The EB process is found to depend on both the graphene type and on the ambient conditions. For the mechanically exfoliated graphene, performing EB under vacuum leads to a higher yield of nanometer-gap formation than working in air. Conversely, for graphene on SiC the EB process is not successful under vacuum. Finally, the EB is possible with turbostratic graphene discs only after the creation of a constriction in the sample using lithographic patterning.Graphene-based electrodes are very promising for molecular electronics and spintronics. Here we report a systematic characterization of the electroburning (EB) process, leading to the formation of nanometer-spaced gaps, on different types of few-layer graphene (namely mechanically exfoliated graphene on SiO The vision of molecular electronics is the exploitation of single molecules as the active units in complex devices offering novel functionalities beyond the present technology –3. To ac2 carbon bonds assure a high mechanical stability even above room temperature. Finally, the possibility to exploit specific functionalizations to attach the molecular units to the graphene electrodes through carbon bonds and/or π-stacking seems to be a promising route to develop well-defined and robust junction configurations. In addition to these advantages, several theoretical papers have investigated the possibility to use graphene as an electrode to contact individual molecules [Recently, the use of graphene as electrode material for molecular electronics has been proposed ,10. Witholecules –16, predolecules , edge efolecules , and supolecules .2. Therefore it is important to test it also on other types of graphene. In particular, since many envisaged applications require the use of a gate electrode to tune the molecular junctions properties, it seems appealing the use of few-layer graphene, which is still thin but much less gate-dependent than the single layer [Recent works have successfully made use of graphene for the realization of electrodes in molecular devices ,17. Specle layer ,20.2 substrate and we show that the yield of nanometer-gap formation can be increased significantly when working at a reduced pressure. In addition, we report, for the first time, the EB process (in air and under vacuum) also on other types of few-layer graphene, which are better suited for large scale integration, namely multilayer graphene grown on the C-face of SiC [In this work we compare the EB process in air and in vacuum in few-layer graphene flakes exfoliated on SiOe of SiC  and thine of SiC –25.−4 mbar) so that the final comparison between the results in air and under vacuum is made independent on the geometrical factors. In total, we measured 23 graphene flakes of thickness of 1–20 layers, corresponding to 115 junctions.We first consider the case of few-layer graphene flakes obtained by the mechanical exfoliation technique . A typical flake is shown in the inset of I–V curves become strongly non-linear, probably related to the removal of contaminants induced by the current annealing [21To open a nanometer-size gap in the graphene devices, we polarized the junctions with a single voltage ramp with a fast feedback loop in order to stop the current immediately after the opening of the junction. We used the same method previously employed for the electromigration of gold nanowires . A typicnnealing ,20–2127.ealing 21. When thI–V measurement keeping the bias voltage below ±1.5 V to avoid any modification of the gap [After the EB process, the device is characterized by an  the gap ,28. An eI–V curves according to the Simmons model [d, the junction area A and the height of the tunnel barrier Φ as the fitting parameters. For the fit shown in d = 1.83 nm, A = 5 nm2, Φ = 0.6 eV. It is known that the value of d is a quite robust fitting parameter, while the other two are not, since good fits are still possible even with very different values. Indeed, we can fit our data with almost the same accuracy by using a very broad range of values for A and Φ  the values for d are in the range from 0.5 nm to 3 nm. This narrow distribution, which is expected since in the model the current depends exponentially on the width of the gap, is a strong indication that we can obtain electrodes with distances in the true nanometer range.In order to characterize the gap size in more details, we fit the ns model  (see SupI–V characterization after the EB process. The first class (A) is composed by the devices showing a sizeable tunneling current (I > 10−13 A) within the range of ±1.5 V. In these devices, the EB process created a gap with sizes of a few nanometers. In the second group (B) we consider junctions that underwent an EB process but no tunneling current is observed in our measurement range. These are “open” devices, with a gap larger than about 10 nm, which can be still useful for contacting long molecules or nanoparticles. Finally, the last group (C) comprises those devices that did not break under the maximum voltage/current we applied. The results of these studies are summarized in the upper panel of In the following, we discuss the yield of this EB process, with particular attention on the differences encountered between air and vacuum working conditions. In total, 58 junctions were processed in air and 57 under vacuum and they have been divided in three categories depending on the Considering the junctions processed in air, 25 of them (ca. 43%) exhibited a tunneling current after the EB, in 20 (ca. 35%) no detectable current was measured and finally 13 (ca. 22%) did not break during the process. Such a yield is comparable with what has been found by Burzurì et al. [V = 0.5 V in different devices. We stress that the I–V tests after the EB process were repeated also under different conditions (air or vacuum) and we did not find a significant dependence on the measurement conditions (only the conditions of the EB process are relevant). Since the tunneling current is inversely proportional to the size of the gap, we conclude that under vacuum conditions the EB is more controllable, leading to higher yield of success and generally smaller gap sizes.Interestingly, the tunneling currents measured after the EB process generally display a much higher value when the process is carried out under vacuum than in air, as shown in the bottom panel of It is also worthy of note that the number of devices that were not burned is almost the same for air and vacuum conditions. We noted that very thick or very large flakes led to devices displaying very small two-probes resistance (≤200 Ω), which did not undergo the EB process for our maximum applied voltage (10 V) and/or led to the saturation of our current-meter (30 mA). Indeed, the largest part of the devices falling in the category C of If we consider only the flakes that were effectively electroburned, i.e., neglecting devices falling in category C, the difference in yield between the EB process in air and in vacuum is even more evident: the rate of success (junctions displaying tunneling current after EB) is 55% in air and 80% in vacuum. In addition, the gaps obtained by the vacuum process generally display much larger currents implying a smaller size of the gap itself. Similarly, an improved yield in the fabrication of nanometer-sized gaps in graphene junctions by changing the environment conditions from air to vacuum was also reported by Nef et al.  who, howIbreak (upper panel) and the voltage Vbreak (lower panel) at which the EB process occurs as a function of the initial resistance of the junctions for both the working conditions (air and vacuum).In order to analyze further the dependence of the EB process on the environment conditions, in Ri while it is clearly higher when the process is performed under vacuum. The dependence of Ibreak on Ri, which is primarily determined by the number of layers of the flakes , suggests that the EB occurs at a constant current density (which only depends on the geometry of the device), regardless of the environment conditions, and in agreement with previous works [We found that the current is generally similar in the two cases (air or vacuum) and it always scales with the initial resistance. On the other hand, the voltage at which the device rupture is found does not show a clear dependence on us works . On the us works . We willI–V curve is shown in Vbreak > 40 V for vacuum processed devices, while it is in the range 10–15 V when working in air) and no detectable currents were ever measured after the EB process. Interestingly, the SEM images revealed that the rupture is always at the metal contacts and not in the graphene devices  in large quantities (hundreds of discs with a diameter of about 1–2 µm in a single deposition).We also tested the EB process in multilayered graphene microstructures, employing thin discs of turbostratically stacked graphene . These discs are comprise up to 100 graphene layers exhibiting a rather large charge carrier mobility in the range of 10ed discs . They caI–V-characteristics of such an electroburned device can verify the actual presence of a gap. Typically a bias voltage of 1 V results in a current of 80 pA, as exemplified in In a first attempt, simple pairs of source–drain electrodes have been defined on the discs by electron beam lithography. More than 10 discs have been contacted and underwent the EB process in air. In all of these experiments the EB took place at the contact areas as exemplary shown in 2 and achieve a better control on the gap formation.We now compare the different behaviors encountered in the three types of graphene analyzed. Concerning the exfoliated few-layer graphene flakes, we found that performing the EB in vacuum leads to a higher yield of nanometer-gap fabrication. At the same time the EB process happens at higher voltages under vacuum than in air, the current being almost identical under the two conditions. This means that the dissipated electrical power and the local temperature of the devices are similar or even slightly higher in the case of vacuum. Since a higher local temperature is usually considered to be detrimental for a controlled gap formation, we conclude that the environment pressure must be the key parameter to optimize the EB process of few-layer graphene on SiO2 vs graphene/SiC).Graphene grown on the C-face of SiC exhibits a similar behavior under ambient conditions, while under vacuum the oxygen–carbon reaction does not even initiate. This difference must be related to the different stacking feature (turbostratic vs graphite-like) of the graphene itself and/or to the different characteristics of the graphene/substrate systems (exfoliated graphene/SiO2 (which is known to be more reactive than SiC). Such impurities may be working as catalysts, initiating the EB process of graphene even under a reduced oxygen atmosphere.Focusing on these parameters, it is interesting to consider the case of turbostratic microdiscs, which have a morphology similar to that of graphene on SiC but are deposited on the same substrate as the exfoliated graphene. Here, we found that the EB process leads to a breaking of the graphene devices only after the patterning of a constriction. This hardness with respect to exfoliated graphene suggests that the different morphology of the edges also plays a role to initiate the burning. Indeed, the presence of nonsaturated carbon bonds makes the edges the most reactive part of the device. Edges cleaved during the exfoliation (exfoliated graphene), edges created during the oxygen plasma (graphene on SiC and turbostratic discs after the patterning), and edges of the untreated microdiscs  do have We presented a systematic study of the electroburning (EB) process in few-layer graphene devices for different graphene types. We focused our attention on exfoliated graphene, graphene epitaxially grown on the C-face of SiC and turbostratic graphene microdiscs comparing the results obtained when working in air and under vacuum. We showed how the process strongly depends on the specific type of graphene and on the environment pressure in the chamber. For mechanically exfoliated graphene, the vacuum process leads to the formation of nanometer-sized gaps with a higher yield and generally smaller sizes. On the other hand, for graphene on the C-face of SiC, the EB process creates a gap in the graphene devices when it is performed in air, while under vacuum conditions it simply leads to the blow-up of the metal contacts. As-deposited turbostratic graphene discs are found to be extremely resilient against the EB process: The opening of a gap in the device is possible only after creating a hot spot in the discs, as demonstrated in our work after the patterning of a constriction.Our work suggests that the oxygen pressure is a key factor in the EB process but also other factors such as the type of graphene stacking, the morphology of the edges and the specific graphene/substrate system play an important role. Further studies will focus on tuning the partial oxygen pressure during the electroburning to determine the optimal working conditions for the different types of device. We believe that these results will contribute to the realization of reliable graphene based electrodes for molecular electronics and spintronics.Few-layer graphene flakes were deposited by the standard “scotch tape” mechanical exfoliation method from natural graphite pieces on top of a p-doped silicon wafer coated with 300 nm of oxide. Flakes of suitable thickness (1 to approx. 20 layers) were located with an optical microscope on the basis of their contrast with the substrate. In some cases, the effective number of layers is also checked by micro-Raman spectroscopy, see Turbostratic graphene was obtained on on-axis SiC(000−1) semi-insulating wafer dice. First, the SiC dice were hydrogen-etched in order to obtain atomically flat surfaces . This prTo assure good ohmic contacts, the first fabrication step was the thermal deposition of 3 nm Cr/30 nm Au as the initial metal contacts. Successively, graphene was patterned in the desired device geometry  by electron beam lithography and oxygen plasma etching (30 s in a Diener Femto plasma system at maximum power). Finally, the remaining metal parts (the bonding pads and the connections from the contacts to the pads) were obtained by the evaporation of 10 nm Cr/100 nm Au and lift-off.−4 mbar). After the electroburning process, I–V measurements were taken with an AdWin-Pro system (16 bit output and input ) using a FEMTO pre-amplifier.For mechanically exfoliated and epitaxial graphene the electroburning process was performed by applying an increasing voltage ramp (20 mV/s) while continuously measuring the conductance of a junction. The burning of carbon atoms initiated the formation of a gap and increased the resistance. Once the resistance overcome the chosen value of 200,000 Ω, the voltage was immediately reset to zero (<100 µs). The process was performed at room temperature either in air or under vacuum ; the AFM characterization of the graphene grown on SiC; additional details about the fitting procedure of the non linear I–V curves according to the Simmons model.See Supporting Information for the Raman spectra of the different graphene materials (exfoliated, graphene on the C-face of SiC, turbostratic micro-discs deposited on SiOFile 1Titel: Material characterization.


Through careful control of the subtle balance between the linear periodic potential of rippled graphene and the C60 surface mobility, we demonstrate that C60 molecules can be arranged into a quasi-1D C60 chain structure with widths of two to three molecules. At a higher annealing temperature, the quasi-1D chain structure transitions to a more compact hexagonal close packed quasi-1D stripe structure. This first experimental realization of quasi-1D C60 structures on graphene may pave a way for fabricating new C60/graphene hybrid structures for future applications in electronics, spintronics and quantum information.We report the preparation of novel quasi-one-dimensional (quasi-1D) C Due to its many exciting chemical and physical properties, numerous applications have been demonstrated in single-electron transistors, superconductivity, and photovoltaics660 often serves as a model system for understanding molecule-surface interactions. C60 on epitaxial graphene, a system exclusively formed of elemental carbon, provides an ideal model system to study interactions at the molecular level, electromigration, molecular electronics and spintronics. On most metal or graphite surfaces, C60 molecules self-assemble into isotropic thermodynamically favorable hexagonal close-packed monolayers8910Buckminsterfullerene, C60 molecules is quasi-one-dimensional (quasi-1D) structures. Highly ordered quasi-1D molecular configurations are excellent model systems and prototypes of 1D quantum confinement of electronic states, and thus have potential importance in electronic nanodevices, spintronics and solid-state quantum computation12131460 nanostructures have been rarely realized experimentally due to their highly anisotropic configuration. A few previously reported experimental results include C60 peapods by embedding C60 molecules in carbon nanotubes, 1D C60 structures aligned along step edges on vicinal surfaces of metal single crystals, and C60 chains on self-assembled molecular layers161718192021222324252760 nanostructures on rippled graphene by utilizing the linear periodic potential in graphene as a template. Through carefully controlling the subtle balance between the weak periodic potential and the surface mobility of C60, C60 molecules on rippled graphene are arranged into a novel quasi-1D C60 chain structure with widths ranging from two to three molecules wide. At a higher annealing temperature, the quasi-1D chain structure transitions to a more compact hexagonal close packed quasi-1D stripe structure. To the best of our knowledge, this is the first experimental realization of quasi-1D C60 structures on graphene. Our experimental results may pave a way for fabricating new C60/graphene hybrid structures for future applications in electronics, spintronics , and quantum information.A particularly interesting arrangement of C60 molecules were deposited onto graphene grown by the CVD method on Cu foil at the background pressure below 2.0 × 10−8 mbar (more details are in the Method section). After post-deposition annealing at 423 K, we observed small atomically flat facets with typical sizes ranging from 50 nm to 200 nm on the surface. These facets can be classified into two predominant types based on the distinguishable surface characteristics, labeled as type A and type B  plus a trimolecular cell is 5.08 ± 0.02 nm. From these models, we can clearly see the large spacing (1.23 nm) between adjacent chains 60 on graphite based on the thermal desorption spectroscopy study60 chains and stripes are grown, show a well-defined linear periodic graphene rippled superstructure with an orientation close to the chain and stripe orientation. 60 chains and stripes, as indicated by the blue arrow in 30272827283334353637To gain detailed insight into the growth mechanisms of the quasi-1D C60 chain structure. Based on the rippled graphene structure, it is possible for us to explain the formation of the chain structure from the point of view of mechanical stability60 on flat graphene is the hexagonal close-packed structure due to the weak diffusion barrier on flat graphene and the much stronger C60-C60 interaction4160 to arrange into more stable quasi-1D configurations. The linear periodicity of rippled graphene, 0.75 ± 0.01 nm, is smaller than the typical inter-row distance of 0.87 nm in the C60 closed packed arrangement. C60 molecules randomly occupy the valley sites of the linear periodic potential at the initial stage of growth. It is possible for C60 molecules to settle down in the adjacent ripples at a less stable state because C60 molecules are forced to stay slightly off the valley sites. In this way, two or three adjacent linear graphene ripples are filled up. This is consistent with the observation that almost all C60 chains consist of two or three rows. Based on a geometrical and energetic argument, it is easy to see that chains with 4 or more rows are less likely to form. Compared with the formation of the bimolecular chains, the possibility to form trimolecular chains (both sides need a C60 molecule sitting down around the middle C60 molecule) is about half the probability of the formation of the bimolecular chains, which only need one side to be occupied with C60 molecules. From a Monte Carlo statistical approach, C60 molecules adsorbed on such a linear graphene sheet will be stabilized in one of the valleys60 molecules with less thermal energy can form monomolecular chains. That is one of the possible directions to further tune and control the formation of these quasi-1D structures.The orientation of the graphene ripples is fully consistent with that of the C60 stripe structure after annealing at 487 K, we will start by considering the diffusion of Cu atoms underneath graphene and the arrangement of C60 molecules on graphene. When annealing at a higher temperature of 487 K, C60 molecules have a higher probability to overcome the potential energy difference between the hilltop and the valley sites of rippled graphene and rearrange into chains with larger width due to the higher thermal excitation. The junction structure in the bimolecular and trimolecular chains probably provides diffusion channels for C60 molecules to transfer between adjacent chains resulting in emergent chains with larger widths. Considering that the arrangement of bimolecular and trimolecular chains is usually 2-3-2 or 3-2-2 as described above, it is reasonable to assume that the emergent chains with width of five C60 rows are the configuration with highest probability to form, compared with chains of other widths. The distribution of chain widths and stripe widths may imply that at the annealing temperature of 487 K, the thermal excitation can only drive C60 molecules to cross over one or two energy barriers, but not more. When adjacent bimolecular and trimolecular chains merge into a 5-row chain, the interchain spacing between the merged 5-row chain to its neighboring chains can be generated from the initial interchain spacing in the chain configuration (1.23 nm), which is 1.59 nm. This spacing is slightly smaller than the interchain spacing of 1.73 nm required for a new C60 row to fill in the empty space. It has been previously reported that Cu atoms can easily diffuse underneath a graphene sheet even at room temperature. These detailed observations suggest that copper atoms can diffuse along smooth step edges and kinked sites, and even dissociate from step edges, indicating a weak interaction between the graphene sheet and the underlying copper atoms60 row to form along the straight step edges and results in hexagonal close packed quasi-1D C60 stripes with the typical width of six rows. The origin of C60 stripes with other number of rows can be attributed to the emergence of different combinations of C60 chains, such as two bimolecular chains emerging into a stripe with five C60 rows.In order to understand how the bimolecular and trimolecular chains transform into the C60 chain and stripe structures  was loaded into the homemade Knudsen cell mounted on the load lock of the LT-STM system. The C60 source was degassed at a pressure of low 10−8 mbar before deposition. C60 molecules were then sublimated at a deposition rate of ~0.9 ML/min with the background pressure below 2.0 × 10−8 mbar. During the deposition process, the substrate was kept at room temperature. The samples were then subsequently annealed at 423 K, 487 K, and 578 K respectively in the preparation chamber of the STM system with the base pressure at low 10−10 mbar. All the STM measurements were performed at room temperature in the STM characterization chamber with a base pressure of about 3.0 × 10−11 mbar.All experiments were carried out in an ultra-high vacuum (UHV) scanning tunneling microscope system (Omicron LT-STM). Before CHow to cite this article: Chen, C. et al. Temperature Evolution of Quasi-one-dimensional C60 Nanostructures on Rippled Graphene. Sci. Rep.5, 14336; doi: 10.1038/srep14336 (2015).


There are over 0.8 million people living with dementia in the UK. The needs of people with dementia are increasingly complex as the illness progresses. Eating and drinking difficulties are a major source of ill health and stress across the stages of dementia in multiple settings.The evidence on what interventions support people with dementia in continuing to eat and drink well needs to be updated, and the full set of interventions assessed. To ensure people with dementia and their carers have access to the best current evidence, which address the questions that are important to them, we consulted with stakeholder groups and formulated specific questions to be addressed by a systematic review. The evidence will be summarised in light of these questions.http://www.crd.york.ac.uk/ PROSPERO/ DisplayPDF.php?ID= CRD42014007611) was shared with members of two patient and public involvement groups. The members were asked to comment on the protocol and suggest questions they would like the review to address. Compiled questions were added to the protocol to inform our search strategies.The review protocol (We have conducted a comprehensive search of 13 databases for studies that assess the effectiveness of interventions to improve, maintain or facilitate oral food and drink intake, nutrition and/or hydration status, in people with dementia. Screening studies for eligibility, assessing risk of bias and extracting relevant data is underway and duplicated independently. Studies included in the review will be used to address the specific questions asked.The stakeholders’ questions, that were not originally specifically addressed by the review protocol included issues around personalisation of interventions, the relationship with the carer, meaningful activity around food as well as specific issues around oral hygiene and swallowing difficulties. These questions are guiding the way that the review is conducted and will help focus the way that information from the review is provided to people with dementia and their carers.This systematic review aims to review the effectiveness of interventions to improve, maintain or facilitate oral food and/or drink intake in people with dementia and address specific questions raised by stakeholders. By soliciting and then addressing questions within the systematic review that are important to our stakeholders we will improve the usability of our findings - the questions and their summary answers will be used to inform best practice.


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).