Magnetic graphene-based materials have shown great potential for developing high-performance electronic devices at sub-nanometer such as spintronic data storage units. However, a significant reduction of power consumption and great improvement of structural stability are needed before they can be used for actual applications. Based on the first-principles calculations, here we demonstrate that the interaction between tungsten atoms and nitrogenized-divacancies (NDVs) in the hybrid W@NDV-graphene can lead to high stability and large magnetic anisotropy energy (MAE). More importantly, reversible switching between different magnetic states can be implemented by tuning the MAE under different electric fields, and very low energy is consumed during the switching. Such controllable switching of magnetic states is ascribed to the competition between the tensile stain and orbital magnetic anisotropy, which originates from the change in the occupation number of W-5d orbitals under the electric fields. Our results provide a promising avenue for developing high-density magnetic storage units or multi-state logical switching devices with ultralow power at sub-nanometer. Magnetic nanostructures with giant magnetic anisotropy energy (MAE) are of great interest for developing high-density spintronic data storage devices with the reduction of bit size1234689101112131415161718202122242627The stability of the TM atoms on graphene can be greatly enhanced if they are absorbed on the graphene with defects. Our recent work shows that defective graphene containing single vacancies (SVs) can anchor the TM atoms firmly at the centers of SVs while the MAE of the system remains largeIn this work, based on the first-principles calculations we demonstrate that the hybridization between the NDV-graphene and tungsten atoms can result in both high stability and large MAE for the W@NDV-graphene, meanwhile the whole hybrid system exhibits facile manipulation of MAE by external electric fields. We show that the conversion between different magnetic states in the system can be easily realized by controlling the switching of intermediate magnetic state upon the application of moderate electric fields. Such conversion of magnetic states hardly consumes energy, providing us a promising avenue to develop high-density magnetic storage units or multi-state logical switching devices with ultralow power at sub-nanometer. We also give an insight about the origin of such tunable MAE manipulated by external electric fields. We reveal that the change in the occupation number of W-5d orbitals under different electric fields plays an important role in determining the change in magnetic properties of the hybrid system, and the competition between the tensile stain and orbital magnetic anisotropy is responsible for the change in MAE under external electric fields.B. According to the spin density isosurface displayed in the inset of Our calculations of magnetic moments show that the W@NDV-graphene possesses a total spin magnetic moment of 2.24 μB at 1.0 V/Å. x − Ey = 16.36 meV, Ey − Ez = −21.17 meV, and Ez − Ex = −4.803 meV, respectively. Accordingly, the MAE of the hybrid system is determined to be 21.17 meV with an easy magnetization direction parallel to the graphene plane  and a hard magnetization direction perpendicular to the graphene plane . Applying external electric fields below 0.6 V/Å only gives rise to a perturbation of energy difference between each of the two magnetization directions, and the easy magnetization axis also does no change. Surprisingly, when the applied electric field is close to 0.8 V/Å, not only the energy difference between each of the two magnetization directions changes remarkably, but also the easy and hard magnetization axes are switched completely, i.e. the easy magnetization axis changes to z axis whilst the x axis becomes the hard magnetization axis. At 0.8 V/Å, the energy differences between each of the two magnetization directions become Ex − Ey = 47.979 meV, Ey − Ez = 12.284 meV, and Ez − Ex = −60.263 meV, respectively. The MAE of the whole system thus reaches a very high value of 60.26 meV, about three times of the original MAE value at zero electric field. When further increasing the electric field, the MAE only slightly decreases and both the easy and hard magnetization axes remain unchanged.To examine the feasibility of manipulating the MAE by external electric fields, we calculated the response of the magnetic properties of the W@NDV-graphene to the electric field. An electric field was applied to the W@NDV-graphene by the dipole layer methodBT for 1-bit storagems = 6.1 K, larger than those of most single-molecule magnets in the current stage such as Fe4-organometallic complexes (Tms = 0.5 K)ms = 5.5 K)34Both high stability and large MAE indicate that the W@NDV-graphene hybrid system can be used as a single-molecule-magnet-like unit for magnetic storage. Considering that the size of the W@NDV-graphene supercell used in this work is about 1.5 nm, if the W@NDV-graphene units are assembled on this scale, the spacing between each unit is large enough to avoid the coupling interaction with each other (see F). When the electric field is increased from 0 to 0.8 V/Å, the PDOS of W-5d orbitals has a drastic change, indicating that the electric field induces variation of occupation number of W-5d orbitals. Based on the second perturbation theoryF is related to the change of MAE as follows: x, Ly and Lz are the angular momentum operators, respectively. According to F, and its behavior is substantially determined by the denominator. We have calculated the spin orbital coupling (SOC) matrix elements between different d orbitalsyz, dxz and dz2 orbitals have large contributions to the MAE. Moreover, from the second term in yz and dz2 orbitals dramatically through Lx operator, leading to the change in MAE.We now turn to explore the mechanism of manipulating MAE by electric fields for the present hybrid system. The influence of the electric field on the occupation number of W-5d orbitals in the W@NDV-graphene was first considered. F contribute to the MAE are dz2, dyz and dxz. The magnetization direction is mostly determined by the SOC of up-spin dz2 with dxz state through Ly operator, and then the easy axis is parallel to the graphene plane. When an external electric field is applied, the occupation number of the down-spin dyz orbital of W atom is reduced, consequently shifting the dyz orbital to the higher energy side. The reduced electron occupation number of dyz orbital is also visualized by the charge-density difference between 0 and 0.8 V/Å, as plotted in the inset of z2 orbital. As a result, the coupling interaction between down-spin dyz and dz2 through Lx operator is reduced. In addition, both the occupied up-spin dyz and unoccupied up-spin dxz increase, indicating that the coupling interplay between dyz and dxz through Lz operator is enhanced while the coupling interplay between dz2 and dxz is reduced due to the decrease of up-spin dz2 and the shift to higher energy side of up-spin dxz. Therefore, the MAE of the hybrid system increases and the magnetization direction of the easy axis changes.We can classify the contributions of PDOS to MAE into two groups, one involving the coupling between the same spin states and the other including the coupling between different spin states (i.e. the spin-flip terms). Previous investigations of FeCo/MgO(011) thin films have revealed that the spin-flip terms are much smaller when the exchange splitting between majority and minority states is largeOur further calculations show that applying electric fields to the W@NDV-graphene can cause an increase in the distance between the W atom and graphene plane (e.g. the increment reaches 0.537Å at 0.8 V/Å). This indicates that a tensile stain perpendicular to the graphene plane is induced. 2/Pt(100) have revealed that the charge transfer has great influence on the orbital magnetic anisotropy B. According to Bruno's relationshipW is an average value of the SOC coefficients of W atom, and Δml,W is the difference of orbital magnetic moment between the easy and hard magnetization directions. Our results shown in B), indicative of a sufficient exchange splitting of the W atom. After applying external electric fields, the reduction in charge of W atom leads to an increase in full spin magnetic moments . At high electric fields, the suppression of MAE by tensile strain becomes stronger than the increase of MAE caused by OMA, causing the drop of total MAE. Evidently, from the view of applications, the MAE manipulation by electric fields is more efficient if the hybrid system is under stress-clamping state.Overall, we have demonstrated that the amplitude of MAE and magnetization direction of the W@NDV-graphene can be facilely manipulated by external electric fields. The hybrid system has both high stability and large MAE value. The change in the occupation number of W-5d orbitals under electric fields plays an important role in determining the change in magnetic properties of the hybrid system. Applying electric fields can cause the change in both tensile stain and orbital magnetic anisotropy. The competition between these two parameters determines the change in MAE under electric fields. Considering the tunable MAE manipulated by external electric fields allows us to largely reduce the energy consumption during operating the magnetic states, we expect such kind of hybrid system is promising for developing high-performance magnetic storage and multi-state logical switching devices with ultralow energy consumption at sub-nanometer scale.−7 eV. The NDV-graphene was modeled by a 6 × 6 graphene supercell with a divacancy defect, in which four carbon atoms around a divacancy was substituted by four nitrogen atoms. To avoid interactions between the layers, vacuum region of 15 Å in the direction normal to graphene was employed.The first-principles calculations were performed in the framework of spin-polarized density functional theory (DFT) using the projector augmented wave (PAW)42G.X.G. performed the theoretical calculation. J.G.W. conceived and provided advice on the analysis. J.G.W., G.X.G. and Y.H. wrote the paper. H.B.S., F.Q.S., J.J.Z. and G.H.W. participated in analysis and discussion. All authors participated in the data discussion.Supplementary Information


High broad‐band photoresponsivity of mechanically formed InSe–graphene van der Waals heterostructures is achieved by exploiting the broad‐band transparency of graphene, the direct bandgap of InSe, and the favorable band line up of InSe with graphene. The photoresponsivity exceeds that for other van der Waals heterostructures and the spectral response extends from the near‐infrared to the visible spectrum. Eg = 1.26 eV at T = 300 K, which can be increased due to quantum confinement by reducing the number of atomic layers in the crystalline sheet.8) and high carrier mobilities (μ = 0.1 m2 V−1 s−1) at room temperatureThe development of van der Waals (vdW) heterostructures fabricated by mechanically stacking 2D crystals has led to the discovery of fundamental physical phenomena and to the realization of 2D functional devices ranging from sensitive phototransistors to tunnel diodes.This burgeoning research field is still in its infancy and offers exciting opportunities for discoveries and the realization of functional devices that implement InSe in combination with other vdW crystals. Crucial to these future developments is the formation of good interfaces and Ohmic contacts to the optically active InSe layer. Devices currently used are mostly prepared from InSe with metal contacts.n‐InSe–graphene heterostructures with a high photoresponsivity (up to ≈105 A W−1 at λ = 633 nm), not yet achieved in other 2D vdW crystals, and with a spectral response that extends from the near‐infrared to the visible spectrum. The highest photoresponsivity is observed in device architectures where the InSe and graphene layers are vertically stacked and an optical window is created by overlapping the two graphene electrodes, thus enabling sensitive photodetection. These heterostructures provide innovative device architectures that enable access to fast electron speeds and high broad‐band spectral response.Here, we report on van der Waals graphene/InSe heterostructures formed by mechanical contact due to attractive vdW forces at the graphene/InSe interface. We exploit the broad‐band transparency of graphene and the favorable band line up of graphene with InSe to create high‐performance photodetectors. We demonstrate vertical and planar graphene–1.03Se0.97. The γ‐polytype crystal structure of InSe was probed by X‐ray diffraction (XRD) using a DRON‐3 X‐ray diffractometer in a monochromatic Cu–Kα radiation of wavelength λ = 1.5418 Å. The primitive unit cell contains three InSe layers each of which has a thickness of 8.320 Å and consists of four covalently bonded monoatomic sheets in the sequence Se‐In‐In‐Se; along the c‐axis, the primitive unit cell has a lattice constant of c = 24.961 Å and, within each a–b plane, atoms form hexagons with lattice parameter a = 4.002 Å. In its bulk form, InSe contains native donors due to In‐interstitial atoms.T = 300 K, we obtain an electron density n ≈ 1021 m−3, an electron mobility μ = 0.1 m2 V−1 s−1, and a Fermi energy EF ≈ 0.21 eV below the conduction band minimum.Our InSe crystals were grown using the Bridgman method from a polycrystalline melt of Int that range from 20 to ≈100 nm, over which InSe retains a direct bandgap−2–10−1 m2 V−1 s−1).2/Si substrate and patterned by electron beam lithography (EBL) into two contacts prior to the transfer of an InSe nanoflake . Figuren‐InSe–graphene planar device structure with Au‐metal contacts on two graphene layers. The n‐Si layer serves as a gate electrode and the graphene layers, g1 and g2, serve as source and drain to an InSe channel of length l ≈ 2 μm, width w ≈ 10 μm, and thickness t ≈ 30 nm.The InSe nanosheets were prepared from the as‐grown crystals by mechanical exfoliation using adhesive tape and have thicknesses Vs applied between the two graphene electrodes. The I–Vs characteristics are symmetric relative to negative and positive values of Vs and, for each Vs, the current increases with increasing gate voltage Vg. Separate Au‐contacts on each graphene layer enable us to assess their conductivity: they exhibit a linear dependence of the current on the bias voltage and a minimum conductance at around a gate voltage Vg ≈ 60 V, corresponding to p‐type doping and hole concentration of p ≈ 5 × 1012 cm−2, typical for CVD graphene  and decay (τd) times of the current of τd ≈ τr < 0.1 ms.As shown in Figure I–Vs characteristics are preserved under optical illumination, see FigureI, is weakly modulated by the gate voltage Vg and its dependence on Vg is opposite to that observed for the dark current, i.e., ΔI decreases with increasing Vg from −60 to +60 V. As shown in Figure I is slower than that of the dark current with a rise time τr ≈ 1 ms and decay time τd ≈ 10 ms, which are not affected by Vg.The linearity and symmetry of the High‐quality interfaces, crystal lattice alignment and control over the relative band alignment are essential for tuning and controlling electrical injection of carriers from an electrode into a semiconductor layer. The cleaved facets of InSe layered crystals are atomically smooth, as confirmed by our atomic force and transmission electron microscopy studies (not shown), and contain a low‐density of surface states,Figureχg = −4.5 eV) is larger than that of bulk InSe (χInSe = −4.6 eV);φg) can be increased relative to that of n‐type InSe (φInSe = – 4.81 eV) by applying a gate voltage. Compared to InSe, this gate effect is stronger in graphene due to the low density of states of graphene when the chemical potential is close to the Dirac point. In particular, over the range of applied gate voltages Vg from −60 to +60 V, we find that φg remains higher than or comparable to φInSe, with the Fermi level increasing from ca. −0.3 eV (φg = −4.85 eV at Vg = −60 V) to 0 eV (φg = −4.5 eV at Vg = +60 V) relative to the neutrality point of the Dirac cones. Thus, we infer that at equilibrium, electrons tend to transfer from graphene to InSe where they can form an accumulation layer at the interface with graphene  over the range of Vg from −60 to +60 V, we estimate that l changes by ca. ±20%. This dependence was not observed in the Au–InSe–Au planar devices where the dark current and photocurrent both decrease or increase with varying the gate voltage.The band alignment is supported by the opposite dependence on the applied gate bias observed for the dark current and the photocurrent. To explain the monotonous increase of the dark current with increasing g Figure b and theg Figure b, we shoVs, electrons and holes that are photoexcited in InSe are swept by the electric field in opposite directions and are extracted at the graphene electrodes to generate a photocurrent ΔI = [etαP/hv]lτ/τt, where P is the incident power, α is the absorption coefficient of InSe at the photon energy hv, e is the electron charge, t is the thickness of the InSe layer, and τl/τt is the ratio of the minority carrier lifetime (τl) and transit time (τt) of electrons in InSe . Thus the photoresponsivity R of our device can be approximately described by R = ΔI/P = [etα/hv]τl/τt. Furthermore, we can express the external and internal quantum efficiencies as EQE = R(hv/e) = [tα]τl/τt and IQE = τl/τt, respectively. These relations indicate that large values of R, EQE, and IQE can be achieved if the lifetime of the minority carriers (holes) is longer than the transit time of electrons.Under an applied bias FigureVs = 2 V, Vg = 0 V, and λ = 633 nm: R is strongly dependent on the optical power P and reaches a maximum value of R = 4 × 103 A W−1 at the lowest incident power P = 10 pW, which corresponds to τl/τt ≈ 3 × 105 for α = 106 m−1 at hv = 1.96 eV (λ = 633 nm) and t = 30 nm. From the measured values of R, we estimate a maximum external quantum efficiency EQE = R[hv/e] ≈ 5 × 103 and a specific detectivity D* = R(A/2eI)1/2 ≈ 1010 m W−1 s–1/2, where A = 20 μm2 is the area of the device and I = 0.6 × 10−6 A is the dark current at Vs = 2 V and Vg = 0 V. As shown in Figure R and EQE decrease by a factor of ≈30 and 50, respectively, for hv decreasing from ≈2 to ≈1.3 eV.R with increasing P is analogous to the P‐dependence reported previously for the photoresponsivity of other graphene‐based photodetectors.τl/τt. Thus the decrease of R with increasing P suggests a decrease of τl and/or an increase of τt. The transit time of photoinduced carriers can be increased by the power due to enhanced scattering.τl. Since this behavior is not observed in devices based on thick bulk InSe flakes, we infer that Auger‐like carrier recombination on traps is enhanced in these 2D vdW crystals due to stronger Coulomb interactions. Finally, R decreases linearly with decreasing applied bias voltage Vs, as expected from the increase in the electron transit time τt = l2/μVs: for Vs = 2 V, l = 2 μm, and μ = 0.1 m2 V−1 s−1, we estimate a transit time τt ≈ 2 × 10−11 s and a carrier lifetime τl ≈ 5 × 10−6 s.The decrease of R and EQE for the InSe/graphene photodetector are significantly larger (by a factor of 104) than those measured in devices in which Au‐contacts replace the two graphene electrodes on the InSe flake, see Figure φAu = −5.2 eV) is lower than that for n‐InSe (φInSe = −4.8 eV), thus resulting in a Schottky contact at each interface of the heterostructure and in non‐linear I–Vs (not shown). A barrier may also form due to a defective interface between InSe and the metal contact, as reported recently by Feng et al.Our values of n‐InSe. Since graphene is optically transparent, high photoresponsivity can also be achieved in multilayer systems where the InSe and graphene layers are vertically stacked with the top graphene layer acting as a broad‐band optical window. These structures have the additional advantage that the separation of the graphene electrodes, determined by the thickness t of the InSe flake, is smaller leading to a more sensitive photodetector through a decrease of the carrier transit time. However, compared to the planar devices, for the fabrication of the vertical graphene/n‐InSe/graphene structure, additional process steps are required to pattern the upper graphene contact layer. Here we have developed different processing steps, corresponding to devices architectures of type A  and bottom‐graphene (g2) layers overlap an InSe nanoflake of thickness t = 27 nm. The size of the optical window on the top surface of the InSe is approximately 4 μm × 4 μm 1/2 ≈ 1013 m W−1 s–1/2, where A = 16 μm2 is the area of the device and I = 0.4 × 10−9 A is the dark current at Vs = 2 V and Vg = 0 V. Finally, we note that a power law relation of the form R = −nCP, with n ≈ 0.7 where C is a constant, provides a good empirical fit to data of R for both vertical  has been achieved in phototransistors based on graphene and colloidal nanocrystals,Overall, the photoresponsivity measured in our planar and vertical devices is significantly larger than that reported for photodetectors based on graphene and/or vdW 2D crystals at similar wavelengths and laser powers, which do not exceed values of 10In summary, our findings demonstrate that mechanically formed heterojunctions of InSe and graphene have optical and electrical properties with potential for applications in optoelectronics. Our innovative fabrication methods could be extended to other material systems and device architectures. It should be possible to achieve different band alignments and potential profiles by combining InSe with other layered semiconductors, and by selecting the doping, n‐ or p‐type, of the component layers, which is difficult to achieve in other van der Waals 2D crystals. For example, further developments include systems in which the InSe layer is replaced by multiple bandgap heterostructures, such as InSe/GaSe or InSe/hBN junctions and quantum wells. This type of vertical heterostructure architecture would enable fast electron speeds; photonic applications are also attractive due to the potential to access an even wider spectral range than that achieved in our current devices.λ = 633 nm), an XY linear positioning stage and an optical confocal microscope system. The laser beam was focused to a diameter d ≈ 1 μm using a 100× objective and the maps were measured at low power (P < 0.1 mW) to avoid lattice heating.Images of the InSe flake and graphene topography were acquired by atomic force microscopy (AFM) in noncontact mode under ambient conditions. The experimental setup for the photoconductivity maps comprised a He–Ne laser (f = 7 Hz) and focused onto the device. The photocurrent signal was measured using a Stanford SR830 lock‐in amplifier (integration time constant of t = 3 s).For the measurement of the photocurrent spectra Figure b, light f = 10−1–105 Hz. The temporal dynamics of the photocurrent  and illumination by a mechanically modulated He–Ne laser with λ = 633 nm, P ≈ 10−5 W cm−2 and frequency f in the range 1–400 Hz. The photocurrent signal was measured using a Tektronix DPO 4032 digital oscilloscope and a Keithley2400 was used as a DC voltage source. For these AC measurements, the device was connected in series with a 1 MΩ resistor. We measured the voltage drop across the resistor, which enabled us to measure voltage signals with a low noise level.The measurements of the DC dark current and photocurrent versus the applied voltage Figures b and 2a Fabrication of Planar Graphene/InSe Heterostructures: The exfoliated InSe flakes were integrated into 2D planar structures incorporating either Au or graphene contacts fabricated by EBL. In these devices, CVD‐graphene, grown on a copper substrate, was transferred to a SiO2/Si substrate (SiO2 layer thickness of 300 nm) by first coating it with poly(methyl methacryrate) (PMMA) and placing on the surface of a FeCl3 etchant  to remove the copper. The PMMA/graphene film was then placed in a HCl bath and subsequently rinsed in deionized water and then mechanically placed on the substrate, before removing the PMMA by immersion in acetone, rinsing in isopropyl alcohol and annealing in an Ar:H (95:5) mixture at 400 °C for several hours. The graphene was then patterned using electron beam lithography and etching in an Ar/O2 plasma and, following removal of PMMA, as described above, transferred to a N2 atmosphere where the InSe was exfoliated and transferred to the substrate. Au/Ti contact pads were formed to provide contacts to the graphene layers, resulting in the structures shown in Figure Fabrication of Vertical Graphene/InSe Heterostructures: The fabrication of the vertical devices is more complex than for the planar devices and requires several steps. We have used two methods, referred to as methods A and B, both described below.Method A: Devices of type A  which had been cleaned in acetone, isopropyl alcohol and an Ar/O2 plasma. A suitable monolayer was then located using a microscope. The InSe flake (thickness t = 27 nm) and top monolayer graphene electrode were also formed by mechanical exfoliation using the low‐tack low‐residue tape, but were deposited onto PMGI (poly dimethyl glutarimide)/PMMA coated substrates for later transfer. Suitable flakes of each material were selected and isolated (using a sharp implement to score a circle in the PMMA around each flake); the layer of PMGI underneath each of the isolated flakes was then dissolved using MF‐319 developer. The flakes, along with their supporting PMMA membranes, were then floated onto the surface of a beaker of distilled water and deionized water. Each membrane was captured on a PMMA‐coated stainless‐steel washer and the washer was attached to a “transfer arm” connected to a set of micromanipulators; these micromanipulators were then used to position each flake over the previous layer(s) of the device before the membrane was brought into contact with the substrate. Finally, the membrane was cut free of the washer, baked onto the substrate at 130 °C and the PMMA was washed away in an acetone bath to leave the desired flake in place. Au/Cr contact pads (50/3 nm) were formed to provide contacts to the two graphene layers.A Figure a were foMethod B: Devices of type B , graphene was transferred and cleaned, as described above. InSe flakes were deposited by exfoliation and a flake of thickness t was selected using optical microscopy . The graphene layer was patterned and etched in an Ar/O2 plasma to leave a graphene strip, which provided a continuous connection to the (unetched) graphene under the selected InSe flake. Au/Ti contact pads (110/15 nm) were then deposited at each end of the lower contact strip. In the same step Au/Ti pads were also formed to provide contacts to the upper graphene strip, which was deposited in a later fabrication step (see below). An isolation layer formed by exposing a negative resist, AR‐N, using electron beam lithography was then formed; this covered the edges of the InSe flake and the regions of the lower graphene contact close to the flake, to avoid the formation of electrical shorts between upper and lower graphene strips. A window with an area of approximately 2 μm × 2 μm was formed in the AR‐N layer on the top surface of the InSe to facilitate the mechanical contact with the upper graphene layer. Finally, a second layer of graphene was transferred as described above; in this case the supporting PMMA film, which was added to the graphene/copper before etching in FeCl3, was not removed, but instead used as the resist layer in a further stage of electron beam lithography in which the upper graphene contact strip was defined by etching in an Ar/O2 plasma. The resulting upper contact runs at right angles to the lower strip to form a graphene/InSe/graphene vertical heterostructure and overlaps the Au/Ti contacts formed at an earlier stage of the process .I(V) characteristics of the devices A and B is possibly related to the inclusion of organic residues at the graphene/InSe interface. This may arise from polymeric layers used during the process (including the negative resist layer for type B devices); in addition the InSe layer is submerged in water during the transfer of the upper graphene contact.The non‐linearity of the As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


Sarcoidosis is a multisystem granulomatous disorder invariably affecting the lungs. It is a disease with noteworthy variations in clinical manifestation and disease outcome and has been described as an “immune paradox” with peripheral anergy despite exaggerated inflammation at disease sites. Despite extensive research, sarcoidosis remains a disease with undetermined aetiology. Current evidence supports the notion that the immune response in sarcoidosis is driven by a putative antigen in a genetically susceptible individual. Unfortunately, there currently exists no reliable biomarker to delineate the disease severity and prognosis. As such, the diagnosis of sarcoidosis remains a vexing clinical challenge. In this review, we outline the immunological features of sarcoidosis, discuss the evidence for and against various candidate etiological agents (infective and noninfective), describe the exhaled breath condensate, a novel method of identifying immunological biomarkers, and suggest other possible immunological biomarkers to better characterise the immunopathogenesis of sarcoidosis. Sarcoidosis is a multisystem, inflammatory disorder of obscure aetiology. Its defining histopathology is the existence of noncaseating epithelioid granulomas with accompanying mononuclear cell infiltration and microarchitecture destruction , 2.Although sarcoidosis involves the lungs in >90% of cases, it also affects the heart, skin, eye, and central nervous system . This acSarcoidosis has been reported in all ethnic and racial groups with the majority of studies recording a peak incidence of 20–39 years of age for both males and females and a bimodal distribution whereby women have another peak incidence at 65–69 . Disease remission occurs in as many as two-thirds of patients, usually in the first 3 years after diagnosis. Other patients have chronic unremitting sarcoidosis which may subsequently lead to lung fibrosis . The errex vivo lymphocyte stimulation in the peripheral blood and bronchoalveolar lavage fluid (BALF) of sarcoidosis patients, describes the exhaled breath condensate, an innovative method of identifying immunological markers, and proposes novel immunological markers to better characterise the immunopathogenesis of sarcoidosis.This review outlines the current understanding of sarcoidosis, with reference to reg cells), notably CD4+CD25brightFoxP3+ cells . . 4]. β (TGF-β), MMP, and insulin growth factor-1 (IGF-1)) [Persistent granulomatous inflammation can lead to fibrosis. The immunological mechanisms leading to a fibrotic outcome remain undetermined. Nonetheless, various cytokines which are able to support a fibrotic response have been found at disease sites in patients with sarcoidosis ) , 36. H1 to TH2 cytokine predominance may occur in chronic sarcoidosis in response to persistent inflammation. TH2 cytokines such as IL-13 increase TGF-β production . NKT cells are mostly CD4+ and express an invariable TCR [γ, therefore suggesting that NKT cells exert regulatory activity which prevents disease progression [NKT cells produce Table TCR . Moreovegression .The aetiology of sarcoidosis remains unclear. A myriad of observations have supported the notion that sarcoidosis can be caused by environmental and infectious agents. Moreover, based on chronic beryllium disease, an analogous granulomatous lung disease, it has been speculated that one or more antigenic stimuli may be involved in the pathogenesis of sarcoidosis. Therefore, it is highly likely that the development of a sarcoidosis reaction to an antigen depends on a combination of genetic polymorphisms, the host's immune status, and exposure to environmental agents . A Case Control Etiologic Study of Sarcoidosis (ACCESS) demonstrated that sarcoidosis patients were 5 times more probable than controls to report a parent or sibling with sarcoidosis [Both family and genetic host studies have recognised genes that are responsible for this genetic susceptibility. Twin studies prove that monozygotic twins are more concordant for sarcoidosis than dizygotic twins. Moreover, familial aggregation of sarcoidosis can be seen worldwide. The multicentre study entitled coidosis .Genetic linkage studies on German families revealed a strong linkage to chromosome 6p. This led to the discovery of butyrophilin-like 2 (BTNL2), a costimulatory molecule within the MHC locus. Single nucleotide polymorphisms (rs2076530 G → A) in BTNL2 may affect T-cell regulation and activation .The genome-wide association study conducted by Hofmann and colleagues  revealedHLA class II are cell surface proteins that prime the adaptive immune system to antigens. Sarcoidosis is associated with the DR subtypes of class II antigens. HLA-DRB1∗01 and HLA-DRB1∗04, are negatively associated with sarcoidosis, whereas HLA-DRB1∗03, HLA-DRB1∗11, HLA-DRB1∗12, HLA-DRB1∗14 and HLA-DRB1∗15 have been shown to increase the risk of sarcoidosis. HLA-DRB1∗03 is associated with Löfgren's syndrome (~80% of patients with Löfgren's syndrome experience disease remission). Finally, the HLA-DRB1∗1501/DQB1∗0602 haplotype was associated with severe and chronic pulmonary sarcoidosis .α is an essential mediator for granuloma formation. Variants of the TNF gene confer a 1.5-fold increased risk of having sarcoidosis [TNFcoidosis . Apart fcoidosis –52.In some populations, variations in the gene that encodes for RAGE (a transmembrane receptor) have been associated with an increased risk of sarcoidosis. However, the close proximity of this gene to the MHC region makes it difficult for one to ascertain if this association is due to linkage with neighbouring HLA genes .There were no associations between polymorphisms in genes for vitamin D receptor or serum angiotensin-converting enzyme , 54.γδ or αβ genes. The expression of specific Vα, Vβ, or γδ+ TCR genes in blood, lung, and at sites of Kveim-Siltzbach skin reactions implies that sarcoidosis is an antigen-driven disorder. There is a subpopulation of T cells (AV2S3+ (Vα2.3) CD4+ T cells) from BALF of HLA-DR∗0301 sarcoidosis patients which is unique to sarcoidosis. Moreover, it has been shown that the amount of AV2S3+ BALF T cells at the onset of sarcoidosis correlates positively with prognosis, suggesting that AV2S3+ T cells may offer some protective function against sarcoidosis [The T-cells at sites of inflammation in sarcoidosis exhibit a restricted repertoire of TCR coidosis .Numerous pathogens have been implicated and investigated in the etiology of sarcoidosis. Moreover, spatial crowding of unrelated sarcoidosis cases suggests that sarcoidosis can also be a result of exposure to environmental agents , 3. Noneex vivo studies that employed flow cytometry to investigate peripheral blood lymphocytes in sarcoidosis patients demonstrated a greater activation of nonstimulated CD4+ and CD8+ BALF T cells compared to peripheral blood lymphocytes. This showed that the sarcoid immune response is largely compartmentalised to disease sites [Preliminary se sites . H1/TH2 model is under scrutiny as it oversimplifies the immunopathogenesis of sarcoidosis [+ T cells expressing IL-4 and IFNγ obtained from the peripheral blood of sarcoidosis patients did not differ significantly from that of healthy controls [H1 and TH2 levels of cytokine positive CD4+ T cells compared to healthy controls [The Tcoidosis . For inscontrols , 58. Othcontrols , 60, empγ secreting CD4+ lymphocytes in BALF and peripheral blood of sarcoidosis patients is insignificant [+ lymphocytes were stimulated with ionomycin and phorbol 12-myristate acetate, there was an appreciable increase in secreted IFNγ but a decrease in IL-4 expression in sarcoidosis patients compared to controls [γ+/IL-4+ CD4+ T cell ratios in sarcoidosis patients. Lower ratios were demonstrated in scleroderma and in patients with idiopathic pulmonary fibrosis. It has also been shown that upon stimulation, compared with controls, there are increased numbers of CD4+IFNγ+ cells in both BALF and induced sputum of patients with sarcoidosis [Under unstimulated conditions, the difference in the percentages of IL-4 and IFNnificant , 58. Aftcontrols . Moreovecoidosis , 89.H1 than TH2 cytokines in both the BALF and peripheral blood of sarcoidosis patients and more CD4+ T cells in BALF express TH1 receptors  than CD4+ T cells in the peripheral blood. [+ T cells largely express TH1 cytokines, interestingly, following stimulation, only 80% and 40% of CD4+ IL-4+ cells concurrently produce IFNγ and IL-2 respectively, thus demonstrating that activated BALF lymphocytes of sarcoidosis patients are capable of a complex, concurrent production of TH1 and TH2 cytokines [After stimulation, more T cells express Tl blood. . Althougytokines . H1/TH2 equilibrium, Nureki and colleagues [H1 and TH2 chemokines (interferon-inducible protein-10 (IP-10) and thymus and activation-regulated chemokine (TARC)) were both increased in the serum of sarcoidosis patients. This was in agreement with previous findings that demonstrated elevated BALF and peripheral blood IL-13 (a TH2 cytokine) mRNA levels [H2 cell preponderance occurs in the peripheral blood of sarcoidosis patients and that this, together with the generalised intensification of TH1 activity, gives the appearance of an increase in both TH1 and TH2 circulating cytokine expression in sarcoidosis patients compared to healthy controls [To further explore this dichotomy of blood Tlleagues  showed tA levels . Therefocontrols .H17 cells have also been implicated in the induction of granuloma formation [H17 related cytokine levels in both BALF and peripheral blood [+ T cells in patients with Löfgren's syndrome compared to healthy controls [H1 cytokine expression at disease sites, indicate that TH17 cells have a systemic role in patients with non-Löfgren's disease and is involved in sarcoidosis progression [H1/TH2 balance, the regulatory mechanisms in the peripheral blood of sarcoidosis patients, and the functional significance of TH17 cells.Tormation . Flow cyal blood . Anothercontrols . These dgression . TherefoThe diagnosis of sarcoidosis is never secure. Clinico-radiological findings alone are often insufficient to confirm the diagnosis of sarcoidosis. It needs to be supported by histological evidence showing noncaseating granulomas. This warrants a tissue biopsy which is invasive . This maExhaled breath condensate (EBC) has been subjected to intensive research as it provides a noninvasive alternative for sampling the airway and alveolar space a promising source of biomarkers for a variety of lung conditions –97. DuriA number of immunological biomarkers have been recognised in EBC. However, there exists no sufficiently sensitive and specific marker for diagnosing and predicting the prognosis of sarcoidosis . α, PAI-1, and IGF-1 levels in EBC were closely positively correlated with BALF samples from sarcoidosis patients. Conversely, IL-6 levels were negatively correlated with that which is in BALF. The propensity of IL-6 to form complex molecular forms of higher molecular weight could account for this discrepancy [β1, PAI-1, TNFα, IL-8 and vascular endothelial growth factor in sarcoid EBC. However, the small sample size and the failure to make comparisons with healthy controls limited the usefulness of this study [TNFcrepancy . Anotheris study .Exhaled eicosanoids ) and cysteinyl leukotrienes were also found to be elevated in the BALF and EBC of sarcoidosis patients . In a laβ, and the chemokine ligand (CCL18) [Cellular and molecular biomarkers previously discovered in BALF and serum of sarcoidosis patients could also serve as biomarkers. These include eosinophils, neutrophils, serum angiotensin converting enzyme (ACE), neopterin, chitotriosidase, TGF- (CCL18) , 104–107 (CCL18) , 109. Se (CCL18) . Amongst the above-mentioned biomarkers, ACE is the most contentious as it has been shown to have poor sensitivity and specificity  and its Given the multifactorial nature of sarcoidosis, no ideal markers for detecting and monitoring the clinical course of sarcoidosis exist. It is very likely that a combination of markers will be required. Although EBC has advantages over BAL  , the lacγ plays a pivotal role in the immunopathogenesis of sarcoidosis. MicroRNA-29 (miR-29) and T-bet have been shown to modulate its production [IFNoduction , 115. γ-secreting T cells. This reduction skews the immunological response towards a TH1 lineage by initiating a positive feedback loop which enhances IFNγ production. This also suggests that the up-regulation of miR-29a and miR-29b can mitigate IFNγ expression [MicroRNAs are noncoding RNA that can inhibit the production of mRNA. The miR-29 family is made up of four members. Amongst these four, miR-29a and miR-29b were found to be downregulated in IFNpression , 116. Abpression –119. It pression . γ production. It binds to a number of enhancers and to the promoter region of the IFNγ gene to promote IFNγ transcription [γ expression [As previously mentioned (see immune reactions in sarcoidosis), T-bet is a transcription factor necessary for IFNcription . T-bet epression , 123 in pression , coeliacpression , Crohn'spression , and Behpression , 127. Mopression . Howeverpression , 130. UnDespite nearly 140 years of extensive research, the aetiology and pathogenesis of sarcoidosis and the mechanisms that regulate the immune reactions in the peripheral blood remain undetermined. Moreover, given its variable clinical manifestation and the lack of a reliable diagnostic test with uniformed reference values and measurements, diagnosing sarcoidosis remains a clinical conundrum for many physicians. Given that the majority of sarcoidosis patients have pulmonary involvement, EBC could be used as a non-invasive method to diagnose sarcoidosis. Besides being a potential immunological biomarker of sarcoidosis, interferon modulator levels in EBC and the peripheral blood can be compared to elucidate the regulatory mechanisms in the peripheral blood. Results from such studies may also explain the pathology underpinning the peripheral anergy seen in sarcoidosis.


Exfoliated monolayer graphene flakes were embedded in a polymer matrix and loaded under axial compression. By monitoring the shifts of the 2D Raman phonons of rectangular flakes of various sizes under load, the critical strain to failure was determined. Prior to loading care was taken for the examined area of the flake to be free of residual stresses. The critical strain values for first failure were found to be independent of flake size at a mean value of –0.60% corresponding to a yield stress up to -6 GPa. By combining Euler mechanics with a Winkler approach, we show that unlike buckling in air, the presence of the polymer constraint results in graphene buckling at a fixed value of strain with an estimated wrinkle wavelength of the order of 1–2 nm. These results were compared with DFT computations performed on analogue coronene/PMMA oligomers and a reasonable agreement was obtained. These experiments involved the simple bending of a tiny flake by an indenter on an AFM set-up and the force-displacement response was approximated by considering graphene as a clamped circular membrane made by an isotropic material. To date there are no reported data, as yet, on pure axial stretching of graphene monolayers to fracture. Furthermore there is still an uncertainty concerning the ultimate tensile strain to failure which expected to be higher than even 30% making graphene a very ductile material indeed in spite of its very high stiffness.Graphene consists of a two-dimensional 2D) sheet of covalently bonded carbon and forms the basis of both 1D carbon nanotubes, 3D graphite but also of important commercial products, such as, polycrystalline carbon (graphite) fibres. As a single defect-free molecule, graphene is predicted to have an intrinsic tensile strength higher than any other known materialD sheet oPrevious studies by us and others have reported the effect of applied strain on the Raman spectrum G and 2D bands) under compression and tension loading histories3D bands uIn this present study we extend the work reported previously by examining a wide range of graphene sizes (length-to-width ratio) embedded in the SU8/PMMA matrix system. Special attention is given to the efficient transfer of stress from the polymer to the inclusion (monolayer graphene) as assessed by the value of Raman wavenumber shift per % of applied axial strain. For the interpretation of the acquired data, we consider the general problem of buckling of an embedded plate for which the “bonding” between the flake and the surrounding medium is modelled by linear elastic springs that act only in the z-direction, perpendicular to the graphene plane (Winkler's approach). A constitutive equation yielding the critical strain to compression failure as a function of the graphene elastic constants, the dimensions of the graphene flakes and the Winkler's modulus is derived. The same methodology has been used earlier in the modelling of instability problems in the case of embedded carbon nanotubes7From the full range of experimental results the model predicts a universal strain value for failure, for all reported flakes regardless of their specific dimensions. This is because the critical strain to failure in compression is mainly affected by the presence of polymer and the magnitude of graphene-polymer interaction (interface) which is expressed by the Winkler modulus. Thus, it is concluded that the graphene mode of failure in the embedded state differs dramatically from that observed in the air. Hence, significant differences in the values of the half-waves determined for the two media (air and polymer) are expected in the case of graphene flakes.K, of PMMA (in its various forms) and graphene.Finally, in order to obtain an independent value for the Winkler modulus we have performed calculations of the interaction between monolayer graphene and PMMA within the framework of density functional theory (DFT). These calculations can be separated in two parts, (a) geometry optimisations of isotactic PMMA (i-PMMA) and syndiotactic (s-PMMA) helixes with a length of two rings each, and (b) rigid potential energy surface (PES) scans of various relative configurations of a PMMA monomer and coronene. The optimised structures are in good agreement with the experimental findings of Kumakix and y directions). However, if the dimension of the flake along the loading direction, x, is less than that required for efficient stress transfer then the applied strain values given by the beam formula are not necessarily attained at the positions from which Raman measurements are taken. The reason for this is that the stress/strain in an inclusion embedded in a polymer matrix is built through shear at the interface. As found already by a number of authors, in order to transfer efficiently the applied stress or strain from the matrix to graphene a transfer length of the order of 1–2 μm is needed−1/%, through the formula . Also the above formula is only valid up to the first inflection point as failure processes that are triggered at that point will also affect the shift of the Raman wavenumber.The graphene samples were prepared by mechanical exfoliation of HOPG using the well-established scotch-tape methodmula see : The meal = 30 μm) and another (given in l = 4 μm), which corrections need be applied to. The results in th degree polynomial which accurately captures both the slopes at zero and critical strains and the slopes near the origin. The critical strain value for buckling shown at the plateau corresponds to the point of zero slope. For the case of In given in  with a stL) and the initial slope is – as expected - 38 cm−1/% i.e. outside the bounds of the required value of 60 cm−1/%. As seen the corrected data through formula (1) yield a critical value of compressive failure of ~−0.45% for this flake. A full collection of the data obtained in this work and in Ref. In contrast, the results of A plate with low flexural rigidity fails under compression in air by elastic buckling. When the plate is supported/embedded, out-of-plane buckling is restricted by the presence of the surrounding material which, in effect, provides a strong support against any buckling instability. The result is that the critical stress for buckling of the plate is orders of magnitude larger than that of the free plate. Analogous phenomena have been observed in embedded microtubules18l/w of the rectangular flake. As shown the data fluctuate around a mean value of about −0.6% regardless of aspect ratio. With reference to experimental scatter, this increases considerably for values of l/w close to the origin (very short flakes) for which certain corrections were made based on the initial slope of the Raman wavenumber vs. strain line. The scatter reflects the difficulty in measuring the graphene strain at that region due to the inefficient stress transfer as mentioned above. However, there is no doubt that the average value even for very short flakes is within the range of the mean value obtained from all the data points of the graph.In rp is the reaction pressure of the polymer matrix, wK is the Winkler's foundation modulus in units of stress per unit deflection (N/m2/m), and u is the deflection in the z- direction. We note that nonlinear frameworks can also be found in the literature for modelling the substrate effect39A widely used method for modelling plates resting on elastic foundation is that of Winkler'sA, of the compressed plate which is given byE is the total energy of the system- assumed zero at the point of failure-bU is the plate bending energy, fU is associated with the elastic strain energy as the flake is pushing/pulling against the surrounding matrix and T is the axial compression energy released by the flake buckling. The corresponding expressions for the terms of xN is the compressive force per unit length applied in x-direction, D is the bending stiffness and v is the Poisson's ratio of the plate and wK is the Winkler modulus.Here, for the embedded graphene sheets, the flake is assumed to be pushing on both surfaces against linear springs as shown in wK is the spring constant that describes the foundation and the interaction between the plate and the foundation and w is plate's width. The boundary conditions that should be satisfied are x = u'y = 0u = u = 0, u'. For the out-of-plane displacement u we make the assumption6394041m, n are the half-waves of the buckling mode in the x and y directions, respectively. Such an assumption for the displacement seems to be reasonable even though we have no direct observation of how the graphene plate is deformed due to the intervening polymer layer. However, useful information can be extracted from simply supported flakes that confirm the sinusoidal nature of graphene buckling under axial compression  the second term is zero and therefore the problem reduces to the Euler buckling formula for a freely suspended (graphene) plate in air.By inserting m, is evaluated by equating the force expressions 2crε is insensitive to the ratio of l/w and retains a value of ~−0.6% for a wide range of sizes and axial geometries  interactions and the other describing the influence of the elastic deformations of the matrix itself at the vicinity of the interface. For each case we have ascribed a spring of certain stiffness. These two springs are considered to be connected in series and with the overall stiffness equal to the reduced stiffness of the two springs  and syndiotactic PMMA (s-PMMA) over a graphene substrate. These MMA monomer–coronene relative configurations for the PES scans are shown in i-PMMA and s-PMMA with trans-gauche and trans-trans backbone conformations, respectivelyli-PMMA chain  and a 32ls-PMMA chain  using the B97-D functional and employing the def-SVPIn order to get an independent estimate of the strength of the van der Waals interactions between the PMMA polymer and graphene we have performed calculations within the framework of density functional theory (DFT). The primitive potentials see , u1–4, wu1–4 from the PES scans, the parameter z corresponds to the transverse distance from the coronene plane of the monomer's carbon atom nearest to the plane. In the case of the composite potentials, the function U corresponds to the potential per unit area (Å−2).All of the potentials, i.e. the initial potentials from the PES scans and the final composite potentials, have been fitted to the same modified Lennard–Jones type potential of the general form: zeq, we have σ/zeq = (1/2)c1/6, and i-PMMA and s-PMMA helices  system.It is straightforward to find the non-linear spring stiffness per unit area: R1, R2 between either of the two monomers to the coronene plane. The resulting scan is shown in K are 21.77 and 40.17 GPa/nm, for s-PMMA and i-PMMA respectively. An independent evaluation of the second spring length is non-trivial. Thus, we have based our treatment of this term on experimental results while also taking advantage of our (DFT based) calculation for the first spring. Specifically, we form a family of curves  multiple wrinkling (buckling) on the imposition of an axial strain of approximately −0.6%. The agreement between experiment and the DFT work was satisfactory bearing in mind the limitations of the DFT methods.2) to ensure that the Raman wavenumber of the examined region corresponds to the stress-free values and no fluctuations of Raman frequency that would possibly indicate graphene wrinkling, were detected.Graphene monolayers were prepared by mechanical cleavage of HOPG (High Order Pyrolitic Graphite) and transferred onto PMMA bars. The PMMA bars have dimensions of 2.9 mm thickness, width of 12 mm, length of 10 or 12 cm and covered on the top by a ~200 nm thick layer of SU8 photoresist . The graphene samples were first located using an optical microscope and monolayers with the desired dimensions were chosen for testing. The number of layers was identified with Raman measurements. Finally a layer of PMMA  was spin coated on the top with 6000 rpm. It is noted that in the previous work the samples were spin coated with another polymer S1805 photoresist (Shipley). Graphene and the polymer matrix do not chemically interact so the results are not influenced and no change in the Raman measurements was observed. For compression measurements on an embedded graphene it is of paramount importance for the flake or the region investigated to rest absolutely flat on the polymer substrate. Bearing in mind that the surface of polished polymers might exhibit a roughness of at least tens of nanometers, care was taken to map carefully the area under investigation . The reason for switching from cantilever to 4-point-bend testing is to allow a much better control of the strain imparted to graphene regardless of its position on the beam and to thus induce a higher level of applied strain than previously. The measurements were recorded with an excitation wavelength of 785 nm and laser power ~1 mW. More details can be found in references cOS = 0.17 and cSS = 1.75ζ def2-TZVPP3037All-electron density functional theory calculations including dispersion corrections (DFT-D2) were performed using the generalized gradient functional B97-D3233−7 au), the one electron density (rms of the density matrix up to 10−7), as well as, the norm of the Cartesian gradient (10−4 au). All of the DFT calculations were performed using the Gaussian package. The SCS(MI)-MP2 calculations were performed using Turbomole program package.Tight convergence criteria were enforced on the SCF energy (10Project planning, C.G.; sample preparation Ch.A., K.S.N.; spectroscopic measurements and analysis, Ch.A., O.F., G.T., J.P., K.P.; data interpretation and modelling Ch.A., D.S., E.N.K., N.P., C.G.; manuscript drafting, C.G., Ch.A., E.N.K., D.S.Supporting Information


The interface between organic semiconductor and graphene electrode, especially the structure of the first few molecular layers at the interface, is crucial for the device properties such as the charge transport in organic field effect transistors. In this work, we have used scanning tunneling microscopy to investigate the poly (3-octyl-thiophene) (P3OT)-graphene interface. Our results reveal the dynamic assembling of P3OT on single layer graphene. As on other substrates the epitaxial effect plays a role in determining the orientation of the P3OT assembling, however, the inter-thiophene distance along the backbone is consistent with that optimized in vaccum, no compression was observed. Adsorption of P3OT on ripples is weaker due to local curvature, which has been verified both by scanning tunneling microscopy and density functional theory simulation. Scanning tunneling microscopy also reveals that P3OT tends to form hairpin folds when meets a ripple. Beside its assembling on flat terraces, we pay special attention to the influence of unique characteristics of chemical vapour deposition (CVD) grown graphene, to name a few, the different facets and step edges of copper substrate, ripples and wrinkles along the boundaries between crystalline and noncrystalline areas or on step edges of crystalline areas1214In this work we focus on the interfacial structure of poly(3-octylthiophene) (P3OT) on a single layer graphene grown on multi-crystalline copper foil (SLG-copper). The assembling behaviour of P3OT at the solid-liquid interface was monitored In this work we performed STM characterization at the solid-liquid interface. P3OT was dissolved in phenyloctane with a concentration of 0.1 mg/g and a droplet of this solution was deposited on SLG-copper. As shown in About one hour after deposition, P3OT can cover almost the entire surface of SLG-copper, even randomly adsorbed second layers can be observed . In the According to the literature, after annealing of the polycrystalline copper foil, low index facets, Cu (111) and Cu (100) are dominant facets121312trans conformation, while cis conformation is a prerequisite for forming folds. When adsorbed on the ripple, the inter-chain distance is significantly larger, (2.63 nm vs 1.70 nm), suggesting that the inter-chain interaction is weaker. This is consistent with the observation that P3OT molecules adsorbed on ripples are unstable, and also in accordance with the previous observation of enlarged inter-chain distance on the curvature surface of single wall nanotubes20During CVD growth of graphene on copper foil, defect-like wrinkles and ripples tend to develop due to strain relief either along the boundaries between crystalline and noncrystalline areas or on step edges of crystalline areas8910et al. have attributed the observed compression to the epitaxial effectIn the high resolution images, periodic modulations along the P3OT backbone can be clearly distinguished . After cThe experimental results showed that the interaction of P3OT on ripples is weaker than that of P3OT on planar graphene. To verify the intrinsic interfacial interactions between P3OT and graphene or ripples, the adsorption capacity of a model oligomer of poly P3OT on a single layer graphene and ripples were respectively studied using the first principle calculations. For the calculations, the model molecule, P3MT composed of four 3-methyl thiophenes were used to model the P3OT, a piece of graphene with 370 carbon atoms was adopted to simulate the graphene layer, and single-walled carbon nanotubes  and SWCNT-30) were constructed to simulate the different curvature surface of ripples. All the models are large enough to avoid the edge effect.The calculation results showed that all the most stable conformations of the P3MT on graphene are obtained with sulfur atom above the centre of a six-carbon ring of the graphene or SWCNTs shown in  due to tTo summarize, we have investigated the assembling behaviour of P3OT on the SLG-copper surface. At the solid-liquid interface the adsorption of P3OT on SLG-copper appears quite dynamic. The orientation of P3OT domains is governed by the epitaxial adsorption of P3OT on the three-fold symmetric graphene lattices, while the underlying copper lattice has neglectable effect. The P3OT chains can across even the multi-atomic step edges of underlying copper lattices, however, tend to form hairpin folds when met a ripple along the step edge. The adsorption of P3OTs is weaker on the corrugated ripple surface, even the inter-chain distance is larger, which means that the local curvature also affects the molecule-molecule interaction. The epitaxial adsorption and effect of local curvature were both confirmed by our DFT simulations, which indicates the π-π stacking and polarization play decisive roles in the polythiophene-graphene interaction. The information obtained in this work is relevant to the understanding of graphene-semiconductor interactions which is crucial in many technologically important applications such as molecular electronics, organic solar cells and field-effect transistors.n ~34000) was purchased from Aldrich, and used without further treatment. The graphene was prepared with 25 μm thick copper foils as substrate. The single layer graphene on copper foil (SLG-copper) used here was synthesized by ambient pressure chemical vapor deposition reported elsewhereoC for 10 mins to remove the native oxide and rinsed with deionized water. Then, loaded into fused silica tube and heated to 1035 oC under the H2:Ar = 50:450 sccm (standard-state cubic centimetre per minute), followed by annealing for 20 min. Subsequently, 0.1 sccm methane was introduced into the system for 20 min. At last, the furnace was cooled to room temperature under the same gas flow. The SLG-copper was used as prepared for STM substrate.The solvent, 1-phenyloctane (≥99%) and P3OT  of the solution of P3OT in phenyloctane (0.1 mg/g) on SLG-copper. STM measurements were performed under room temperature at the liquid-solid interface with a constant current mode . Detailed imaging conditions are given in the figure captions. Mechanically cut Pt/Ir (80/20) tips were used. The STM images of the adlayers were corrected for XY drift against SLG-copper lattice.All geometry optimizations of the P3MT adsorbed on graphene and SWCNTs were respectively carried out using the density functional theory (DFT) method with the SVWN functional282930How to cite this article: Jiang, Y. et al. The Assembling of Poly (3-Octyl-Thiophene) on CVD Grown Single Layer Graphene. Sci. Rep.5, 17720; doi: 10.1038/srep17720 (2015).


This use of a chemical waste product as a solid carbon source instead of commonly used explosive hydrocarbon gas sources for graphene synthesis has the dual benefits of converting the waste to a valuable product, and reducing pollution.We have developed a simple, scalable, transfer-free, ecologically sustainable, value-added method to convert inexpensive coal tar pitch to patterned graphene films directly on device substrates. The method, which does not require an additional transfer process, enables direct growth of graphene films on device substrates in large area. To demonstrate the practical applications of the graphene films, we used the patterned graphene grown on a dielectric substrate directly as electrodes of bottom-contact pentacene field-effect transistors (max. field effect mobility ~0.36 cm The CTP film was passivated with a Ni capping layer, then annealed under high temperature, low vacuum and a reducing atmosphere. The method produces multi-layer graphene on the top and bottom surfaces of the Ni layer. This result suggests that carbon atoms are released during pyrolysis of the CTP, then diffuse into the Ni layer during annealing, and precipitate to form graphene on both sides of the Ni layer upon cooling4344454643452/Si substrate in an electronic device. We demonstrated an organic field-effect transistor (OFET) as the first example of organic electronic devices using graphene electrodes based on direct patterned growth from solid carbon source on the device substrate without using any additional graphene-transfer process; the Ni capping layer was removed by wet etching, thereby leaving the graphene films on SiO2/Si substrate.Here we introduce a simple, scalable, transfer-free, ecologically sustainable, value-added method to synthesize patterned graphene electrodes for high-performance electronic devices by recycling coal tar pitch (CTP) as an inexpensive carbon feedstock. CTP as an industrial by-product of steel production and is mainly composed of a complex mixture of aromatic hydrocarbons and heterocyclics2 (500 nm)/Si substrates (2 cm × 2 cm) were treated with UV/ozone for 30 min to establish good wetting between CTP solution and substrate, and then a 20-nm thickness of CTP film was deposited on SiO2/Si substrates uniformly by spin coating . The thickness of CTP film was measured using ellipsometry and the film uniformity was confirmed using atomic force microscopy (AFM) (2 (10 sccm) gas, and low vacuum (~330 mTorr) at 900–1100 °C for 1–4 min to find the optimal conditions for graphene synthesis. After annealing, the samples were cooled to room temperature under the same Ar/H2 flow. A Raman spectroscopy system with excitation of 532 nm was used to evaluate whether CTP films had turned into graphene films512/Si substrate in various growth conditions, the Ni layer and the graphene films grown on top of the Ni layer were removed by simple immersion in aqueous FeCl3 solution for 1 min; the underlying graphene films remained on the SiO2/Si substrate. This process enabled direct growth of graphene films on dielectric substrates, and avoided damage to the graphene that could occur if the graphene were transferred using additional transfer process. In the average Raman spectrum (n = 2500 points), the D peak (~1350 cm−1) corresponds to defects in the graphene films. At annealing time of 4 min, the D peak intensity decreased as annealing temperature increased in the range of 900–1100 °C  . A 200-n–1100 °C . When an–1100 °C . The temDc is diffusion coefficient, Dco is the maximum diffusion coefficient at infinite temperature and QID is the activation enthalpy for the diffusion, R is the gas constant and T is absolute temperature. This equation is Fick’s first law of diffusion expressed in terms of an Arrhenius type. Based on this equation, increase in annealing temperature increases the amount of carbon that can dissolve in the Ni film, and therefore increases the amount of carbon precipitated to the Ni surface upon cooling, and increase in annealing time gives sufficient supply of carbon atom to form high-quality graphene at high temperaturewhere The effects of various softening points (SPs) and concentrations of CTP solution were investigated, because these factors also affect the graphene quality. As the SP of CTP increased (due to an increase in the portion of large molecules in CTP solution), the film thickness increased at the same concentration (8 wt%) and spin coating condition  . The num2/Si substrates to measure the average Raman spectrum. We used a poly[methyl methacrylate] (PMMA)-layer-supported graphene transfer process with metal etching solution (FeCl3) described elsewhereFurthermore, CTP-derived films on top of the Ni layer were transferred onto SiO2 gas. Graphene films formed even under low-vacuum condition with Ar atmosphere (~3.8 mTorr), although the average Raman spectrum had higher D peak intensity  indicate that of mono-layer to few-layer graphene had formed. This synthesis method does not use any explosive gases (CH4 and H2), and can be used to mass-produce graphene easily and safely.We also obtained graphene films from the growth process without using Hntensity  than did2/Si substrates at 1100 °C for 4 min were prepared for use in further characterization. This annealing condition was chosen based on the results of preliminary trials. In the Raman scanning over a large area , pronounced peaks occurred at ~1580 cm−1 (G peak) and ~2700 cm−1 (2D peak), in addition to the D peak51D to G peak intensity IG ~0.1)  , then transferred to the TEM grids. TEM images at the folded edge of graphene films show that the multilayer graphene consisted of mono-layer to few-layer graphene with regular interlayer spacing (~0.34 nm)  in agreement with our Raman analysis data 26CTP-derived graphene films grown on the SiOIG ~0.1) . The qua1) (D/IG . On most(I2D/IG) . Raman mportions , and thasis data . Hexagonsis data  and elecsis data , inset26The graphene growth using Cu catalyst was also investigated. The Ni capping layer was replaced by a Cu capping layer (200 nm) and the samples were annealed at same growth conditions but at 1000 °C for 4 min because the melting point of Cu (1085 °C) is lower than that of Ni (1455 °C). After annealing, the Raman spectrum of Cu surface did not show any peaks related to carbon . After Cu etching, the Raman spectrum of the film that remained on the substrate had a large D and G peak without 2D peak; this spectrum is associated with amorphous carbon . Because the Cu is a surface catalyst due to the low solubility of carbon in it, the Cu layer did not act as an efficient catalyst to form a graphene when our metal capping layer structure was used432 gas, 0.2 Torr, 10 s) to remove the CTP film from regions not covered by the Ni pattern. During this process, the patterned Ni layer functions as a passivation mask that protects the CTP film underneath the Ni pattern from RIE treatment. After applying this method, Raman spectrum of the gap between the Ni patterns did not reveal any peaks related to carbon materials ; this means that RIE treatment removed the unprotected CTP film effectively. After etching the Ni patterns, the patterned graphene source/drain electrodes (gaps ~100 μm) can be obtained directly on the 500 nm SiO2/Si substrate. Finally, a pentacene layer (50 nm thick) was deposited on the patterned graphene electrodes at 50 °C through a shadow mask to complete OFET device fabrication.We used the multi-layer graphene films obtained from our transfer-free method to fabricate bottom-contact OFETs  based oninset) were characterized by measuring their output and transfer characteristics . We calculated RC of Gr-P and Au-P FETs by using the transfer line method with channel lengths of 30, 50, 80, and 100 μm . As gate voltage varied from −60 to −150 V, RC of the graphene electrode, normalized by channel width (1500 μm), decreased from 0.14 MΩ·cm to 0.043 MΩ·cm (545556μFET in the saturation regime was an order of magnitude higher in Gr-P FETs (0.05–0.13 cm2·V−1·s−1) than in Au-P FETs (0.011–0.017 cm2·V−1·s−1). The transfer curve of the Gr-P FETs showed a high on/off current ratio (1.1 × 107) with small hysteresis . In this case, the fabrication process was the same as used to fabricate the Gr-P FETs, except for the use of the large area substrate. The distribution of the μFET of the Gr-P FET large-area arrays showed ~95% operation and maximum μFET 0.13 cm2·V−1·s−1 (average μFET ~0.07 cm2·V−1·s−1)  , inset weristics . For comt levels . The Gr-43 MΩ·cm , which icm 545556. Calculasteresis , so theysteresis , inset. V−1·s−1) .RC in Gr-P FETs. RC between the electrode and the organic channel is primarily affected by the magnitude of the carrier injection barrier at the interface54WF of each electrode and electrode-pentacene interface. The calculated hole injection barrier of graphene electrode was 0.43 eV, which is 0.23 eV lower than that of Au electrode (0.66 eV), even though the WF of the graphene electrode (4.32 eV) was lower than that of Au electrodes (4.60 eV)  . We attr2/Si substrates that had graphene and Au electrodes patterns were immersed in an OTS SAM solution (0.1% in toluene) for 3 min, then dried under blowing using N2 gas. The OTS-treated Gr-FETs showed increased drain current and μFET (0.25–0.36 cm2·V−1·s−1), and the OTS-treated Au-FETs were also affected (μFET 0.04–0.07 cm2·V−1·s−1)  .μFET, ~0.36 cm2·V−1·s−1) in a bottom-contact geometry than did Au-electrode pentacene FETs. Using the same method, we also fabricated large-area graphene-electrode pentacene FET arrays (144 devices on a 4-inch wafer) to demonstrate transfer-free fabrication of graphene-based electronic devices. Our results provide a promising way to use graphene films converted from the carbon waste (CTP) as a reliable conducting material and as an alternative electrode for electronic devices. Furthermore, using this industrial waste product of steelworks to produce graphene will contribute to a clean environment and sustainable graphene growth.In conclusion, we demonstrated a simple, scalable, low-cost, ecologically sustainable, value-added process to synthesize graphene films from inexpensive CTP films under a Ni capping layer without using explosive hydrocarbon gas sources and physical transfer process. We synthesized multi-layer graphene films underneath the patterned Ni layer in a tube furnace and then achieved patterned graphene films as source/drain electrodes directly on the OFET device substrate without physical transfer process after removing the patterned Ni layer. The pentacene FETs using graphene source/drain electrodes synthesized using our transfer-free patterned growth method showed higher carrier mobility (max. 2 (500 nm)/Si substrates (NAMKANG HI-TECH). After baking at 240 °C for 30 min, the Ni layer (200 nm) was deposited  on a coal tar pitch film at room temperature. Then the sample was annealed using thermal CVD  with flowing 50 sccm Ar and 10 sccm H2 at a total pressure of ~0.3 Torr, then cooled to room temperature by removing the tube from the furnace while maintaining the same Ar/H2 flow. The Ni layer was etched away by dipping the sample into FeCl3 solution (Iron(III) chloride solution 45°Be’, JUNSEI Chemical) for 1 min and the sample was rinsed with DI water. Raman spectra and mapping images of graphene films were investigated using a Raman spectroscopy (WITEC Alpha 300R Raman spectroscope equipped with a 532 nm diode laser). The CTP-derived graphene was further characterized by HR-TEM .The coal tar pitch  was diluted with quinoline solvent  to a concentration of 8 wt%. The Coal Tar Pitch solution was spin coated  on a SiOA 50-nm-thick pentacene (Sigma Aldrich) layer was deposited on the electrodes through a shadow mask at 50 °C. The electrical properties of the graphene and Au-electrode pentacene FETs were characterized using a parameter analyzer (Keithley 4200). The work function of each electrode and electrode-pentacene interface was measured using ultraviolet photoelectron spectroscopy (UPS) .How to cite this article: Seo, H.-K. et al. Value-added Synthesis of Graphene: Recycling Industrial Carbon Waste into Electrodes for High-Performance Electronic Devices. Sci. Rep.5, 16710; doi: 10.1038/srep16710 (2015).


The intriguing structures have several advantages for straightforward fabrication on patterned substrates, high-performance graphene sensors without post-annealing process.We introduce a high-performance molecular sensor using self-corrugated chemically modified graphene as a three dimensional (3D) structure that indicates anisotropic charge distribution. This is capable of room-temperature operation, and, in particular, exhibiting high sensitivity and reversible fast response with equilibrium region. The morphology consists of periodic, “cratered” arrays that can be formed by condensation and evaporation of graphene oxide (GO) solution on interdigitated electrodes. Subsequent hydrazine reduction, the corrugated edge area of the graphene layers have a high electric potential compared with flat graphene films. This local accumulation of electrons interacts with a large number of gas molecules. The sensitivity of 3D-graphene sensors significantly increases in the atmosphere of NO Specifically, graphene-based gas sensors can detect individual molecular levels of gas adsorption3437891114151617Molecular detection using graphene is a promising method for use in gas sensors and biosensors because of its outstanding electrical properties and large specific surface area12 gas molecules. Moreover, the sensitivity of the gas sensor increased significantly at various NO2 gas concentrations.Thus, the goal here is to develop an easily fabricated, high-performance rGO molecular sensor capable of room-temperature operation, and, in particular, exhibiting high sensitivity and reversible fast response with saturation region. For this purpose, we introduce self-corrugated, chemically modified graphene as a 3D structure with anisotropic charge distribution. The accumulated electrons on the corrugated-edge area of the graphene layers exhibit a high electric potential compared with two-dimensional rGO (2DrGO) films. This local accumulation of electrons interacts with a large number of gas molecules, resulting in a reversible faster response and higher sensitivity than with conventional rGO sensors. The density of states of the self-corrugated graphene surface was confirmed by Scanning Kelvin Force Microscopy (SKFM), and the gradient of electric potential as a function of wall height on the 3DrGO was modeled by computational simulation as a function of anisotropic charge distribution. The field-enhancement factor is five times higher than that of 2D graphene film, and a fast response with saturation of resistance change was confirmed for the adsorption and desorption of NO2 thickness of 300 nm for the fabrication of a resistive-type gas sensor. In order to produce a well-aligned, corrugated rGO structure, the GO film was placed in a prepared chamber with a relative humidity of 80%. The rapid evaporation of toluene reduces the temperature at the air–solution interface because the toluene is highly volatile and immiscible with water vapor. Thus, condensation and nucleation of micro-sized water droplets results on the substrate. Rapid evaporation and airflow across the surface of the ODA-functionalized GO solution led to convective flows in the solution; the water droplets then created regular corrugation by packing of GO nanosheets. The GO nanosheets encapsulate the water droplets and precipitate at the water–solution interface. Following the evaporation of water from the GO nanosheets, self-corrugated regular vertical GO structures were produced on their surfaces. It has been established that solvents with a low viscosity have a low surface energy and their droplets have a tendency to adopt a spherical shape on the surface of high-viscosity solvents to minimize surface energy2H4) at 80 °C overnight. The both materials are deposited on the same interdigitated structure. This process constitutes a simple means for the fabrication of resistive-type rGO-based devices with various flexible substrates. 2 at thickness of 300 nm. The length and width of interdigitated electrodes are 50 mm, respectively. The −1, with lower intensity ratio of I2D/IG than single layered graphene, confirming heterogeneous structureD/IG ratio of 3DrGO reveals a large number of structural changes or defects in the 3DrGO. However, the I2D/IG of 3DrGO is lower than that of 2DrGO in 2, ~284.6 eV), C–C/C–N , C–O , C = O , and O–C = O  bonds. With subsequent reduction by fuming hydrazine, the presence of oxygen functional groups was dramatically reduced. These results suggest that the structural changes of rGO were not affected by the functionality of oxygen moieties. ds–Vds of the interdigitated rGO electrodes for both 2DrGO and 3DrGO. The resistance of 3DrGO (37 Ω) is slightly higher than that of 2DrGO (29 Ω), which agrees with the Raman analysis of structural changes of rGO. Although the morphology is significantly different, the resistance is only slightly altered because of the similar quantities of oxygen functional groups as insulating sites and the presence of nearly flat rGO films on the bottom layer of 3DrGO, as shown in Structural changes, such as mechanical deformation, wrinkles, doping, and the number of layers of rGO, can be described by Raman spectra, as shown in 2 substrate with respective Cr and Au source and drain electrodes of 30 mm length and width. ds–Vgs characteristics for both samples. It is clear that the strain and mechanical defects lower the on/off ratio due to the delocalization of electrons. In addition, the threshold voltage is similar for both samples. However, the on current of 2DrGO is higher than that of the 3DrGO, while the on/off ratio is lower. This is caused by mechanical deformation of graphene surfaces such as corrugation, breakage, and wrinkles29There are several issues involved in producing high-performance molecular sensors. A large surface area with corrugated surfaces facilitates the adsorption of large amounts of gas molecules and accelerated charge accumulation, which can improve gas detection in terms of sensitivity and fast-response time. In particular, the charge transfer induces significant changes in resistance due to the relation between electrostatic potential and charge density which can be described by Coulomb interaction. Also, the corrugated surfaces induce an inhomogeneous charge distribution. For instance, electron charge density increases on the rippled area of graphene surface2 gas concentration of 20 ppm. The gas sensitivity (S) is defined as S = ΔR/R0, where R0 is the initial resistance when exposed to pure N2 gas, and ΔR is the resistance change after exposure to the N2 and NO2 mixture. After exposure to NO2 gas at 20 ppm for 300 s, the resistance decreases due to the p-type property of rGO and the electron acceptor of the NO2 molecules0 × 100 (%). The sensitivity of 3DrGO (65%) is roughly three times greater than that of 2DrGO (20%). 2 gas concentration for 20, 50, and 100 ppm; the sensitivity increases linearly with increasing gas concentration. An extrapolation to lower gas concentrations predicts a sensitivity of about 0.2 at 1 ppm. This value is three times higher than that predicted for 2DrGO. The observed sensitivity values indicate that resistance increases with NO2 adsorption. Since the NO2 molecule is electrophilic, charges are expected to transfer from rGO to the physisorbed NO2, which confirms the p-type characteristic of rGO shown in 2 gas exposure can be separated into three regions, namely, (i) a sharp initial in gas response (fast-response region), (ii) a nearly linear intermediate region (slow-response region), and (iii) a region in which the sensor completely saturates (saturation region), as shown in 2 gas exposure, the first and second cycle are stable. The base-line is saturated as shown in 2 carbon and gas molecules, while the slow response is due to mechanical defects or oxygen functional groups, as described earlier2 sites. Conversely, the 3DrGO exhibits a reversible response with the saturation region (iii). This indicates that a certain limit adsorption of gas molecules over the surface could be present. The excess NO2 molecules that did not reach the surface active sites of the sensor could then be expected to remain in the surrounding region. This establishes a constant gas response. These results demonstrate that a gas molecules interacting with high-energy binding sites during are adsorbed continuously during gas exposure. Meanwhile, the adsorption of gas molecules were saturated by accumulated gas molecules with fast response owing to high electric potential on the graphene.The response curves for NOWe introduce self-corrugated, chemically modified graphene as a 3D structure that governs anisotropic charge distribution. This morphology can be formed by condensation and evaporation of water on a GO solution using the “breath figure” technique. In particular, the accumulated electrons on the edge sites of the corrugated graphene layers exhibit a strong electric field compared to flat graphene films. The sensitivity, stability, and fast response of 3DrGO were significantly improved in terms of gas-molecule detection. This method is a promising candidate among practical molecular sensors and flexible devices using chemically modified graphene without post-annealing process.2SO4 are mixed in a flask. Subsequently, 60 g of KMnO4 is slowly added for 1 h. For safety, stirring was carried out for 2 hours in ice bath. Again, the mixture is stirred vigorously for 24 hours at room temperature. DI (deionized) water is added and stirred for 10 min. A volume of 50 mL H2O2 (30 wt% aqueous solution) was then added, and the mixture was stirred for 2 h at room temperature. The mixture was precipitated and filtered to obtain the graphite oxide powder. This was then exfoliated into GO nanosheets in deionized water by homogenizer for 1 h2 substrate with thickness of 300 nm. The interdigitated electrode was fabricated by simple shadow mask with deposition of metal source such as Cr and Au. The channel width and length of metal electrodes were 50 μm. The Cr/Au electrodes were deposited using thermal evaporator at 2 × 10−6 torr with a thickness about 20 nm and 100 nm respectively. In order to produce the regular structure of 3DrGO, humid air which has a relative humidity of 80% and a flow rate of 0.1 L/min was supplied from a nozzle affixed to a funnel. The vertical distance between the substrate and the end of the funnel was kept at 6 mm inside a desiccator. After complete evaporation of the solvents, vertically aligned patterns of the ODA-functionalized GO were obtained. Subsequently, the film was transferred to a furnace to eliminate the ODA from the GO. Pyrolysis was carried out at 250 °C for 1 h under atmospheric pressure. The self-corrugated rGO with uniform array structures were achieved after fume reduction using hydrazine monohydrate (N2H4) at 200 °C for 24 h. The 3DrGO gas sensor could be fabricated with high uniformity as shown in Graphenen Oxide (GO) was obtained by natural graphite  according to a modified Hummers method−1 for Stokes shift. To confirm the alteration in the atomic ratio of carbon to oxygen with presence of functional group with respect to the 2DrGO and 3DrGO samples after fuming reduction by hydrazine, the X-ray Photoelectron Spectroscopy (XPS: Thermo VG Scientific Inc. MultiLab2000) analyses were carried out on an X-Ray photoluminescence spectrometer with an Al cathode as the X-ray source with power of 150 W. The fitted peaks of XPS spectra were determined by considering a combination of Gaussian and Lorenzian distribution. The Ids-Vds and Ids-Vgs characteristics were measured using a two-probe method (Keithley 4200-SCS). The total resistance with respect to 2DrGO and 3DrGO devices with interdigitated electrodes was measured by Ids-Vds. In order to measure the gating effect as depending of structures in 2 gas was mixed with N2 gas. The concentration of NO2 gas with respect to 20, 50, 100 ppm was controlled by Keithley 2700 multimeter with solenoid valve system. The resistances of sensors were measured by Keithley 2000 multimeter, recording electrical resistance as a function of time for gas exposure. Two-probe DC resistance measurements was used by Au wires for current and voltage leads. The measured values were carried out under atmospheric pressure.The morphologies of rGO were obtained by field emission scanning electron microscopy . The optical image of interdigitated electrodes for gas sensor was also revealed by optical microscope (Nikon Eclipse LV100). The structural and chemical characteristics in the 2DrGO and 3DrGO sheets were investigated by confocal Raman spectrometer  with an excitation wavelength of 532 nm and Rayleigh line injection filter with a spectral range of 100–3600 cmHow to cite this article: Yol Jeong, S. et al. Enhanced response and sensitivity of self-corrugated graphene sensors with anisotropic charge distribution. Sci. Rep.5, 11216; doi: 10.1038/srep11216 (2015).