TDP-43 is a DNA/RNA binding protein associated with TDP-43 proteinopathies. Many mutations have been identified in the flexible C-terminal region, which is implicated in the disease pathology. We investigated four point mutations in the amyloidogenic core region (residues 311–360) of TDP-43 by biochemical and spectroscopic methods. We found that the G335D mutation enhances the aggregation and inclusion formation of TDP-43 and this mutant in TDP-35  exaggerates the antagonist effect on RNA processing by endogenous TDP-43; whereas Q343R gives an opposite effect. As a comparison, M337V and Q331K have very little impact on the aggregation and inclusion formation of TDP-43 or TDP-35. NMR structural analysis showed that the G335D mutant in the core region forms a loop linker between the two α-helices and promotes α-to-β transition, but Q343R loses the second helix and consequently the structural transformation. Thus, the propensity of structural transformation in the amyloidogenic core of TDP-43 determines its aggregation and inclusion formation. This study may provide a molecular mechanism of the TDP-43 proteinopathies caused by genetic mutations. TDP-43 (TAR DNA-binding protein of 43 kDa) is a DNA/RNA binding protein containing two RNA recognition motifs (RRMs) that participate in RNA binding and a long C-terminal glycine-rich region (GRR) that is involved in protein-protein interactions1426TDP-43 is mainly localized in nucleus and plays multiple roles in transcriptional repression, pre-mRNA splicing, translational regulation and stability8910121351517181921222728More than 40 mutations in TDP-43, mainly in the C-terminal GRR domain, have been identified in familial and sporadic cases of ALS and FTLD723242731in vitro and in cells. Our findings implicate pathological impact of the mutations on the inclusion formation and dysfunction of TDP-43, and bring light to the pathogenesis of the related neurodegenerative diseases.Previously, we identified an amyloidogenic core (residues 311–360) in the C-terminal flexible region of TDP-43E. coli to report the aggregation properties of the fusion proteins. As our previous reportAs preciously described, many mutations in the C-terminal GRR domain of TDP-43, especially in its amyloidogenic core, have been identified in familial and sporadic cases of ALS and FTLD7in vitro by ThT assay. Firstly, we analyzed time courses of the core fragment and its mutants fused with thioredoxin (Trx-TDP(311–360)) upon thrombin cleavage ) forms a helix-turn-helix structure 25. The tWe next investigated structural transformation of the mutant peptides of TDP311–360) during the aggregation processes as detected by circular dichroic (CD) spectroscopy . Previou0 during It has been demonstrated that the C-terminal 35-kDa fragment (TDP-35) of TDP-43 triggers formation of cytoplasmic inclusionsTo confirm the above observation, we transfected FLAG-tagged TDP-35 and its mutants respectively into HEK 293T cells, and then performed supernatant/pellet fractionation analysis . The datAs reported previously, TDP-43 is mainly localized in the nucleus that can be extracted into the supernatant. However, when treated with staurospine (STS), TDP-43 is degraded to fragments that cause formation of the cytoplasmic inclusions and deposit in the pelletAs previously reported, TDP-43 can promote CFTR exon 9 alternative splicing and TDP-35 loses the function of RNA processing811As known, TDP-35 may act as an antagonist against endogenous TDP-43 in RNA processing by a dominant-negative mechanismThere are many pathologically identified mutations in the amyloidogenic core of TDP-4333383932397Structurally, the amyloidogenic core region of TDP-43 forms a helix-turn-helix structure in solution and it is prone to transformation into a hairpin-like β-sheetTDP-43 is a nucleic-acid binding protein with RNA processing function in nucleus. Mutation, fragmentation and mislocation will alter RNA splicing and cause neurodegeneration45Although cytoplasmic inclusions of wild-type TDP-43 are presented in 90% of ALS and 50% of FTLD patients4343313846Nde I/BamH I cloning sites to make the GFP-fusion constructs. The cDNA of TDP(311–360) was subcloned into pET32M to get a thioredoxin (Trx)-fused protein, and it was also cloned into pHGB vector to generate GB1-TDP(311–360) with an N-terminal His tag. The TDP-43 and TDP-35 cDNAs were constructed into pcDNA3.1-Myc/His and pcDNA3.1/ FLAG vectors (Invitrogen) respectively for eukaryotic expression. All mutants  were generated by PCR site-directed mutagenesis. All constructs were verified by DNA sequencing.The cDNAs of human TDP-43 and its fragments were digested and ligated into a pET22b-GFP plasmid by using The method for assaying the aggregation abilities of GFP-fused TDP-43 fragments and their respective mutants were performed as described previously2+-NTA following an FPLC Superdex-75 column (GE Biosciences) respectively, and their concentrations were determined spectrophotometrically by using each extinction coefficients.The Trx and GB1 fused proteins were expressed and purified as described previouslyThe time course of the aggregation process was monitored by ThT fluorescence2/dmol). For the solid-state CD measurements, about 150 μL of the protein solution (ca. 0.50 mg/mL) was cast onto a 2-cm diameter cylindrical quartz glass for evaporating overnight at room temperature, and then the CD spectrum was recorded. The solid-state spectra were presented as ellipticities (mdeg). The secondary structure contents were calculated by using the Model JWSSE-J700 program based on data deconvolution algorithmAll Far-UV CD spectra were recorded on a JASCO J-715 spectropolarimeter (JASCO) at room temperature as described previously15N/13C-labeled GB1-TDP(311–360) mutant (G335D or Q343R) was dissolved in 20 mM phosphate (pH 6.5), 50 mM NaCl and 8% D2O for NMR data acquirements. All NMR spectra were recorded at 25 °C on a Bruker Avance 600-MHz spectrometer equipped with a TCI CryoProbe (Bruker Biospin). The backbone and side-chain chemical-shift assignments were obtained from the spectra of HNCO, HNHA, HNCACB, CBCA(CO)NH, CC(CO)NH, and HCCH-TOCSY. NOE restraints for structure calculations were obtained from 15N- and 13C-edited NOESY spectra. The NMR data were processed by using NMRPipe and analyzed with SPARKY. The backbone dihedral restraints were derived from TALOS program. The structures were calculated using ARIA2.0 and CNS program, assessed by PROCHECK and displayed by MOLMOL. The structural calculation was performed for 9 cycles and a total of 200 structures were finally obtained. Ten lowest-energy structures were selected and displayed.NMR data acquirement and processing, and structure calculation were performed as reported previously2552Cell culture, transfection, Western blotting, immunocytochemistry and confocal microscopy were carried out as described previouslyThe transfected cells were lysed in 100 μL of a RIPA buffer ) on ice for 30 min and centrifuged at 15,000 g for 15 min. The supernatant was added with 100 μL of the loading buffer (2% SDS), while the pellet was sufficiently washed with the RIPA buffer for three times and then added with 50 μL of the loading buffer (4% SDS). Equal volume of supernatant and pellet fractions was subjected to SDS-PAGE with 12% acrylamide gel and transferred onto PVDF membranes (PerkinElmer). The mouse monoclonal antibodies against FLAG or Myc, goat anti-actin (Santa Cruz), goat anti-mouse IgG-HRP antibody and rabbit anti-goat IgG-HRP antibody (Jackson Immuno-Research) were used. The proteins were detected by an ECL detection kit (Thermo scientific). The band intensities were quantified by Scion Image software (Scion Corp).The method for assaying the CFTR exon 9 splicing was same as the literatures1554Accession codes: The coordinates and structure factors have been deposited in the Protein Data Bank with accession codes 2N3X (TDP(311–360)), 2N4G (G335D) and 2N4H (Q343R), respectively. How to cite this article: Jiang, L.-L. et al. Two mutations G335D and Q343R within the amyloidogenic core region of TDP-43 influence its aggregation and inclusion formation. Sci. Rep.6, 23928; doi: 10.1038/srep23928 (2016).


We have quantitatively observed a reduced but still sizable magnetic moments of the epitaxial Fe ML on graphene, which is well resembled by simulations and can be attributed to the strong hybridization between the Fe 3dz2 and the C 2pz orbitals and the sp-orbital-like behavior of the Fe 3d electrons due to the presence of graphene.Successful spin injection into graphene makes it a competitive contender in the race to become a key material for quantum computation, or the spin-operation-based data processing and sensing. Engineering ferromagnetic metal (FM)/graphene heterojunctions is one of the most promising avenues to realise it, however, their interface magnetism remains an open question up to this day. In any proposed FM/graphene spintronic devices, the best opportunity for spin transport could only be achieved where no magnetic dead layer exists at the FM/graphene interface. Here we present a comprehensive study of the epitaxial Fe/graphene interface by means of X-ray magnetic circular dichroism (XMCD) and density functional theory (DFT) calculations. The experiment has been performed using a specially designed FM As a prototypical two-dimensional quantum system, graphene displays a combination of exceptional properties including large charge carrier mobility, high thermal conductivity, strong mechanical strength, excellent optical characteristics, electrically tuneable band gap, as well as the recently discovered long spin coherence length123568910111213111/FM2/SC structure that to a large extent simulates the realistic FM/graphene interface of the proposed graphene-based transistors45In any proposed graphene-based transistors, the best opportunity for spin transport could only be achieved when no magnetic dead layer exists at the FM/graphene interface. Previous studies on various FM/semiconductor (FM/SC) heterojunctions revealed the possibility that the magnetic ordering near a region of the surface or interface of FM/SC may be modified due to interdiffusion, termination and hybridization; and controversial reports make this issue rather complex151617181920Tc)221/FM2/SC structure . The thick FM1 layer provides the FM2 ML with a source of exchange interaction and these two together restore the interfacial behavior of the thick ferromagnetic FM2/SC. Combined with the unique elemental selectivity of XMCD, such structure allows direct observation of the magnetization of FM2 at the FM2/SC interface.Fundamentally all the intriguing spintronic phenomena observed in the FM/graphene heterojunctions strongly depend upon the interfacial hybridization and magnetic exchange interaction8910111213L2,3 absorption edges were performed on beamline I10 at the Diamond Light Source, UK. Circularly polarized X-rays with 100% degree of polarization were used in normal incidence with respect to the sample plane and parallel with the applied magnetic field, as shown in − − σ+, obtained by flipping the X-ray helicity at a fixed magnetic field of 3 T, under which the sample is fully magnetized with little paramagnetic contribution. The X-ray absorption (XAS) and XMCD of the Fe and Ni L2,3 absorption edges were performed and typical spectra obtained at 5–300 K are presented in The XAS and XMCD of the Fe and Ni mspin) and orbital magnetic moment (morb) were calculated by applying the sum rulesL2,3 edges based on3d-band holes, nh, was assumed as 3.7 accounting for the charge transfer from Fe to the atop stabilizing Ni layer (nh = 3.4 in bulk)26mspin, morb, and total magnetic moment  of the interface Fe at temperatures from 5 K to 300 K were plotted in mtotal of Fe displays a decreasing trend with the increasing temperature from 1.23 μB/atom at 5 K to 1.12 μB/atom at 300 K, i.e., ~9% reduction, pointing to a TC close to bulk-like Fe. At the lowest temperature (5 K), we obtained mspin = (1.06 ± 0.1) μB/atom, which is ~50% reduced from the bulk-like Fe (whose mspin = 2.2 μB/atom) and morb = (0.18 ± 0.02) μB/atom, corresponding to an enhancement of ~200% compared to morb = 0.085 μB/atom of that in the bulkThe spin  μB/atom at 5 K to (0.52 ± 0.06) μB/atom at 300 K, pointing to a TC close to the bulk-like Nimorb shows a significant enhancement compared to that of bulk-like Fe, the Ni stabilizing layer shows a nearly quenched morb. This is expected for bulk-like Ni, where the effect of the crystal field plays a dominant role. As an additional note, the graphene is likely to carry induced magnetic moment at the Fe/graphene interface252930The calculated 2), a Fe bcc(100) 5 × 3 supercell (14.33 × 8.60 Å2), and a Fe bcc(110) 5 × 1 supercell (14.33 × 4.06 Å2), respectively, matching the graphene 6 × 2 supercell (14.82 × 8.56 Å2), a 6 × 1 supercell (14.82 × 4.28 Å2), and a 1 × 1 cell (2.47 × 2.47 Å2), are used to model Fe fcc(111), bcc(100) and bcc(110) MLs on graphene. For each superstructure, all atomic positions are relaxed until the atomic forces were smaller than 0.02 eV/Å. Figure S4 presents the initial (upper row) and the relaxed (lower row) geometries of Fe fcc(111), bcc(100) and bcc(110) stacking on graphene, respectively. The calculations suggest that the Fe prefers to follow the fcc(111)-like structure of the graphene substrate. Significant deformations were observed for both Fe bcc(110) and (100) on graphene due to the large lattice mismatch. The averaged Fe-Fe bond length after relaxation changes from 2.73 to 2.35 Å for Fe bcc(110), and from 2.91 to 2.38 Å for Fe bcc(100) ML, respectively, whilst that for Fe fcc(111) remains unchanged from 2.47 Å.In order to determine the electronic and magnetic ground state of the Fe ML on graphene, first it must be addressed where Fe likes to reside with respect to the C atoms. Periodic DFT calculations were performed to obtain the most energetically stable stacking of ML Fe on graphene. Three initial superstructures, namely, a Fe fcc(111) 1 × 1 primitive cell  μB/atom, is more likely a mixture of Fetop (mspin = 1.23 μB/atom) and Febridge (mspin = 0.67 μB/atom), given the small energy difference (∆E = 43 meV) of their calculated total energies.It was found that the Fe atoms deposited on graphene prefer to follow fcc(111) stacking and there exist three inequivalent positions of the fcc Fe allocations on graphene as presented in top/graphene (right column), respectively, together with their corresponding partial density of states (DOSs) are presented in dz2 orbitals (see the white shed areas), which is responsible for the large mspin (2.76 μB/atom). In contract, the Fe 3dz2 orbitals in ML Fetop/graphene become strongly delocalized in the reciprocal space and tend to show an sp-orbital-like partial band character. The exchange splitting sharply reduces from 2.67 eV to 0.94 eV due to the presence of graphene, pointing to a weakened exchange interaction between the 3d electrons of the Fe in Fetop/graphene and consequently a reduced spin polarization near the Fermi level (EF). Further examinations of the symmetry matching and spatial overlap suggest that the strong Fe-C hybridization predominantly originates from that between the C1 pz orbitals and the Fe 3dz2 orbitals, although the interactions between the C2 pz orbitals and the Fe 3dxz and 3dyz orbitals cannot be neglected  and the ML Fetop/graphene. The 3D charge accumulation and depletion was obtained by subtracting the charge density of Fetop/graphene from that of the freestanding Fe ML and graphene as shown in e is transferred from Fe to C atoms per unit cell. This substantial charge transfer mainly occurs at the interfacial region of the Fetop/graphene and predominately via the C1 pz orbitals. The Fe-C hybridization, in turn, slightly polarizes the C atoms in Fetop/graphene in opposite direction to the Fe moments (mspin = −0.02 μB/atom for the C1 and −0.01 μB/atom for the C2 atoms) as can be seen from bridge/graphene (not shown) are qualitatively same but quantitatively different as that of Fetop/graphene (not shown) and mspin of −0.008 μB/atom are induced for both the C1 and C2 atoms.mspin of 1.53 μB/atom for Fe. The inclusion of the topmost 7 MLs Ni has not modified the favorable stacking of the ML Fe in conjunction with the graphene. Furthermore, the calculated average mspin of the 7 ML Ni atop Fe is 0.59 μB/atom, well reproducing the experimental value of mspin = (0.59 ± 0.06) μB/atom obtained from XMCD.As an additional note, the effect of the topmost Ni was also considered. The bulk-like Ni atop the Fe ML serves as a stabilizing layer, in order to make the ML Fe in such configuration to be representative of the interface Fe of a bulk-like Fe on graphene. The two FM layers are expected to ferromagnetically couple with each other, provided an atomically clean interface is prepared, e.g., using UHV growth techniques. The calculation suggests that the stabilizing Ni layer has a limited influence on the underneath Fe by offering only a small charge transfer, i.e. 0.017 e/atom. In the fully relaxed structure, the Ni tends to pull the Fe slightly away from the graphene substrate, resulting in a higher top/graphene interface up to room temperatures. We obtained a reduced but still sizable magnetic moment of the ML Fe on graphene, i.e., (1.23 ± 0.1) μB/atom, which is well resembled by the DFT calculations. The observed suppression of the Fe magnetization is attributed to the strong hybridization between the Fe 3dz2 and the C 2pz orbitals, leading to greatly reduced exchange splitting, and the sp-orbital-like behavior of the Fe 3d electrons due to the presence of graphene. Our study restores a realistic scenario of the proposed graphene-based transistors and addresses the open question of the magnetic and electronic nature of the Fe/graphene interface. Although possibilities such as the discontinuous growth of the ML Fe and the contact of Fe with the SiO2 substrate cannot be fully excluded from the experiment, our study has demonstrated a valuable model of exploring the interfacial magnetism of FM on graphene and the results are encouraging based on the contemporary sample preparation technique. Future work to explore the tuning of the spin polarized band structure of both the FM and graphene via the interface engineering will be of great interest and have strong implications for both fundamental physics and the emerging spintronics technology.To summarize, we have demonstrated an unambiguous ferromagnetic epitaxial Fe/graphene interface by XMCD. This was aided by the inclusion of a Ni stabilizing layer atop the Fe film, which facilitated the study of the temperature dependence of the Fe/graphene interface. We observed strong dichroic XAS spectra of Fe, suggesting a well-maintained ferromagnetic order of the Fe2/Si substrate35The single layer graphene used in this study was prepared by chemical vapor deposition (CVD) on Cu foil and then transferred on top of 300 nm SiOab initio simulation package (VASP)37k-points was used to sample the two-dimensional Brillouin zone for the thin film electronic structure calculations. The total energies were obtained using the tetrahedron method with the Blöchl corrections. In each unit cell, all atomic positions are fully relaxed with the conjugate gradient procedure until the residual forces vanished within 0.02 eV/Å. The vacuum space of 15 Å was set to separate the interactions between neighboring slabs of the unit cells. Test calculations have been performed for the bulk bcc Fe (a0 = 2.87 Å) and the freestanding fcc monolayer Fe (a0 = 2.55 Å) using this method and the results compare well with the pioneering reports25293031414244The first-principle calculations were performed using the projector augmented wave methods implemented in the Vienna How to cite this article: Liu, W. Q. et al. Atomic-Scale Interfacial Magnetism in Fe/Graphene Heterojunction. Sci. Rep.5, 11911; doi: 10.1038/srep11911 (2015).


Transparent electrodes with superior flexibility and stretchability as well as good electrical and optical properties are required for applications in wearable electronics with comfort designs and high performances. Here, we present hybrid nanostructures as stretchable and transparent electrodes based on graphene and networks of metal nanowires, and investigate their optical, electrical, and mechanical properties. High electrical and optical characteristics, superb bendability , excellent stretchability , strong robustness against electrical breakdown and thermal oxidation were obtained through comprehensive study. We believe that these results suggest a substantial promise application in future electronics. Rs) of the randomly distributed mNW networks (lower than approximately 80 Ω sq−1) [Rs of undoped, chemical vapor deposition (CVD)-grown graphene (higher than approximately 1 kΩ sq−1) [−1 with transmittance of approximately 90% at 550 nm) [Transparent conductive films have been widely used in electronics such as touch screen panels, electrochromic smart windows, and displays. In addition, stretchable electronics have attracted tremendous recent attention due to their potential advantages including outstanding suitability on a nonplanar surface and excellent portability. Therefore, encoding stretchability into transparent electrodes has been actively explored as the key to emerge flexible optoelectronic devices. Indium tin oxide (ITO) is the most commonly used material in transparent electrodes. However, it has several critical disadvantages, such as (i) the increasing costs for indium, (ii) need for a complex vacuum process, and (iii) its mechanical brittleness under external stress. These weak points of ITO limit its use for flexible electronics. Thus, there is a clear and urgent need for new transparent conductive materials with superb mechanical properties. Various alternative candidates for ITO have been intensively pursued, including carbon nanotubes -6, condu Ω sq−1) ,20,26 wikΩ sq−1) ,9-11 and 550 nm) . In addi 550 nm) ,27. Seve 550 nm) ,20-22,27Rs of pristine graphene has been estimated to be as low as approximately 30 Ω sq−1 [Rs of undoped, synthesized graphene by using various methods [Rs of ITO. Although chemical doping processes [Rs of the CVD-grown graphene by increasing carrier density, the instability of the interactions between graphene and dopant typically limits the lifetime of this doping effect and leads to a time-variant increase of Rs [Another promising alternative to ITO comes in the form of graphene that can be deformed up to a strain of approximately 4% with negligible cracking  as well 0 Ω sq−1 , but the methods ,10,11,32rocesses ,33 can fse of Rs .In this paper, we report graphene-mNW hybrid nanostructures as high-performance, stretchable, and transparent electrodes. Especially, we focused on the study of the mechanical flexibility of the hybrid film, so various flexible tests such as bending, folding, and uniaxial and multi-axial directional stretching tests were performed. These fabricated hybrid structures present superb mechanical flexibility (folding with bending radius of approximately 3.7 μm) and stretchability .−1) were purchased from Nanopyxis Co., Ltd . The solution was stored at 5 °C, and then spin-coated on substrates. The spin coating condition of AgNWs was fixed at 500 rpm for 30 s, except for the density study of AgNWs, which was done according to the spin rates. After spinning the AgNW suspension on substrates (with or without graphene as a top layer), the samples were annealed at 150 °C for 90 s to evaporate the solvent completely.Silver nanowires (AgNWs) with average length of 30 μm ± 7 μm and average diameter of 20 nm ± 5 nm dispersed in ethanol (3 mg mL2 (500 sccm). The CVD growth was carried out under CH4 (12 sccm) and H2 (500 sccm) flows for 5 min, and then the reactor was rapidly cooled to room temperature under Ar (500 sccm) flow. To transfer the synthesized graphene, we used a 200-nm-thick poly(methyl methacrylate)  film as a supporting layer which was spun on the graphene. The Cu foil was dissolved in a diluted etching solution of FeCl3:HCl:H2O (1:1:20 vol.%) with the PMMA/graphene layer floating on the surface of the solution. Subsequently, the sample was rinsed by floating it on the surface of deionized water (DI water) and then the sample was transferred onto the desired substrates. In addition, the PMMA supporting layer was removed with acetone.The synthesis and transfer of graphene were carried out under the following conditions. A Cu foil  was loaded onto the center of a quartz CVD reactor under low vacuum (100 mTorr). Subsequently, the reactor was heated up to 1,000 °C with introduction of the flow of Ar (200 sccm) and HWe have employed two approaches to assemble the hybrid structures (i) AgNWs on graphene or (ii) graphene on AgNWs: (i) After the CVD synthesis of the graphene layer, the graphene was transferred onto a target substrate. After spin casting of the AgNW dispersions onto the graphene surface, this sample was annealed at 150 °C for 90 s. (ii) AgNW network films were coated on substrate, as described above. Then the synthesized graphene layer was transferred onto the films by using PMMA. Subsequently, the supporting layer was removed with acetone.3PO4:C2H4O2:C6H4NO5SNa:H2O (55:1:4:40 vol.%).The graphene-AgNW hybrid electrodes can be photolithographically patterned by using an etch-back process. First, the hybrid films were dry-etched using oxygen plasma  to remove carbon-based materials. Because the residues of the AgNWs were oxidized and disconnected by oxygen plasma exposure, the exposed areas became electrically nonconductive. Subsequently, these AgNW residues can be completely eliminated by using an etching solution of HThe preparation of a 2-μm-thick polyimide substrate for the bending test was carried out by the following procedures. After casting the 200-nm-thick PMMA sacrificial layer on a bare Si wafer, a polyimide precursor ; Aldrich, St. Louis, MO, USA) was subsequently spin-coated at 3,000 rpm for 30 s. The thermal curing was processed at 250 °C for 6 h. After fabricating graphene-AgNW hybrid electrodes on the prepared polyimide surface, this 2-μm-thick polyimide film together with the hybrid electrodes could be delaminated from the bare Si wafer by eliminating the PMMA layer using acetone.Rs using a probe station with a Keithley 4200-SCS semiconductor parametric analyzer . In addition, the two-point probe method was also used for the R measurements and electrical breakdown test using a Keithley 2425 source meter. Rs can be acquired by calculating it from the measured R by considering the aspect ratio of patterns (Rs = R × (Pattern width)/(Pattern length)). Furthermore, the change in R of hybrid films during uniaxial and multi-axial stretching was measured after fixing the sample on home-built stretching stages.For the electrical, optical, and mechanical characterization of the prepared samples, a four-point probe method was used to measure the The optical film transmittance was assessed by ultraviolet–visible near-infrared (UV–vis-NIR) spectroscopy . The transmittance of the substrates was excluded in all our data. The optical images were taken using an optical microscope . The morphologies of the fabricated films were characterized using field emission scanning electron microcopy (FE-SEM)  and atomic force microscope (AFM) .Rs than higher spin rates  , and this rc can be estimated by measuring the outer diameter of the supports. In this manner, the hybrid samples were wrapped on different cylinder-shaped objects, such as a plastic straw and a glass capillary tube. For extreme bending, the hybrid films on the substrates were folded in half by pressing two pieces of Si wafer together using tape to attach both ends of the polyimide to the wafer  [ε = (ts + tf)/2rc, where ts and tf are the thickness of the substrate and film, respectively. Although this folding leads to approximately 27% bending strain, a negligible change in R occurs for bending to radii of curvatures as small as 3.7 μm.Mechanical flexibility is another key factor of transparent electrodes for next-generation stretchable and wearable electronics. We achieved ultimate flexibility including bending and folding capability of the hybrid structures in Figure rain (ε) ,12 accorR under various tensile strains (σ) was plotted in Figure In addition to the extreme folding, the excellent elasticity of graphene  and the Rs values were measured at specific stretching strains. Figure Rs according to increase of ε, and this graphene-AgNW (spin-coated) hybrid electrode can also be stretched up to 100% tensile strain without significant change. To investigate the durability against consecutive stretching and releasing, the Rs of the sample was measured during cyclic stretching tests  and it was almost constant without notable deformation, as plotted in Figure Rs of our hybrid sample during stretching toward the multi-axial direction at up to 30% by using a homemade cylinder-shaped stretching stage, and the relative difference Rs was recorded at less than 5%  hybrid electrode. Here, a multi-axial direction means the centrifugal force direction parallel with the plane of a fabricated hybrid films. The biaxial direction stretching of film has been reported  but that2 on Si wafer, and then graphene, AgNW networks, and their hybrid structures were formed on this substrate with metal pads, respectively. The patterning of the aforementioned three materials with various widths and lengths was performed by using photolithography and dry and wet etching processes, followed by measurement and comparison of their R . Figure The electrical conductance of the nanostructured films can be changed by pattern size, such as via the widths and lengths of a pattern . For thi2). Figure I-V characteristics of graphene, AgNWs, and the hybrid structures with various channel lengths (L) were measured in an ambient condition. As one of the results, the top graph of Figure y-axis graphs are presented together at the bottom of Figure I-V curves indicate electrical conductance, and the narrow pattern size (channel length = 10 μm) caused the relatively high resistance of AgNW networks, compared to the resistance of the hybrid structures, because only small portions of the conductive pathways along the NWs remained inside the patterns. For all the three cases, currents increase almost linearly with the applied voltage before their breakdown. Based on these results, we compared breakdown bias (Vbreakdown) according to various L of these three different materials. In the case of AgNW networks that can be degraded by electromigration [Vbreakdown < approximately 4 V) for a channel length of approximately ≥30 μm (average length of AgNWs) due to the large aspect ratio of NWs and NW-NW contact resistances. On the other hand, the Vbreakdown of graphene is higher than that of AgNW networks. This polycrystalline graphene can be damaged by Joule heating and oxidation at the defect sites [Vbreakdown of graphene also decreases when the channel length falls below the threshold related to the grain sizes. In the case of graphene-AgNW hybrid nanostructures, the graphene underneath AgNWs can create a path to dissipate heat and electrical stress, which can make AgNWs sustainable against breakdown at relatively high electric bias. Therefore, the hybrid structures exhibit the highest values of Vbreakdown among the three materials . Also, the hybrid electrode can flow the highest current against its breakdown due to its low resistance. The hybridization of the graphene and AgNWs give the hybrid structures strong endurance against electrical load, so this improved robustness is another advantage of the hybrid electrode.In addition, the robustness of conductive materials that can withstand electrical load becomes an important feature for next-generation electronics, as devices are highly integrated and miniaturized. This is because the electric field is inversely proportional to the length of pattern. To assess and compare the robustness of nanostructured films against electrical loads, patterned samples of graphene, AgNWs, and their hybrid structures as channels with a width of 10 and 20 μm were prepared on a metal electrode deposited onto a Si wafer , 10,000 times in stretching cycles with 100% tensile strain toward the uniaxial direction, and 30% stretching in tensile strain toward the multi-axial direction. We believe that the hybridization with two-dimensional and one-dimensional nanomaterials presents a promising strategy toward stretchable, wearable electronics and implantable biosensor devices.The results presented here demonstrate the superior characteristics of graphene-AgNW hybrid nanostructures as flexible, stretchable, and transparent electrodes. From the fundamental study, we confirmed that the hybridization between AgNWs and CVD synthesized graphene could significantly enhance electrical properties such as low


We obtained a profound light modulation over a broad spectrum by using the fabricated large area transparent graphene supercapacitors and we compared the performance of our devices with the one based on graphene from standard copper. We propose that the copper foils used in the lithium-ion batteries could be used to obtain high-quality graphene at much lower-cost, with the improved performance of electrical transport and optical properties in the devices made from them.This work demonstrates an attractive low-cost route to obtain large area and high-quality graphene films by using the ultra-smooth copper foils which are typically used as the negative electrodes in lithium-ion batteries. We first compared the electronic transport properties of our new graphene film with the one synthesized by using commonly used standard copper foils in chemical vapor deposition (CVD). We observed a stark improvement in the electrical performance of the transistors realized on our graphene films. To study the optical properties on large area, we transferred CVD based graphene to transparent flexible substrates using hot lamination method and performed large area optical scanning. We demonstrate the promise of our high quality graphene films for large areas with ~400 cm Carbon based nanomaterials have provided a new perspective in electronics owing to their very high charge carrier mobility and nanoscale dimensions. In particular, the carbon nanotubes7Graphene layers can be isolated chemically from bulk graphite with intercalation of atoms or molecules through the graphite201319et al.et al.26et al.The surface quality of the Cu foils and the transfer process greatly influence the performance and reliability of graphene based devices. Previous works have showed that the morphology of the Cu surface is one of the key parameters to have better quality graphene films2 flexible graphene electrodes and enlarged 160 times the previously reported (2.5 cm2) supercapacitor structureIn this paper, we demonstrate that the commercially available ultra-smooth Cu foils  can help to open new avenues for large area device applications with the improved quality of the graphene films. We studied the synthesis of graphene on commercially available smooth Cu foils and we provided a detailed comparison between the graphene synthesized on ultra-smooth Cu foils and the one synthesized on the standard Cu foils in terms of structural, electrical and optical properties. We observed a stark improvement in the electrical performance of the transistors realized on our graphene films. Then, to demonstrate the advantage of high quality graphene films from ultra-smooth Cu foils, we fabricated 400 cmAlfa Aesar Cu foil which is commonly used for CVD of graphene    in the s B1-SBS)  and scan B1-SBS) . On the  B1-SBS) . We inve2 in the form of a snow-flake  on flexible, 125 μm thick PVC substrate. After the CVD synthesis on ultra-smooth Cu foils, we transferred the graphene to the flexible PVC substrates by using hot lamination method. During the Cu etching, we protected the Cu foils at the edges of the sample to use them as contact electrodes for resistivity measurements and application of bias voltage. The measured total resistance of graphene is superposition of contact resistance and the sheet resistance terms  above the half of the incoming photon energy. This blocks the inter-band transitions and makes the graphene more transparent by blocking optical absorption mechanism of the material. By blocking the inter-band transitions, we achieved the light modulation over a broad range of wavelengths with the large area graphene supercapacitors. We achieve more efficient transmittance modulation in the case of smooth-Cu-graphene based supercapacitors due to the lower sheet resistance and high coverage of the smooth-Cu-graphene film. F. Therefore we were able to extract the Fermi energy values corresponding to each curve. Previously we showed that the supercapacitor device structure can provide an efficient electrostatic doping of graphene which cannot be achieved by graphene based transistors using solid insulating dielectric layersHere the limiting factor for the bias voltage is the electrochemical window of the electrolyte used in the supercapacitor structure. Since our devices use mutual gating between top and bottom graphene electrodes, they are limited with the two times of the electrochemical window of the electrolyte (1-Butyl-3-methylimidazolium hexafluorophosphate) which is around 4.5 V. In mutual gating; the Dirac points of the top and bottom graphene shifts with respect to each other leading to two different dips in the capacitance modulation. As a result, we obtained a more efficient transmittance modulation up to two times (at 700 nm) by using smooth-Cu-graphene in supercapacitors operating as optical modulators. We also observed 27% more efficient Fermi level shift compared to rough-Cu-foils (at 4.4 V). By using electrolytes with broader electrochemical windows, the enhancement in transmittance and Fermi energy shift by using smooth-Cu-graphene can be further demonstrated in optoelectronic devices.2) that we used in this study makes the costs affordable for large scale production of graphene. Cu surface modification techniques also add to the overall cost to the graphene growth with the standard Cu foils which itself costs around 115 $ per m2.The lower price of commercially available ultra-smooth Cu foils  in CVD synthesis holds great promise for high performance in large area graphene based devices. Recent industrial progress of large scale production of smooth Cu foils, for application such as lithium-ion batteries, can also provide opportunities to improve the quality of graphene grown by CVD. After a detailed structural comparison between the rough and smooth-Cu-graphene, we studied the performance improvements in the transistors based on smooth-Cu-graphene by comparing them to transistors based on rough-Cu-graphene. To demonstrate the promises of the smooth-Cu-graphene on large area, we fabricated flexible graphene electrodes and mapped the transmittance by optical scan. These electrodes enabled us to fabricate graphene-based flexible supercapacitors operating as voltage controlled optical modulators. By enlarging the previously demonstrated supercapacitor structure, we examined the effects of usage of rough and smooth-Cu-graphene in the large area electro-optic devices. As against previous works relating the Cu surface morphology to resulting graphene quality, our method does not require chemical treatment or deposition of Cu foils. The effectiveness of our approach is two-fold:We believe that large scale and low cost synthesis of high quality graphene films and the compatibility of our method to the roll-to-roll fabrication would open an avenue through the realization of graphene based flexible optoelectronic systems such as cell phones with roll-up displays, e-paper, radio-frequency identification (RFID) tags, as well as medical patches that can be attached to the skin to deliver drugs or monitor vital signs, and tactile or electronic skin for robotics and prosthetics3637444 and H2 gases are 285 mTorr and 40 mTorr and corresponding rate of flows are 25 sccm and 4 sccm respectively). We annealed the Cu substrates for half an hour at 1035 °C before the growth. Hydrogen flow was kept constant during the heating, annealing, growth and cooling. After 10 minutes of growth, samples were left for fast cooling to room temperature. The growth took place in a 8 inch diameter quartz chamber. We used both the commercially available 20 μm thick smooth Cu foils  and the standard 25 μm Alfa Aesar Cu foil (item #13382) as growth substrates in the same CVD chamber.For single layer graphene, at 1035 °C methane was sent into growth chamber in company with hydrogen  drop casting method. We drop casted thick droplets of S1813 photoresist on graphene holding Cu surfaces and waited overnight at 70 °C to gently harden the PR. Then hardened PR on graphene holding Cu was put into the dense FeCl3 solution yields the graphene on flexible PVC substrates.We also used lamination technique to transfer large area graphene on PVC. At 100 °C PVC becomes active and exhibits conformal sticking to the graphene holding side of ultra-smooth Cu foils when laminated through the hot rollers. Etching the Cu foils in FeClWITec Alpha 300 M Raman microscope system with 50x microscope objectives was used for Raman mapping. The excitation wavelength was 532 nm.In order to fabricate the graphene based FETs, first conventional UV-photolithographic techniques were used to shape the source and drain electrodes on graphene. Then, thermal evaporation technique was used for metallization of the electrodes. The source and drain electrodes are 50 nm gold on 5 nm titanium adhesive layer. After lift-off step, we isolated the graphene layer by using reactive ion etching so that the graphene layers is aligned between the source and drain electrodes to define the channels. Transfer and output characteristics of the graphene based transistors were performed with HP4142B semiconductor parameter analyser. We recorded the drain current as a function of gate voltage at constant drain voltage. We repeat the measurements for devices having various channel lengths.The transmittance measurements were performed by using Bruker Vertex 70 V FTIR spectrometer integrated with silicon photodetector working in visible and near-IR. We worked in the wavelength range between 400 nm to 1100 nm. We biased the graphene based transistors using Keithley 2400 source measure unit during the transmittance measurements. The leakage current is also recorded during the measurement.How to cite this article: Polat, E. O. et al. Synthesis of Large Area Graphene for High Performance in Flexible Optoelectronic Devices. Sci. Rep.5, 16744; doi: 10.1038/srep16744 (2015).


Controlling tunneling properties through graphene vertical heterostructures provides advantages in achieving large conductance modulation which has been known as limitation in lateral graphene device structures. Despite of intensive research on graphene vertical heterosturctures for recent years, the potential of spintronics based on graphene vertical heterostructures remains relatively unexplored. Here, we present an analytical device model for graphene-based spintronics by using ferromagnetic graphene in vertical heterostructures. We consider a normal or ferroelectric insulator as a tunneling layer. The device concept yields a way of controlling spin transport through the vertical heterostructures, resulting in gate-tunable spin-switching phenomena. Also, we revealed that a ‘giant’ resistance emerges through a ferroelectric insulating layer owing to the anti-parallel configuration of ferromagnetic graphene layers by means of electric fields via gate and bias voltages. Our findings discover the prospect of manipulating the spin transport properties in vertical heterostructures without use of magnetic fields. Graphene, a honeycomb-like single layer crystal of carbon atoms, has been attracting a lot of attention in the recent decade both in terms of fundamental interests and technology. Amongst the various aspects of graphene, one of the most promising potentials is that it has extraordinary transport properties such as high carrier mobility and long mean free path124578Recently, an alternative platform has emerged for graphene FETs ,where graphene and other two dimensional layers are stacked vertically1010101015171319222324In this article, we investigate the spin-resolved transport through the vertical heterostructures with ferromagnetic graphene (FMG). We show that the control of the spin transport through the structure can be achieved by electrically manipulating the spin configurations in FMG sheets. The spin-resolved band structure is taken into account to describe the electronic states of FMG, and the spin-resolved tunneling current density is calculated for two different combinations of heterostructures: FMG–normal insulator(NI)–FMG and FMG–ferroelectric insulator (FEI)–FMG. We also show that the giant electroresistance emerges for the anti-parallel configuration of FMGs when the sandwiched insulator is replaced by an FEI.Vb can be applied between two graphene sheets, yelding the tunneling current through the insulating layer. The electronic properties of FMG are characterized by its spin-resolved electronic statesThe system studied in this article is a vertically stacked heterostructure which is formed by FMG and an insulating layer see . The san27282931σ = ±1 for spin-up and down states of Dirac fermions, v↑ = 1.15 × vF and v↓ = 1.4 × vF are Fermi velocities for each spin with vF = 106 m/s, Δ↑ = 134 meV and Δ↓ = 98 meV are the spin-resolved bandgaps, as displayed in where As a starting point, we introduce our vertical transport model used in this study. We consider the elastic tunneling of Dirac fermions in terms of energy, and scattering effects are taken into account by applying the current density formula. Also, we assume that the magnetizations of two FMG layers are parallel in the absence of external magnetic fields. The spin-resolved vertical tunnelig current is formulated with the interlayer transition matrix element based on WKB approximationwhere ties see , and  anm* is the effective mass of the tunnel barrier material, Δ is the barrier height of the tunneling insulator, Vb is bias voltage which is applied via two graphene sheets, and d is the thickness of the tunneling insulator. Note that Γ is an energy-independent prefactor which represents the momentum scattering of Dirac fermions by disorders such as defects or phonons inside the tunnel barrier material. In other words, Γ = 1 means no scattering mechanism while Dirac fermions tunnel through the tunnel barrier, and on the other hand, the smaller Γs indicate more diffusive vertical transport through the tunnel barrier.Here, VTG = VBG ≡ VG. This assumption allows us to simplify the problem with fixed chemical potentials on both graphene layers in equilibrium, i.e., n0 is proportional to the gate voltage VG, i.e., n0 = αVG where α is the proportional constant depending on the substrate (superstrate) materials between a graphene layer and the bottom (top) gate electrode. When bias voltage is applied to both graphene layers, their chemical potentials are shifted and equilibrium is broken, resulting in non-zero tunneling current denslty throughout the vertical heterostructure.Carrier density on graphene layers is controlled by field-effects via gate electrodes. In the absence of bias voltage, the chemical potentials on both graphene layers are in equilibrium, leading to no net tunneling current density. The dual-gated platform is considered to fix and maintain the same carrier densities in the top and the bottom gate electrodesm* = 0.5 mel with the bare mass of an electron mel are used for the calculations, which are approximately compatible with typical 2d materials such as MoS2, WS2, etc.333435Vb| < 0.1 V. Remarkably, within this bias voltage range, it is found that the current density can be fully spin-polarized according to gate voltage. This gate-tunable feature of the spin transport is led by the following mechanisms. When the equilibrium chemical potential μ0 is place in the mid-gap, the small bias voltage cannot lead to a sufficient amount of tunneling current densities for both spins. As bias voltage increases, spin-resolved current densities begin to flow. The spin-down current starts flowing slightly earlier than the spin-up current density because of their different electronic properties, i.e., the amount of band gaps and the position of the band edges. For large bias voltages, the current densities for both spins keep increasing with different increasing ratios associated to the spin-resolved Fermi velocities in SDOS. For VG = 15 V, as plotted in VG = −8 V makes an FMG layer purely spin-up polarized, and the contribution to the tunneling current density is dominated by spin-up Dirac fermions for small bias voltages as shown in VG = 0 V because the majority spin states near both band edges are opposite to each other . For typical ferromagnetic metal (FM)-NI-FM heterojunctions, electrical resistance strongly depends on how the FM configuration is set. While the electrical current flows well with the small resistance in the parallel configuration, a very large resistance is measured in the anti-parallel configuration. To achieve GMR, devices should be asymmetrically fabricated by using different kinds of FM materials, for which magnetization varies with respect to external magnetic fields. This means that controlling the magnetic fields is essential to change FM configuration. In addition, GER has already been introduced in a normal metal (NM)-FEI-NM heterojunction by using the asymmetric electrical response of a sandwiched FEIThe properties of an FEI are described by a simple model of the polarization density as a function of an external electric field,p0 is the saturated polarization density, β is the characteristic coefficient with the physical dimension of inverse electric fields, d, and Ec is the coercive field which is responsible for the hysterisis of the FEI. Here, the factor s = ±1 implies how the electric field varies, i.e. the forward or reverse sweep of electric fields  is the sign function. In equilibrium, the chemical potentials on the FMG layers should be arranged at the same Fermi energy to be consistent with equilibrium in the absence of bias voltage. Therefore, the Dirac cone on each FMG layer is shifted by δμ/2 = (μT − μB)/2, respectively, i.e., δμ/ed. This FEI-induced electric field is reflected in the tunnling probability as below,The shift of the spin-resolved Dirac cones is led by the following mechanism. For dual-gated devices, the carrier densities on FMG layers VG = 0 V where chemical potentials reside near the band edges of FMGs. Further, for the very small bias Vb = 0.01 V, the tunneling current is allowed only by a positive bias, whereas it is strongly suppressed by a negative bias. Such a large GER ratio is led by the following mechanism. For very small bias voltages, the polarization density is almost unchanged from the saturated value, and the resulting Dirac cone shift makes one FMG layer purely spin-up polarized and the other FMG layer purely spin-down polarized, i.e., anti-parallel spin configuration of FMGs is derived. When a small bias is applied in the positive direction, the chemical potential on spin-up polarized FMG becomes lower and touches the lower spin-down band, while the chemical potential on spin-down polarized FMG becomes higher but still reside in the spin-down band only. On the other hand, the FMG configuration remains anti-parallel for Vb = −0.01 V, resulting in the suppression of vertical tunneling by Pauli blocking. Indeed, Vb = 0.01 V is influenced by the spin-down states only. Therefore, the GER ratio ~104 originates from the bias-tunable spin-configuration of FMGs. Also, the GER ratio has dependence on temperature and deteriorates at higher temperatures .


Long Interspersed Nuclear Element-1 (L1) is an oncogenic mammalian retroelement silenced early in development via tightly controlled epigenetic mechanisms. We have previously shown that the regulatory region of human and murine L1s interact with retinoblastoma (RB) proteins to effect retroelement silencing. The present studies were conducted to identify the corepressor complex responsible for RB-mediated silencing of L1.Chromatin immunoprecipitation and silencing RNA technology were used to identify the repressor complex that silences L1 in human and murine cells.Components of the Nucleosomal and Remodeling Deacetylase (NuRD) multiprotein complex specifically enriched the L1 5′-untranslated DNA sequence in human and murine cells. Genetic ablation of RB proteins in murine cells destabilized interactions within the NuRD macromolecular complex and mediated nuclear rearrangement of Mi2-β, an ATP-dependent helicase subunit with nucleosome remodeling activity. Depletion of Mi2-β, RbAP46 and HDAC2 reduced the repressor activity of the NuRD complex and reactivated a synthetic L1 reporter in human cells. Epigenetic reactivation of L1 in RB-null cells by DNA damage was markedly enhanced compared to wild type cells.RB proteins stabilize interactions of the NuRD corepressor complex within the L1 promoter to effect L1 silencing. L1 retroelements may serve as a scaffold on which RB builds heterochromatic regions that regulate chromatin function. Human L1s are non-LTR mammalian retrotransposons that consist of an internal bidirectional promoter, two open reading frames encoding for ORF1p (RNA-binding protein) and ORF2p (reverse transcriptase and endonuclease activities), a 3′- untranslated region (UTR), and a polyA tail [We have previously shown that recruitment of E2F/RB (E2F-retinoblastoma) to the L1 promoter, along with histone deacetylases 1 and 2  , 5 and mMCF7 cells, HeLa cells, primary mouse embryo fibroblasts (MEFs) and RB null MEFs lacking all three family members  (TKO) were cultured in Dulbecco’s Modified Eagle’s medium  supplemented with 10 % fetal calf serum (FCS), 200 mg/ml streptomycin, and 200 U/ml penicillin G at 5 % CO2 and humidified air at 37 °C.Total RNA was extracted and 1 μg used for cDNA synthesis (Invitrogen Superscript II). For each reaction, 25 μL of 2X SYBR green (Biorad) was mixed with forward and reverse primers to give a final concentration of 10 μm. One μL of cDNA mixture was brought up to 50 μL using DEPC water. Cycling conditions included an initial denaturation at 95 °C for 3 min followed by 50 cycles at 95 °C for 30 s, 55 °C for 30 s and 72 °C for 45 s. The homogeneity of PCR products was confirmed using a real time PCR melting curve.Cells were grown in 150 mm petri dishes to 90 % confluence, fixed in 1 % formaldehyde for 7 min and rinsed 3X with cold PBS. Chromatin was sheared to fragments below 1 Kb and 20 μg immunoprecipitated using selected antibodies. After purification, 0.5 to 2.0 μL of DNA were used for PCR. DNA from non-precipitated chromatin samples was extracted and used as input for PCR. Controls included isotype-matched IgG or Mock reactions which included all reagents except chromatin.Cells were grown on glass coverslips to 30 % confluence, rinsed 2X with PBS , fixed in 3 % paraformaldehyde for 15 min at 4 °C and rinsed twice with PBS. Fixed cells were permeabilized with 0.1 % Triton-X100 for 5 min, blocked in 2 % dry milk for 15 min, rinsed with PBS twice and incubated with primary antibody  in PBS containing 1 % dry milk and 0.1 % Triton-X100 overnight. Secondary antibody, mouse mAB-Alexa-488 conjugate, was incubated for 1 h at room temperature. Cells were washed 3X with PBS and incubated in 1 μg/ml Hoescht nuclear dye for 20 min and then washed twice. Cells were visualized in a CARL ZEISS AXIOVERT 200 inverted microscope at a magnification of 63X.Cells were lysed with buffer and centrifuged at 4000 rpm for 2 min. Immunoprecipitation using anti-Mi-2β and anti-RbAp46 were completed as described previously [Cells were transfected with an EGFP-tagged L1RP vector. After the disappearance of EGFP fluorescence by epigenetic silencing, cells were replated and treated at 30-50 % confluence with 4 nM siRNA for 48 h followed by direct measurements of fluorescence.Cells were challenged with 3 μM BaP or 0.06 % dimethyl sulfoxide (DMSO) as vehicle control for 16 h to reactivate L1.No ethics approval was required for the experimental work performed in this study.RP 5′-UTR was compared to Snail and glucose 6-phosphate dehydrogenase (G6PD) gene promoters as positive and negative control sequences, respectively. All NuRD subunits specifically enriched the sample for the human L1RP 5′-UTR DNA sequence, but not the G6PD sequence , RB-associated proteins 46 and 48 (RbAp46 and RbAp48), MBD2 and MBD3, and metastasis-associated proteins 2 and 3 (MTA2 and MTA3) –11. Valince Fig. . The preSubsequent experiments were conducted to examine the role of RB proteins in NuRD complex recruitment comparing wild type MEFs to TKO MEFs in which all three member of the RB family were genetically ablated . In thesSince basal L1 mRNA levels are markedly upregulated in TKO MEFs , we sougRP carrying a chimeric ORF1p-GFP fusion protein. The backbone vector, pGL4.13, lacked an episomal sequence and is known to undergo silencing upon integration, as evidenced by the loss of fluorescence pyrene (BaP), a genotoxic carcinogen that reactivates L1 via epigenetic mechanisms . BaP indThe NuRD complex represents a multiprotein complex that mediates formation of heterochromatin and localized domains of repressive chromatin. Its repressor activity involves assembly of a histone deacetylase macromolecular complex that couples histone modifications with nucleosome-stimulated ATPase activity. Mi2-β, RbAp46, MBD3, MTA 2 and MTA3 have been identified as critical NuRD multiprotein complex components. Mi2-β functions as an ATP-dependent helicase , RbAp46 The epigenetic silencing profile executed by the NuRD complex includes histone deacetylation, histone H3 lysine K9, H3 lysine 27, and H4 lysine 20 trimethylation, and DNA methylation, which together contribute to the recruitment of non-coding RNAs and repressors and corepressors to induce facultative or constitutive heterochromatin formation . Genes sRepetitive L1 sequences in the mammalian genome are extensively silenced via DNA methylation, and L1 methylation status is frequently used as a proxy of global cellular methylation , 25. ThiOur previous findings are consistent with the new data showing that L1 silencing by RB proteins is effected through the NuRD multiprotein complex. All major components of the NuRD complex are recruited to the silenced L1 promoter, and recruitment exhibits promoter-specificity. RB appears to orchestrate the recruitment of NuRD proteins and loss of RB function is associated with delocalization of Mi2-β from the nuclear to the cytosolic compartment. A critical role for RB in regulation of L1 retrotransposon was established in studies showing that genetic ablation of RB family members is associated with highly unregulated overexpression of L1 in MEFs. Recent studies using non-transformed human bronchial epithelial cells have shown that L1 reactivation by BaP is associated with disruption of NuRD complex assembly and function . Similar studies using cancer cell lines are confounded by differences in RB mutant status and RB-related functions. Such differences notwithstanding, comparative analyses of constitutive L1 expression have shown that human cells exhibit remarkably lower expression than murine cells , a findiCollectively, our findings identify RB as a critical regulator of the NuRD multiprotein repressor complex assembly within the regulatory region of L1. NuRD recruitment is in turn responsible for effective silencing of L1 transcriptional activity. Genetic ablation of RB proteins destabilizes protein-protein interactions within the NuRD macromolecular complex to mediate nuclear rearrangement and delocalization of Mi2-β from the nuclear to cytoplasmic compartment. These findings suggest that Mi2-β is critical to transcriptional repression of L1. These findings are highly relevant to our understanding of oncogenesis and malignant progression given that NuRD is a key determinant of cellular differentiation, and that inappropriate recruitment of NuRD to specific loci contributes to tumorigenesis . Given tOur findings indicate that the L1 silencing by RB is effected through the NuRD macromolecular complex and implicate these molecular interactions in the formation of L1 heterochromatin structures. Full-length L1s in the mammalian genome are therefore potential targets for NuRD-mediated silencing and recruitment of silencing proteins such as the polycomb family of proteins. As such, unregulated silencing of L1 may play a central role in the initiation and progression of cancer.


Using nano angle resolved photoemission (nanoARPES) complemented by transport studies and Raman spectroscopy as well as density functional theory (DFT) calculations, we report the direct observation of massless Dirac particles in monolayer graphene, providing a comprehensive mapping of their low-hole doped Dirac electron bands. The micrometric graphene flakes are oriented along two predominant directions rotated by 30° with respect to each other. The growth mode is attributed to the mechanism when small graphene “molecules” nucleate on the Ge(001) surface and it is found that hydrogen plays a significant role in this process.The practical difficulties to use graphene in microelectronics and optoelectronics is that the available methods to grow graphene are not easily integrated in the mainstream technologies. A growth method that could overcome at least some of these problems is chemical vapour deposition (CVD) of graphene directly on semiconducting (Si or Ge) substrates. Here we report on the comparison of the CVD and molecular beam epitaxy (MBE) growth of graphene on the technologically relevant Ge(001)/Si(001) substrate from ethene (C Graphene is widely supposed to be a material on the way to enable new development of many modern technologies13436791112131516171819202122232421μm227Graphene films grown on Ge may be used in two ways. First, if large area graphene (like 300–450 mm in diameter) is obtained, the film may be transferred to the target wafer at any point of the fabrication process  without the hazard that the production tools become contaminated. In practice this means that the growth must be performed on Ge(001)/Si(001) wafers, because Ge wafers of this size are not commercially available and it would most probably be prohibitively expensive to develop and introduce a Ge wafer growth technology solely for this purpose. The second way to use graphene films grown on Ge may be open when graphene can be grown directly on device-size Ge(001) islands pre-deposited on the Si(001) substrate. For example, the width of the active region of a graphene base transistor may be smaller than a micron, and its whole area may be of the order of 4-based CVD2H4 precursor with the Ge(001) surface is then analyzed in the framework of ab initio density functional theory (DFT) calculations and conclusions for the microscopic mechanisms of the growth process are drawn. These mechanisms are then compared to those expected for the growth of graphene on Ge(001) from atomic carbon.Graphene growth on the technologically relevant Ge(001) and Ge(001)/Si(001) wafers has been reported for a CHEF in the CVD sample vanishes there (see discussion below).Morphology of the CVD- and MBE-grown graphene on Ge/Si(001) was studied by means of scanning electron microscopy (SEM) . The graEF the Fermi level (EF) of the CVD-produced graphene flakes on Ge(001)/Si(001) substrate was studied in the micro and nanoARPES experiments with 120 μm and 100 nm spot size resolution, respectively2930ED) coincides with EF. In the vicinity of ED, the energy of graphene electronic states depends linearly on the wave-vector k . Here, EB is the binding energy and kx, ky are two in-plane components of the wave vector of electron. When a neutral graphene has two domains rotated with respect to each other by 30° and the same area of each of them is illuminated by the photon beam, the Fermi surface in the ARPES experiment consists of 12 equivalent spots centred at the 12 equivalent K-points around the BZ centre, Γ. When a neutral graphene is fully disordered or polycrystalline, the Fermi surface forms a circle centred at Γ and similar to the one observed for highly oriented pyrolytic graphite (HOPG).The electronic structure in the vicinity of vector k . These evector k . In caseI measured for graphene/Ge(001)/Si(001) at hν = 100 eV (spot size 120 μm) as a function of two emission angles, polar θ (detector angle) and azimuth ϕ (in-plane sample rotation), and then converted into I. One can clearly see that the photoemission at the Fermi level has a circle-like shape centred at Γ and having equidistant intensity maxima. The radius of the circle in the reciprocal space is 1.7 Å−1 and the angular distance between the maxima is 30°. This indicates that the observed photoemission intensity originates from graphene flakes with orientation distribution that has probability maxima in the directions corresponding to graphene hexagons being parallel and/or perpendicular to the surface dimer rows of the Ge(001) surface (see discussion below).hν = 30 eV along the Γ − K direction for one of the preferential domains is given in A detailed map of photoemission intensity measured with circularly polarized light of EF is located by 0.185 eV below ED, which corresponds to the hole concentration of 2.3 × 1012 cm−2. The linear fit of the Dirac cone in the ARPES intensity map in 6 m/s. This can be compared with the graphene/Ge/SiC(0001) system, where p-doping of graphene was obtained via intercalation of Ge atoms in graphene/SiC at 720 °C and the hole doping of 4.1 × 1012 cm−2 was measured−1.These data allow us to estimate the energy distance between the Fermi level and the Dirac point, and hence the doping level of graphene: s and Ge 3d core levels are marked. The photoemission intensity at the K-point for one of the preferential graphene domains and in the vicinity of the Fermi level , which in turn is determined by the doping level; therefore VFB measures the doping level. Comparison of the VFB value for CVD and MBE prepared films indicates that former films are more strongly p-doped. The measurements in the FET configuration allow to estimate the carrier mobility of the graphene channel. For this purpose, one assumes that the film area controlled by the gate tip corresponds to a square and one measures the sample-tip distance d by approaching the tip towards the sample until mechanical contact is observed. The measured carrier mobilities were lower than for high-quality graphene and amounted to 1000 ± 500 cm2/V ⋅ s at room temperature for CVD and MBE prepared graphene/Ge(001). 2/V ⋅ s are reportedResults of a electrical measurements in the field-effect transistor (FET) configuration in 4-point STM  are consents cf. : the Fer samples . The gatet al.−112 cm−2 for the MBE graphene and around 3 × 1013 cm−2 for the CVD sample. The presence of data points with G-mode positions on the left side of the zero-doping line can be partially explained with the anomalous phonon softening of the G mode for Fermi-level shifts close to the Dirac point353712 cm−2.In accordance with the electrical 4-point STM measurements, the doping level of the CVD graphene film is higher than that of the MBE-grown film as deduced from the Raman spectra measurements. A rough estimate yields the average doping below 1 × 102), whereby a mobile Ge adatom is released. The process is reversible: a Ge adatom may eject the C atom from a C-Ge surface dimer. The mobile and highly reactive C interstitially produced in this process can diffuse under the surface. The mobile species controlling the MBE growth mode of graphene on Ge(001) are C atoms, short chains consisting of few C atoms, and Ge adatoms; the latter are attracted to the edge of growing graphene. Carbon dimers (C2) are trapped by Ge dimer vacancies (DV) to become graphene seeds; a similar role is attributed to longer C chains (like C8) trapped between dimer rows. In contrast to that, results of our calculations suggest that nucleation of CVD graphene on a flat Ge(001) surface takes place predominantly on the DV defects (see Although carbon and Ge form no stable carbide phase (GeC), the chemical interaction between the Ge(001) surface and C atoms is strong. This is particularly pronounced when atomic carbon is deposited2H4 precursor molecules do not give all H atoms away to the substrate. Calculations indicate that the major unpolymerized diffusing species may be, apart from adsorbed H, C2H3 molecules, with small admixture of C2H and C2 (see 2H2 (acetylene) produced during C2H4 decomposition desorb before they decompose further. The surface reactions do not lead to significant production of Ge adatoms. Moreover, although there are Ge adatoms on the surface (due to high temperatures used in these experiments), they are repelled from the edge of a hydrogen-terminated graphene. This is in strong contrast to the MBE process: adsorption of Ge adatoms at edge sites of hydrogen-free graphene is energetically preferred to adsorption at surface steps.The vast majority of Csp3-like bond DV, trapped inside Ge dimer vacancy defects DV is produced also by atomic C in two kick-out eventsGe + Ge(bulk) −0.13 eV, and C(diff) + CGe → (C2)DV + Ge(bulk) + 0.31 eV, where C(diff) is the mobile atomic carbon diffusing under the surface and Ge(bulk) is a Ge atom having the same energy as in the Ge bulk. The total energy gain of 0.2 eV is in this case much smaller than for trapping on a DV (1.0 eV), because a new DV site must be produced, by subsequent ejection of two Ge atoms. The difference between these two energies (0.8 eV) should thus be equal to the DV formation energy (0.6 eV). The discrepancy (0.2 eV) comes from inaccuracies in the calculation of atomic relaxations, caused by the presence of nearly degenerate metastable configurations of defects and by finite size effects. This inaccuracy amounts to at least 0.1 eV per defect.] Dimer vacancies collect carbon; e. g., when the second C2 is trapped by a DV already occupied by C2, the energy decreases by 0.9 eV, which is similar to the energy gained by trapping of the first C2. Trapping of carbon by dimer vacancies provides a means to align the orientation of graphene with that of the substrate  dimers by GeC bonds. But if a polyarene has grown large enough to lean towards the horizontal position as the result of van der Waals attraction, the GeC bonds become weakened and are now vulnerable to attack by H atoms. The latter leave the Ge(001) dimers and stick to C. We estimate that, at the high temperatures needed for the CVD growth, even when the H originates only from C2H4 decomposition and when the partial pressure of C2H4 is only 2 × 10−2 mbar (as in our experiments), the whole edge of such a polyarene becomes saturated with H  to ethylene, DVs have no influence: at the growth conditions, they do not trap Cd with H  and the periment .2H molecules. First, the process of C4H2 formation is significantly more favorable inside a DV (4H2 (DV) traps C4H4; the resulting C8H6 molecule is likely to remain inside the dimer vacancy /Si(001). According to the mechanism proposed above, the domain size is associated with the abundance of Ge dimer vacancies (on which graphene nucleates), and with the presence of surface steps on the substrate surface . When the surface is well oriented and flat , the average step-step distance (or terrace width) equals to as is the surface lattice constant of Ge(001). This is in accordance with the observation that both orientations occur with the same frequency on the area covered by the X-ray spot. The quality of graphene increases with the substrate temperature because the concentration of dimer vacancies NDV  ≈ 0.6 eV) increases with temperature slower than is the decrease of the fraction pocc of DVs occupied with hydrocarbons x. For example at thermal equilibrium, pocc is thermally activated with Etrap(C4H2) ≈ 1.7 eV. In other words, the concentration s~poccNDV of the growing nuclei decreases with temperature.ARPES measurements strongly indicate that the orientational domains are so small that many of them fit into the area sampled by the X-ray beam spot, i. e. into space of about 100 × 100 nmμm range.Small domain size is responsible also for the curved shape of CVD graphene edges . The shaDoping of the MBE film may originate solely from molecules from the ambient air that are adsorbed on top of the graphene. Higher doping of the CVD film may be caused by Ge atoms−6 K−1) and Ge (6 × 10−6 K−1). Cooling from 930 °C to 25 °C would produce strain of about −0.8%, which by its absolute value is higher than the −0.3% estimated under the assumption of biaxial strain (The compressive strain in the MBE film (−0.3%) may possibly be caused by thermal expansion coefficient difference between graphene  technology compatible Ge(001)/Si(001) substrates from CThe major experimental result is that the graphene film on Ge(001)/Si(001) consists of two orientational domains rotated by 30° with respect to each other, and that the grain size is comparable to step-step separation of the substrate. Poor crystallographic quality (high Raman D mode peak) and poor electrical quality (low carrier mobility) of the CVD graphene produced in this process are attributed to the small domain size. Arguments linking the grain size with the size of surface terraces separated by Ge(001) monatomic steps are presented. It was argued that the dimer vacancies omnipresent on the surface act as seeds on which graphene can nucleate and that the orientation of a domain is determined by the orientation of Ge dimers on the surface terrace in which the domain nucleated. Small graphene nuclei released from the dimer vacancy sites by hydrogen atoms are volatile; this contributes to the reduction of graphene growth rate by increased surface concentration of hydrogen and to the increase of the average domain size. For completeness and illustration, the CVD growth mode and mechanism was compared to those for the growth from atomic carbon in MBE chamber.−10 mbar range. The substrate was a Ge(001) film (2 μm) grown by CVD on a Si(001) wafer and then overgrown in situ by epitaxial Ge buffer (about 4 nm). The substrate temperature during graphene growth was close to but below the melting point of bulk Ge (it was around 930 °C). Growth time was varied between 90 min and 200 min in both cases. The substrate was heated from the backside by a graphite filament and the temperature was calibrated by pyrometer. The graphene CVD growth was performed by using ethene (C2H4) as a precursor in a UHV-CVD chamber (Riber). The pressure was in the range of 10−2 mbar during deposition and the flow of ethylene was 5 sccm. In addition, graphene samples were grown by a physical method in a MBE chamber (DCA), as described in ref. e-gun, and the pressure was in the lower 10−7 mbar range.Graphene was grown in a multi chamber ultrahigh vacuum system (UHV) with a base pressure in the lower 10μ-Raman spectroscopy using Renishaw In-viaFlex spectrometer and green laser light (514 nm) with a spot diameter of about 500 nm. Transport measurements were performed with a 4-tip scanning tunnelling microscope  with a high resolution scanning electron microscope (SEM) used as a tip-navigation tool.Quality of graphene was evaluated by hν = 100 eV with linearly polarized light for Fermi surface mapping and at hν = 30 eV with circularly polarized light for the energy scans around the K point of graphene. The circular polarization in these experiments was selected in order to avoid the “dark corridor” problem along the Γ − K direction for graphene−11 mbar.Photoemission experiments were performed at the nano-ARPES microscope of the ANTARES beamline of the SOLEIL synchrotron, equipped with two Fresnel zone plates  for focalization of the synchrotron radiation beam and an order selection aperture (OSA) to eliminate higher diffraction orders. Samples are mounted on a nano-positioning stage placed at the coincident focus of the Scienta R4000 analyser and the FZP focal point, which is responsible for the sample rastering during the image acquisition. The spatial resolution is determined by the FZP resolution and the mechanical stability of the sample stage. Experiments were performed at photon energies p(2 × 2) surface cells, each containing eight dimers on one side and passivated with H atoms on the other side. Further calculation details can be found elsewherePseudopotential density functional theory (DFT) total energies and atomic and electronic structures were calculated by Quantum EspressoHow to cite this article: Dabrowski, J. et al. Understanding the growth mechanism of graphene on Ge/Si(001) surfaces. Sci. Rep.6, 31639; doi: 10.1038/srep31639 (2016).


Many potential applications of graphene require its precise and controllable doping with charge carriers. Being a two-dimensional material graphene is extremely sensitive to surface adsorbates, so its electronic properties can be effectively modified by deposition of different atoms and molecules. In this paper, we review two mechanisms of graphene doping by surface adsorbates, namely electronic and electrochemical doping. Although, electronic doping has been extensively studied and discussed in the literature, much less attention has been paid to electrochemical doping. This mechanism can, however, explain the doping of graphene by adsorbates for which no charge transfer is expected within the electronic doping model. In addition, electrochemical doping is in the origin of the hysteresis effects often observed in graphene-based field effect transistors when operating in the atmospheric environment. Graphene is a monolayer of carbon atoms arranged in a honeycomb lattice . With thIn this paper, we present an overview of two mechanisms of graphene doping by surface adsorbates, namely, electronic and electrochemical doping and consider appropriate atomic and molecular dopants.+/SiO2 substrate [g, the type of carriers and their concentration in graphene can be tuned. The sign of the induced carriers is opposite to the sign of the applied gate voltage. A positive Vg induces electrons while a negative Vg induces holes. It was shown that for graphene on Si+/SiO2 the concentration of charge carriers induced by this method can be as high as 1013 cm−2 [Charge carriers, either electrons or holes, can be induced in graphene by the application of an electric field or by chemical doping. The electric field effect doping is usually performed in graphene-based field effect transistors (FET), in which charge carriers are induced by changing the electric potential between graphene and a gate, which can be, for example, a Siubstrate ,13. By vs, and nitrogen, Ns, substitutional atoms and leads to p- and n-type conductivity, respectively [s-doped graphene behaves like an n-type doped semiconductor but exhibits a low charge carrier mobility [Chemical doping involves interactions of graphene with other chemical species . There aectively –16. HoweThe surface transfer doping is non-destructive and occurs due to the charge transfer between graphene and surface adsorbates. Two mechanisms of charge transfer doping can be distinguished, electronic and electrochemical doping. The electronic doping is a consequence of the direct charge transfer between graphene and an adsorbate. This requires a difference in electronic chemical potentials at an interface, which is determined by the relative positions of the graphene Fermi level and the highest occupied (HOMO) and lowest unoccupied (LUMO) molecular orbitals of an adsorbate. If the LUMO of the adsorbate lies lower in energy than the Fermi level of graphene, G + W for p-doping and ΔG − W for n-doping, where ΔG is the free energy for the molecular reaction and W is the work function of graphene. The work function of graphene is expected to be similar to that of graphite, about 4.5 eV [G can be estimated from the Nernst equation the barrier heights are unknown and it must be assumed that they are sufficiently low for the reaction to occur [The electrochemical doping of graphene occurs when certain surface adsorbates participate in electrochemical redox reactions in which graphene plays the role of an electrode. Such reactions occur spontaneously if the total Gibbs free energy is negative and the diffusion and reaction barriers are sufficiently low for the reaction proceed at room temperature. The total Gibbs free energy change is given by Δt 4.5 eV –18. Althto occur .Eredox to the graphene Fermi level. The Eredox is the equivalent of the Fermi level of an electrolyte solution. It measures the energy required to add or remove an electron in redox couples and is usually referenced to the standard hydrogen electrode (SHE) [G and Eredox measure the same property and are related by ΔG = −n·F·Eredox, where n and F are the number of electrons and the Faraday constant, respectively. Furthermore, while the Eredox level lies higher (lower) than the Fermi level of graphene (EF), electrons (holes) induced by a reaction will flow to graphene until equilibrium is reached, Eredox = EF, making graphene n- (p-)type doped.Whether a reaction will result in p- or n-type doping of graphene depends on the relative position of the electrochemical redox potential de (SHE) . Thus, ΔIn contrast to electronic doping, which occurs instantaneously, electrochemical doping is a time-dependent process, which is affected by the rate of the reaction and diffusion rates of participating species. Therefore, in the case the reaction or diffusion rates are slower than the rate of change of gate voltage electrochemical doping can lead to hysteresis effects which are often observed in graphene-based FET devices –24.Vg with the increase of the potassium concentration is a clear indication of n-type doping. It was also observed that the mobility of the charge carriers decreases with the increase of the doping concentration. This is attributed to the additional scattering caused by the ionized potassium atoms. In agreement with the experimental results, DFT calculations have shown that K atoms act as electron donors [1 electronic level of potassium, The ability to dope graphene either n- or p-type has been demonstrated for various atomic and molecular adsorbates. Electropositive elements that easily donate their outer shell electrons are expected to be n-type dopants. Indeed, density functional theory (DFT) calculations predicted group I–III metals to be efficient electron donors for graphene –28. Elecn donors –28. ElecThe doping properties were also explored for different transition metal clusters  deposited on graphene by molecular beam epitaxy (MBE) . The Ti 4-TCNQ), shown in Whereas one expects that electropositive adsorbates will dope graphene n-type, electronegative species should lead to p-type doping. By using angle-resolved photoemission spectroscopy Gierz et al.  demonstr4-TCNQ molecule was considered to be a good candidate for p-type doping of graphene. The effect of the F4-TCNQ molecule on the electronic properties of epitaxial graphene was studied by synchrotron-based high resolution photoemission spectroscopy (PES) [4-TCNQ  . The inc [4-TCNQ  suggestsmolecule .The p-type doping of graphene can also be achieved by depositing other organic molecules such as pyrenetetrasulfonic acid (TPA) or tetracyanoethylene (TCNE) –38, show6H5CH3), It was observed that the exposure of graphene to a humid atmosphere causes p-type doping whereas exposure to toluene  and acidic (pH 1) conditions, respectively. Therefore, the reaction is inhibited in basic conditions and promoted in acidic conditions since the concentration of OH− is low.The value of Δ−4.82 eV . Under t2 substrate in the presence of O2 and H2O was confirmed experimentally by Raman spectroscopy measurements [2 and H2O leads to a stronger doping than the exposure to O2 or H2O alone. The hysteresis effects commonly observed in graphene FETs doped under atmospheric conditions [The electrochemical p-type doping of graphene on a SiOurements . It was nditions –24 were Similarly, the n-type doping of graphene caused by toluene can be explained by the electrochemical mechanism. In this case it is difficult to determine which particular reaction takes place. One possibility is the electrochemical oxidation of toluene to benzyl alcoholThis reaction is spontaneous as the change of the total Gibbs free energy is −0.5 eV. The released electron can be trapped by graphene and make it n-type doped.Vg at different stages of the toluene-doping experiment. Firstly, the sample was annealed in vacuum and the resistance peak at positive Vg, Vg showing that toluene acts as n-type dopant on graphene, R(Vg) measurement. The position of the resistance peak changes with the direction of the Vg sweep. The hysteresis vanished when the toluene is pumped out, but the doping effect remains, +/SiO2 substrates [2 interface. Indeed, the presence of water layers on the surface of SiO2 films grown on Si has been confirmed by Kelvin probe microscopy, X-ray spectroscopy and FTIR measurements [2. It has been found that for graphene FETs operating under ambient conditions the hysteresis is suppressed when Si+/SiO2 substrate was covered with a thin hydrophobic layer [2 interface. The binding energy of OH− to the SiO2 network was estimated to be about 0.3 eV indicating that OH− can easily diffuse through SiO2 at room temperature [− at the graphene/SiO2 interface and thus the rate of the reaction is influenced by the gate voltage. A positive Vg results in the reduction of the OH− concentration at the graphene interface promoting the further dissociation of H2O molecules, which consequently increases the concentration of holes in graphene, as it was observed experimentally [Hysteresis effects are often observed in graphene FETs based on Sibstrates –24. Suchurements –43. Wateic layer . Anotheric layer , supportperature . Therefomentally .3 and enhanced by NO2 [Some of these redox reactions were first discussed for diamond and carbon nanotubes ,45 and wd by NO2 . Similard by NO2 –50. It w4-TCNQ, have been proved to be very efficient p-type dopants. This type of doping occurs instantaneously and should not cause hysteresis effects. However, the ionized dopants become additional charge scatterers leading to a reduction in the charge carrier mobility.In this article we have revisited the surface dopants of graphene and the charge transfer mechanisms. The deposition of atoms or molecules is an effective method to control the type and concentration of charge carriers in graphene. Surface charge transfer can be mediated by two mechanisms, electronic and electrochemical doping. The electronic doping occurs by direct exchange of electrons between graphene and the adsorbates, either atoms or molecules. Electropositive elements that easily donate their outer shell electrons, such as potassium, are expected to be n-type dopants while adsorbates with a strong electron affinity, such as FThe electrochemical doping of graphene occurs as a result of redox reactions that can take place near the graphene surface. Unlike electronic doping it does not reduce the charge carriers mobility but it may require appreciable time to occur since the reaction and diffusion barriers have to be overcome. Moreover, the electrochemical doping model can explain the hysteresis effects usually observed in graphene based field effect transistors when operating in atmospheric environment.