itself or other degradation substrates [But is the removal of TDP-43 a logical approach in ALS linked to global impairment of protein degradation, when TDP-43 is just one of many substrates of these systems? At the high concentrations caused by global impairment in protein degradation, TDP-43, like other neurodegenerative disease proteins –95, and bstrates . TherefoCertainly, small molecule activators of the UPS or autophagy have been shown to promote TDP-43 clearance and/or mitigate toxicity in models based on TDP-43 overexpression , 98, 99.The identification of disease-linked mutations affecting other RNA-binding proteins suggests an important role for deficits in RNA processing, which is a key TDP-43 function, in disease pathogenesis. ALS-linked mutations in the RBP fused in sacroma (FUS) lead to its mislocalization from the nucleus to the cytoplasm and aggregation , but theC9ORF72 [C9ORF72-linked disease can be distinguished by the additional presence of nuclear foci of repeat-containing RNA [C9ORF72 cases also harbor TDP-43-negative inclusions throughout the central nervous system, which are decorated by ubiquitin [C9ORF72-mediated disease [The most common ALS-linked mutation is an intronic GGGGCC repeat expansion in C9ORF72 , 103, thC9ORF72 . C9ORF72ning RNA . These Rning RNA , 105, lening RNA , 107. C9biquitin , p62 [14biquitin , and/ orbiquitin . These ibiquitin –112. Ovebiquitin , 114, cabiquitin , 116. Hobiquitin , 118. Ma disease , 120, TDMATR3 and hnRNPA1 are also associated with ALS and TDP-43 proteinopathy [FUS and C9ORF72, these ALS-linked genes indicate that RNA processing deficits can be a cause of ALS in the presence or absence of TDP-43 proteinopathy, and that rescue of RNA processing defects could be beneficial to patients. As is the case for C9ORF72-linked disease, the extent to which RNA processing defects are due to TDP-43 proteinopathy will determine whether TDP-43-based therapeutics are effective in patients.Mutations in the RNA-binding protein genes inopathy , 122, poinopathy , 121. TaMAPT [For the majority of ALS cases, no genetic mutations have yet been identified to account for disease; therefore, the development of ALS and TDP-43 proteinopathy in susceptible individuals is thought to also involve environmental factors. This idea was first proposed following observations of unexpectedly high disease incidence rates in certain “hotspots”, such as the Kii peninsula of Japan and the island of Guam , 124. WhMAPT –128, theMAPT , 130, orMAPT .There is an association between US military service and increased risk of ALS , which mIndeed, cellular studies have drawn links between a diverse range of stressors and the properties of TDP-43 protein. Osmotic stress , oxidatiThe vast majority of ALS and FTD cases are of unknown etiology but are linked by TDP-43 proteinopathy, which is defined by cytoplasmic mislocalization, fragmentation, aggregation, and post-translational modification. Here we examine how each of these features of pathological TDP-43 is linked to ALS pathogenesis, and whether preventing these common features might be a valid therapeutic strategy.Mislocalization of TDP-43 in ALS and FTD is evidenced by the deposition of granular, skein-like, and macroaggregated TDP-43 in the cytoplasm, as well as clearing of TDP-43 from the nucleus . Enhancein vivo. However, studies that demonstrated that link had targeted autophagy or interactions with stress granule components rather than TDP-43 nucleocytoplasmic shuttling per se [As a very early pathogenic event cytoplasmic mislocalization is a desirable intervention point, and selectively reducing mislocalized TDP-43 is linked to reduced toxicity g per se , 141. WiPhosphorylated C-terminal fragments (CTFs) of TDP-43 are a major constituent of neuronal protein inclusions in ALS and FTD brains, but less so in spinal cord , 149. ClBut is fragmentation of TDP-43 a target for therapeutic intervention? There is no current consensus as to whether cleavage enhances or mitigates TDP-43 toxicity , 155. ItThe role of TDP-43 aggregation in pathogenesis is one of the most controversial topics in ALS research; thus, preventing aggregation as a therapeutic strategy is equally controversial. Several animal studies have found that mutant TDP-43 causes toxicity in the absence of visible aggregates , and preMisfolded mutant TDP-43 shows a reduced ability to transport RNA appropriately , 19, conThere are several intrinsic cellular mechanisms that can act to either prevent or resolve protein misfolding, namely the chaperone system, autophagy, and the UPS. The chaperone system maintains proper protein folding during synthesis and thereafter, or delivers misfolded substrates for degradation . In the Augmentation of chaperones or protein degradation pathways has been protective in several models of ALS and FTD , 162. BeIn ALS and FTD patient tissue, hyperphosphorylation and ubiquitination are signatures for pathological TDP-43, as they preferentially label TDP-43 that is cleaved, aggregated, and detergent-resistant –166. PhoThere are 29 phosphorylation sites on TDP-43 for casein kinase 1 alone , the besDrosophila, despite ubiquitin chain retention [The influence of ubiquitination on TDP-43 proteinopathy is less well studied. TDP-43 is modified with polyubiquitin chains that are predominantly K48- or K63-linked , 76, 173etention . The proetention , remainsC9ORF72 patients. Similarly, dysregulated RNA processing almost certainly lies at the heart of pathogenesis, but its widespread downstream effects argue against rescue of selected transcripts being useful in patients. Therapeutic strategies would best be directed at a common target proximal to the deficit, which, for most cases, is the misfolding of TDP-43 that is central to its loss of function and gain of function/dominant-negative toxicity. Specifically, enhancers of chaperone-dependent TDP-43 folding, as well as activators of the UPS and autophagy, have shown most promise in model systems. Ongoing refinement of model systems in which to test therapeutics, and recognition of the roles of non-neuronal cell types will be important in bringing these compounds to preclinical and clinical testing. In addition, unbiased large-scale compound screening efforts and the identification of novel causative genes may yield additional insights into disease mechanisms and the role of TDP-43.The contributions of genetic and environmental factors to the etiology of ALS and FTD are complex and interwoven. The increasing accessibility of genotyping and the recent demonstration of safe gene silencing using antisense oligonucleotides may render the targeting of individual gene mutations feasible, but currently this is not the case, except perhaps for ESM 1(PDF 1225 kb)Below is the link to the electronic supplementary material.