Although silicon technology has been the main driving force for miniaturizing device dimensions to improve cost and performance, the current application of Si to soft electronics (flexible and stretchable electronics) is limited due to material rigidity. As a result, various prospective materials have been proposed to overcome the rigidity of conventional Si technology. In particular, nano-carbon materials such as carbon nanotubes (CNTs) and graphene are promising due to outstanding elastic properties as well as an excellent combination of electronic, optoelectronic, and thermal properties compared to conventional rigid silicon. The uniqueness of these nano-carbon materials has opened new possibilities for soft electronics, which is another technological trend in the market. This review covers the recent progress of soft electronics research based on CNTs and graphene. We discuss the strategies for soft electronics with nano-carbon materials and their preparation methods (growth and transfer techniques) to devices as well as the electrical characteristics of transparent conducting films (transparency and sheet resistance) and device performances in field effect transistor (FET) . In addition to discussing state of the art performance metrics, we also attempt to clarify trade-off issues and methods to control the trade-off on/off versus mobility). We further demonstrate accomplishments of the CNT network in flexible integrated circuits on plastic substrates that have attractive characteristics. A future research direction is also proposed to overcome current technological obstacles necessary to realize commercially feasible soft electronics. Since the invention of the transistor, the semiconductor industry has affected nearly every aspect of our daily life ,2. One mFigure Nano-carbons such as one-dimensional carbon nanotubes (CNTs) and two-dimensional graphene layers have been widely studied to open a new technology platform based on flexible electronics requiring high transmittance, bendability, and high mobility –12. Figu2 V−1 s−1 [2 V−1 s−1 [2 V−1 s−1 in suspended structures [9 A cm−2 [−1 K−1 [−1 K−1 [The extraordinary electrical, physical, and chemical properties of CNTs and graphene have been attractive since their discoveries. Both materials exhibit outstanding carrier mobility, which is attractive for applications to electronic devices. The carrier mobility in semiconducting single-walled carbon nanotubes (SWCNTs) has been shown to be as high as ~80,000 cm V−1 s−1 , while t V−1 s−1  on insulructures . These u9 A cm−2 ,17. At r [−1 K−1 , and gra [−1 K−1  with a h [−1 K−1 . In addi [−1 K−1 . For theThe purpose of this article is to summarize the recent progresses of both CNTs and graphene in soft electronics, and furthermore, to provide guidance for future nano-carbon research by clarifying feasible approaches which will most likely lead to soft applications. We first discuss several successful attempts to synthesize CNTs and graphene. Variations in transfer techniques for both materials are discussed thoroughly. For the use of CNTs and graphene for transparent conducting films (TCFs), the characteristics of TCFs using both nano-carbon materials are compared in depth, together with ITO. Furthermore, various types of field-effect devices using different forms of CNT FETs such as single CNT FET, random network CNT FET, aligned CNT FET, and different forms of graphene FETs such as single layer graphene (SLG), bilayer graphene (BLG), and graphene nanoribbon (GNR) are compared. Moreover, the specific FET device performances related to material preparation and fabrication techniques are also discussed. Finally, the logic level, flexibility, and stretchability of devices with a combination of graphene and CNTs along with their utilizations in logic circuits are further discussed. The systematic deep analyses of the device properties of graphene and CNTs highlight excellent opportunities for future flexible electronics. We conclude with a brief perspective on the research directions of soft electronics in future.The preparation techniques for CNTs and graphene are the most important fundamental research areas providing realistic applications. From the discovery of CNTs and graphene, diverse work has been done to improve the quality of the materials  and to control other parameters  and morphology . This section describes some of the most successful methods for synthesis of nano-carbon materials.The CNT synthesis techniques aim to provide control over the tube density, spatial distribution, length, and orientation. Controlling the tube diameter and ratio of semiconducting to metallic SWCNTs have been a critical issue for electrical applications ,23. The 2 substrate. The CVD growth was carried out on various catalysts, including Fe nano-particles and metal thin films  on Si wafers, quartz, and metal foils to synthesize CNT forest [Individual CNTs are horizontally grown on the substrate by CVD, as shown in Figure T forest ,34. DepeT forest . Highly-T forest ,37. EffiT forest ,39. FiguT forest ,41. SoluT forest –46. One T forest . Random T forest ,49. The T forest . NeverthT forest –53. FiguT forest .Since graphene was first electrically isolated from graphite using a mechanical exfoliation method, many efforts have been studied to synthesize thin graphene films such as the CVD method, reduction of graphene oxide (GO), epitaxial growth on SiC, and chemical molecular assembly method.As shown in Figure Most CVD approaches for synthesizing CNTs and graphene require high temperatures which prevent direct growth of nano-carbon materials on plastic and other soft target substrates. CNTs and graphene located on a catalytic substrate need to be transferred onto a target substrate. Transferring graphene from the metal substrates onto desired substrates without degrading the quality of graphene is the critical step to use CVD-grown graphene for most practical applications.3 or (NH4)2S2O8 are often used for removing Cu, and NaOH or KOH for sapphire [2/Si and metal films, still has problems with mechanical damage [2 gas is produced by a water reduction reaction, and the PMMA/graphene layer detaches from Pt substrate due to the H2 bubbles at the interface between the PMMA/graphene and Pt substrate. Damage of the mother-substrate is reduced considerably, and the substrate can be used repeatedly for the next CVD growth. In addition, the transferred graphene is free of metal particles, which are commonly found in graphene transferred by the metal etching process. Figure Wet etching processes are commonly used to detach as-grown materials from the mother substrates using chemical solutions. FeClsapphire ,80. The l damage . The roll damage ,83. The l damage . The tral damage . The PMMCommon electronic devices require conducting, semiconducting, and insulating materials. For conducting elements, several conducting polymers such as polyacetylene, polypyrrole, polythiophene, polyaniline, and poly poly(styrenesulfonate) (PEDOT:PSS) have been investigated for future applications to replace conventional rigid conducting and semiconducting materials ,92. HoweElectrical conducting materials would have potential for consumer applications, such as soft displays, energy generators, and human bio-devices. In such applications, metal oxides such as IZO and ITO are the most widely used materials –96. Howe3+ ions has also been shown to reduce sheet resistance to 110 Ω/sq at a transmittance of 90% [3 doping. This approach demonstrates not only improvement of sheet resistance and uniformity but also better environmental stability compared to topmost layer doping. The optimized LbL-doped four-layer graphene shows a sheet resistance of 54 Ω/sq and a transmittance of 85% (at 550 nm) with a robust bending stability. The performance of the graphene conducting films need to be further tuned and improved to meet different requirements of practical flexible products [During the past few years, much effort has been given in synthesizing CNT films as a conducting element ,97–99. Se of 90% ,110. Whie of 90% . Togethee of 90% ,113, duee of 90% –116. A ne of 90% . Each laproducts ,119. Botproducts . While Iproducts –122. Theproducts . The resMiniaturization is the most important issue not only to increase device integration density but also to improve FET performance for complicated operations. Semiconducting Si technology has given great contributions to society, but now faces scaling which involves heat and power consumption issues due to the fundamental limitations of Si. Atomic-thick nano-carbon materials might satisfy the scaling issue and give great benefits with combination of electrical/mechanical/optical advantages. As an active channel component, SWCNTs and graphene have been studied for fabricating FETs and p − n junctions to demonstrate their potential to outperform established materials for next-generation electronics –128. Her5), at room-temperature operation [2 V−1 s−1 [Figure peration –134. Witperation . For fabperation . The tub V−1 s−1 ,137. Fig V−1 s−1 . SWCNTs  V−1 s−1 ,139. How V−1 s−1 . In grap V−1 s−1 ,141–143. V−1 s−1 .Figure 7One of the key issues in high-performance TFTs is high on/off ratio for efficient switching behavior. In the case of a CNT channel, the as-grown CNT network usually contains both semiconducting and metallic CNTs . These m7 at room temperature [2 in a large sheet of graphene [3 hybridization with atomic H makes this approach impractical.One method to obtain high on/off ratio involves electrical thinning of the thick MWCNTs and CNT bundles, as shown in Figure perature –165. Simgraphene ,167. Howgraphene ,169. It graphene . Yet, agAlthough CNTs and graphene intrinsically have an ambipolar transport property, both show p-type behavior under ambient conditions due to contacts, doping by oxidizing acids, or doping by the adsorption of atmospheric oxygen molecules and/or moisture. It is important to control the carrier type of nano-carbon transistors for applying “complementary metal-oxide-semiconductor (CMOS) technology” because high noise immunity and low static power consumption are critical issues in the modern semiconductor industry. Therefore, it is desired to control the major carrier types of CNTs and graphene FETs by chemical and/or nonchemical doping methods. Here, we introduce several polarity control methods to modify the majority carriers in CNT- and graphene-based transistors such as chemical doping, oxygen doping, electrostatic doping, trap charge-induced doping, and metal work function engineering.4, and AuCl3) are also indicated as dotted lines. As shown in Figure 3+ ion has the large reduction potential of 1.50 V, which acts as p-type doping in CNTs. BV has an oxidation potential of −1.1 V, which implies that BV can act as an n-type dopants. BV donates electrons to the empty conduction band of semiconducting CNTs [2 physisorbs on the CNT surface [2 molecules [In order to have n-type conversion and p-type enhancement behavior in CNTs under ambient conditions, various chemical doping strategies have been investigated –179. Theing CNTs . The riging CNTs . Using βing CNTs . A reducing CNTs . For graing CNTs . Similaring CNTs . For BLGing CNTs . Figure ing CNTs . One of  surface ,188,189.olecules . It has olecules ,191. Figolecules ,192. Potolecules ,175,193.olecules . By varyNumerous efforts have been made to get higher on/off ratios and better control of carrier type in nano-carbon transistors. In order to understand advantages and disadvantages for CNT and graphene FETs, a side-by-side comparison is required. Table Next-generation military and industrial radio-frequency (RF) surveillance systems will benefit from flexibility and stretchability of circuits for increased resilience. A realistic short- and medium-term goal for carbon electronics is utilizing the combination of electrical, mechanical, and optical properties of CNTs and graphene thin films to replace organic semiconductors and a-Si in these flexible/stretchable systems –204. In DD) due to the zero band gap of graphene [Figure graphene . With angraphene –211. Figgraphene . A couplApplications ranging from flexible solar cells, displays, e-papers, wearable and biomedical skin-like devices open up new opportunities in the field of electronics. In this section, we describe applications of several flexible devices possible with carbon electronics jsuch as sensors, LEDs, RF devices, stimulators, and memory devices.−1, which is three times larger than previous NW-based sensors [in vivo optical recordings of signal imaging. In Figure As an example of further applications of flexible devices, Figure  sensors . Figure  sensors . Each pi sensors . Graphen sensors –221. RF  sensors . Noninva sensors . The fleStretchability is a key parameter in the development of wearable devices that can be embedded into clothes and garments or even attached directly to the skin, where high levels of strain will be encountered. Possible applications include the human-friendly devices for detecting human motions, monitoring health system, and healing. In addition to flexibility, all these stretchable applications demand tolerance of large levels of strain (> > 1%) without fracture or significant degradation in electronic properties. The mainstream strategy to realize improved stretchability focuses on the development of stretchable materials including organic polymers, networks of 1-D wires, and nano-carbons –231. Owi−1 at the strain of 150%. This property can be utilized to construct strain sensors, with performance comparable to conventional metal-strain gauges. Using a nonlinear buckling process as shown in the right panel of Figure y and Rx) were stable up to ~11% stretching with only one order of magnitude change at ~25% stretching. 3D-graphene macroscopic structures formed with a foam-like network of graphene was also developed using template-directed CVD  layer was deposited onto rough Cu foil using atomic layer deposition. After coating with PMMA, Cu foil was chemically etched, and the Al2O3 layer was then transferred as dielectric layer. This transferred Al2O3 layer was wrinkled with a “wavy” structure, which was robust under high tensile strain. The resulting TFTs exhibited device-acceptable electrical performance with small gate leakage current due to the build-in air gap between wrinkled Al2O3 and graphene gate. The devices were stretched along the length direction (16% strain) and along the width direction (20% strain), as shown in Figure Figure We have reviewed the current status of CNTs and graphene in diverse applications of soft electronics from material preparation to performance in logic circuits. Low-dimensional carbon materials exhibit superb electronic properties and promising performance and are attractive for future electronics. Methods for synthesizing one-dimensional CNT and two-dimensional graphene films, as well as procedures for device fabrication on soft substrates have been discussed here. Both CNTs and graphene exemplify TCF properties including a high operational flexibility and stretchability that are not accessible with transparent ITO electrodes. Likewise, field effect mobilities of carbon-based transistors have reached levels unfeasible by organic semiconductors/a-Si. CNT FETs, whether composed of a single CNT, aligned CNTs, or random network CNTs, show high on/off ratio and mobility. Graphene FETs provide extremely high mobility but poor on/off ratio due to zero band gap. Engineering for on/off ratio increase and carrier polarity control were summarized. For applications in active electronics, SWCNT and graphene transistors can be assembled on a variety of substrates including flexible plastic and stretchable elastomers. Various complex integrated circuits based on nano-carbon materials have been demonstrated in the literature, as well. Each of these topics requires significant future exploration in order to realize commercialized applications of the immense potential of nano-carbon in next-generation electronics.In spite of recent progress demonstrating the unique advantages of CNTs and graphene, the possible applications, social influence, addressable markets, and related economic issues will eventually decide the success of these nano-carbon materials. Both have unique and superb properties which open the possibility for soft electronics. Nevertheless, applications are limited by a different set of factors. Assemblies of CNTs are practical compared to the use of individual CNTs, but require the positioning of the CNTs in a specific direction, with desired density, and of desired metallicity/chirality. Methods to achieve this control are a current hot topic, but adoption of a particular method will require a high yield for industrial utilization even in niche applications. Conversely, graphene can be prepared in a large-area format. Yet, the transfer to a desired substrate may provoke damage in the graphene layer and degrade device performance. Therefore, developing a smart way of assembling CNTs to maximize the device performance and robust method of transfer of large-area graphene are two key ingredients that are unsolved but required for application. On a systems level, future electronics including biomedical applications with biocompatibility will require further research. For instance, CNTs and graphene combine synergistically, showing better flexibility and stretchability with no degradation of electrical performance when engineered to maximize potential. Additionally, combining both stretchable materials and stretchable geometries can allow for extremely stretchable systems. Aside from the engineering challenges of applying nano-carbon to soft electronics, CNTs and graphene are outstanding materials for demonstrating a number of basic science concepts in the fields of quantum electrodynamics, quantum optics, and quantum chemistry. Controlled synthesis and application of monolayer materials also allows exploration into a new class of vertical tunneling devices. Aside from carbon, other classes of graphene-like 2D materials such as transition-metal dichalcogenide (TMD) materials and boron nitride (BN), might also be promising in the field of soft electronics when a band gap or other electrical/mechanical properties are required. These related engineering opportunities in areas with the broad range of influential research topics provides strong motivation for continued efforts in human-friendly soft electronics.


Copper coordination polymers from cavitand ligands are reported including networked cage-motif structures, one of which takes up C60 from solution. I and/or CuII cations that show a remarkable range of framework topologies and structures. Complex [CuI4CuII1.5(L1)3(CN)6]·CN·n(DMF) 1 features a novel 3,4-connected framework of cyano-linked hexagonal metallo-cages. In complexes [Cu3(L2)4(H2O)3]·6(OTf)·n(DMSO) 2 and [Cu2(L3)2Br2(H2O)(DMSO)]·2Br·n(DMSO) 3 capsule-like metallo-cryptophane motifs are formed which linked through their metal vertices into a hexagonal 2D network of (43.123)(42.122) topology or a coordination chain. Complex [Cu2(L1)2(OTf)2(NMP)2(H2O)2]·2(OTf)·2NMP 4 has an interpenetrating 2D 3,4-connected framework of (4.62.8)(62.8)(4.62.82) topology with tubular channels. Complex [Cu(L1)(NCMe)]·BF4·2(CH3CN)·H2O 5 features a 2D network of 63 topology while the CuII analogue [Cu2(L1)2(NMP)(H2O)]·4BF4·12NMP·1.5H2O 6 has an interpenetrating -b type structure and complex [Cu2(L2)2Br3(DMSO)]·Br·n(DMSO) 7 has a 2D network of 4.82 topology. Strategies for formation of coordination polymers with hierarchical spaces emerge in this work and complex 2 is shown to absorb fullerene-C60 through soaking the crystals in a toluene solution.The cyclotriveratrylene-type ligands (±)-tris(iso-nicotinoyl)cyclotriguaiacylene L1 (±)-tris(4-pyridylmethyl)cyclotriguaiacylene L2 and (±)-tris{4-(4-pyridyl)benzyl}cyclotriguaiacylene L3 all feature 4-pyridyl donor groups and all form coordination polymers with Cu Coordination polymer materials that feature channels or cavities in their framework but are not robust to loss of all guest solvent may also find function as heterogeneous hosts provided guest molecules can be exchanged without substantial loss of framework structure. This principle has recently been used to great effect by Fuijta and others who have developed a crystalline-sponge approach for the determination of otherwise inaccessible crystal structures of guest molecules exchanged into the cavities of coordination polymer hosts.2Coordination polymers and metal–organic frameworks (MOFs) are crystalline coordination compounds with infinite framework structures constructed from metal cations and bridging ligands.2 capture.3 Other well-known classes of molecular hosts including calixarenes and related cone-conformation hosts, and cucurbiturils have also been reported to form coordination polymers, often when suitably functionalised with metal-binding groups. Cyclotriveratrylene (CTV) is a bowl-shaped molecular host with a tribenzocyclononatriene core. CTV forms chain and 2D coordination polymers, most commonly with group 1 metal cations.7 CTV analogues, including the demethylated cyclotricatechylene8 and a range of tripodal ligand-functionalised analogues of CTV have been shown to form coordination polymers that contain embedded molecular hosts. Some of these coordination polymer materials exhibit structures with large channels or cavities and where the molecular bowl of the tribenzocyclononatriene ligand scaffolds are potentially accessible to new guest molecules. In the majority of examples, however, the molecular recognition sites are not accessible due to inter- or intra-network host–guest interactions. These include bowl-in-bowl stacking motifs between the tribenzocyclononatriene scaffolds,11 and a pronounced tendency for the formation of self-complementary interactions between two CTV-type ligands to form a dimeric so-called hand-shake motif, or through inter-network host–guest interactions between a terminal additional ligand group and the tribenzocyclononatriene host core.The ability of coordination polymers to function as a host is a property of the overall assembly and not of the individual molecular or ionic building blocks that make up the framework. Molecules that have an intrinsic ability to bind guest molecules are also well known and are referred to as molecular hosts, and most examples are cyclic in nature. The use of molecular hosts as components of coordination polymers may result in a material which contains both the specific molecular recognition sites of the molecular host and channels and cavities of the coordination polymer framework. As such, they can be regarded as materials with a hierarchical pore structure. Cyclodextrins, for example, form robust metal–organic frameworks which have applications in gold extraction and COa,d,g]cyclononatriene core has been decorated with a tripodal arrangement of 4-pyridyl ligand groups, namely (±)-tris(iso-nicotinoyl)cyclotriguaiacylene L1,15 (±)tris(4-pyridylmethyl)cyclotriguaiacylene L2 ,15 however all three ligands are known to form discrete metallo-supramolecular species. Complexes with ligand L1 have been the most studied and include Pd6L8 stella octangula cage species, capsule-like metallo-cryptophane Pd3L2 species,19 and a Cu6L6 metallacycle formed from CuBr2 that has a topologically unique Borromean-like chainmail arrangement.20 The materials reported here represent a remarkable structural variation and include unusual topologies induced by the distinctive pyramidal shape of these ligands, coordination polymers incorporating cage or capsule-motifs with both large lattice channels and accessible molecular recognition sites, as well as new examples of more common topological types.We report here a series of new coordination polymer materials that have been accessed through the combination of copper cations and CTV-type ligands where the tribenzo[ylene L2  and (±)tI4CuII1.5(L1)3(CN)6]·CN·n(DMF) 1. The crystal structure indicates that there was air oxidation of some copper sites from CuI to CuII. Interestingly, we also obtained complex 1 in small quantities from reaction of L1 in DMF with CuII(NO3)2. The cyanide in the material obtained from Cu(NO3)2 is likely to have resulted from decomposition of the solvent, and it is notable that the IR spectrum of this material did not indicate the presence of any nitrate. Formation of copper cyanide based coordination polymers where the cyanide has come from reagent decomposition has been reported before from cleavage of diaminomaleonitrile.21 Mixed valence copper coordination polymers have also been previously reported, although examples are not numerous.22Reaction of L1 and CuCN in dimethylformamide (DMF) initially results in formation of a pale green viscous solution out of which small red-orange single crystals grow after approximately 3 weeks of standing. The red-orange crystals were of a 3D coordination polymer complex of composition + framework able to be structurally elucidated, counter-anions and any included solvent were not located due to disorder and subsequent weakly diffracting nature of the material. There are three crystallographically distinct Cu sites: a square planar CuII sited on an inversion centre (Cu1), a tetrahedral CuI on a general position (Cu2), and a trigonal CuI on a three-fold axis (Cu3), I with bridging cyanide ligands within coordination polymer materials.23The infrared spectrum of complex I in an unconstrained ligand environment such as this one and the bond valence sum calculation24 for this cation was also consistent with CuII. Each square planar Cu1 site is coordinated by trans cyanide and trans pyridyl groups from two L1 ligands, each Cu2 site is coordinated by the same ligand set but in tetrahedral geometry and each Cu3 is coordinated by three CN– ligands. Cyanide bridges occur between Cu1 and Cu2 sites, and between Cu2 and Cu3 sites, C3-symmetry with one isonicotinoyl group rotated such that the ester carbonyl is sited above the molecular cavity of the CTV-scaffold, and coordinates to Cu centres through all three pyridyl-donors bridging between Cu1 and Cu2 type centres, Square planar is not a common geometry for CuI4CuII1.5(L1)3(CN)6]+ coordination polymer has an unusual 3,4-connected 3D framework structure. A marked feature of the framework is the formation of hexagonal prismatic cages bounded by Cu1 and Cu2 centres, cyanides and six L1 ligands, ca. 3.1 nm. The hexagonal prisms are linked into a 2D network in the ab plane through cyanide bridges to the trigonal planar Cu3, a axis is shown in ab plane that are linked together through bridging L1 ligands. Each copper-cyanide ring in the hexagonal layer has a Cu24(CN)24 composition, The resultant + framework of 1 accounts for less than 40% of the crystal volume.Thermogravimetric analysis (TGA) of crystals of 3(L2)4(H2O)3]·6(OTf)·n(DMSO) 2 were obtained from the reaction of Cu(OTf)2 and L2 in dimethylsulfoxide (DMSO). The crystal structure was solved in space group C2/c. The given formula represents double the asymmetric unit of the structure, although counter-anions and any solvent were not located due to disorder. One crystallographically distinct CuII centre is sited on a two-fold axis whilst the other is on a general position. Each have similar square pyramidal coordination geometries being coordinated by an aquo ligand in the apical sites (Cu–O distances 2.337(4) to 2.407(8) Å) and pyridyl groups of L2 ligands in the four basal positions (Cu–N distances 1.975(6) to 2.047(5) Å). Each L2 ligand bridges between three CuII centres and both crystallographically distinct L2 ligands have similar conformations where each pyridyl group is approximately coplanar with its associated scaffold phenyl group. This is very much in contrast to the structures of complex 1 (and indeed the other complexes of L1 reported here) where the pyridyl (of iso-nicotinoyl) groups were closer to orthogonal with adjoining scaffold phenyls. The L2 ligands are arranged into head-to-head pairs linked together through the CuII cations to give a capsule-like motif, anti isomers or the meso–syn isomer.25 Capsule species where two CTV-ligands are bound together in a M3L2 species are referred to as metallo-cryptophanes, and discrete metallo-cryptophanes are known with 4-pyridyl-decorated CTV ligands. The closest Cu···Cu separations within the metallo-cryptophane motif of complex 2 are of the order of 17–17.5 Å and the distance between L2 centroids is 15.1 Å. Each CuII centre is shared by two metallo-cryptophanes to form a 2D metallo-cryptophane network with large hexagonal cavities, 3.123)(42.122) topology. Each cavity within the network is bound by six cryptophanes and six CuII centres with three aquo ligands directed into the cavity. The longest Cu···Cu separation across the cavity is 34.1 Å and the distance from the centre of the cavity to the O atom of an aquo ligand is 12.1 Å.The ligand L2 differs from L1 only in having a more conformationally flexible methyl ether linkage rather than ester between the pyridyl group and CTV scaffold, however the coordination polymers obtained with L2 are markedly different to those from L1. Plate-like lilac-coloured crystals of complex arenes, calix[4]arenes or cucurbiturils.4 We have previously reported a 2D network of linked metallo-cryptophanes in the complex [Ag3(NCMe)3(L)2Cl]2+ where L = tris{4-(3-pyridyl)phenylester}cyclotriguaiacylene.14 In that case the Ag3L2 metallo-cryptophanes were linked through Ag-μ3-Cl-Ag bridges, and there were no significant channels through the structure due to the manner of packing between the 2D networks. Robson's cyclotricatechylene-based coordination polymer features tetrahedral cages linked together into a network through oxide bridges.8 A number of metallo-cage motifs within metal–organic frameworks are known,28 and include Fujita's coordination polymer comprised of vertex-linked octahedral cage assemblies, materials which have recently been shown to act as crystalline sponges.A coordination polymer with embedded organic cryptophanes has been reported,2 is not a robust material and loses crystallinity on loss of solvent. TGA indicates that the lattice contains at least five additional molecules of DMSO per formula unit that were not located in the crystal structure 2Br2(H2O)(DMSO)]·2Br·n(DMSO) 3 forms from vapour diffusion of ethyl acetate into a DMSO solution of CuBr2 and L3. The crystal structure was solved in the space group P21/n and the given formula represents the asymmetric unit, with the uncoordinated Br– counter-anions modelled as disordered across several sites, and three molecules of DMSO were located in the structure. The two CuII centres both have square pyramidal geometry but have quite different coordination spheres. Cu1 features an axial Br– (Cu–Br 2.707(3) Å) and is also coordinated by four pyridyl groups from four L3 ligands (Cu–N 2.003(11) to 2.054(10) Å), while Cu2 has an axial Br– ligand (Cu–Br 2.793(5) Å), but only two trans pyridyl ligands (Cu–N 1.912(18) and 2.034(9) Å) with the remainder of the coordination sphere occupied by aquo and DMSO ligands, syn isomer metallo-cryptophanes, unlike in complex 3 where they were the chiral anti isomer. DMSO guest molecules were located for both L3 ligands, and are oriented with methyl groups directed towards the hydrophobic cavity, 3 are also distinct in that they do not possess approximate C3-symmetry. For both crystallographically distinct L3 ligands, one of the methylene–phenyl–pyridyl side-arms is bent inwards over the cavitand ligand molecular bowl such that the Caryl–O–CH2–Caryl torsion angle is ca. 75°, whereas the equivalent torsion angle is closer to 180° for the other methylene–phenyl-pyridyl side-arms. This means the three CuII centres that form the equatorial region of the metal-cryptophane form a scalene rather than equilateral triangle , giving the metallo-cryptophane motif a distinctly flattened aspect, Ligand L3 represents an extended-arm analogue of L2, and also forms a copper coordination polymer with linked metal-cryptophane structure. The complex ·2(OTf)·2NMP 4 was isolated from a mixture of Cu(OTf)2 and L1 in NMP where OTf is triflate, CF3SO3–. The crystal structure of 4 was determined in space group P1. The given molecular formula represents the asymmetric unit of the structure, although one OTf– was not located, see Fig. S3 ESI.II centre, both with distorted octahedral geometries. Cu1 is coordinated by two trans aquo ligands along the Jahn–Teller elongated axis (Cu–O distances 2.392(5) and 2.432(5) Å) and four pyridyl groups of L1 ligands (Cu–N distances 2.013(4) to 2.026(4) Å). Cu2 is coordinated by two trans pyridyl groups of L1 ligands (Cu–N distances 1.986(4), 2.008(5) Å), two trans NMP ligands (Cu–O distances 1.983(4), 1.991(5) Å) and two trans triflate anions on the elongated axis (Cu–O distances 2.321(6), 2.342(5) Å). There are two crystallographically distinct L1 ligands which are enantiomers, and one of which has a solvent NMP as an intra-cavity guest molecule. Both L1 ligands bridge between three CuII centres.The crystalline complex [Cu4 features a cationic [Cu2(L1)2(OTf)2(NMP)2(H2O)2]2+ 3,4-connected 2D coordination polymer, noting that Cu2 is topologically trivial in terms of network connectivity and the connecting nodes are the tripodal L1 ligands and the 4-connecting Cu1 which act as a square planar node. The Cu1 centres are linked by L1 ligands into a ladder arrangement shown in b axis. Each connecting centre in the framework is involved in 4-, 6- and 8-membered rings to give a simple framework topology of (4.62.8)(62.8)(4.62.82) noting that there are two types of 3-connecting centre, one in the centre of the ladder motif and the other at the sides, 2.62)(4.62)2 topology which forms a sheet rather than tubular motif and was found in complexes of the ligands tris[4-(4-pyridyl)benzoyl]-cyclotriguaiacylene or tris[3-(4-pyridyl)benzoyl]-cyclotriguaiacylene.12Complex 4 is reminiscent of the tubular 1D coordination polymer carboxylate-appended CTV-type ligands reported by Zheng,9 in both examples the tubes are bounded by four CTV scaffolds with molecular bowls directed inwards and linked by CuII cations. Unlike in Zheng's example, in complex 4 the size of the channels is restricted as the material exhibits 2D → 3D parallel interpenetration, vice versa, shown in  and between calixarene-based chains.6 Channels are evident along the a and c axes of the lattice, see ESI Fig. S4.The framework structure of 1, hence complexes with weakly coordinating anions would be anticipated to form completely different types of network structures. This is indeed the case, and reaction of ligand L1 with [CuI(NCMe)4]·BF4 in acetonitrile results in isolation of single crystals of complex [Cu(L1)(NCMe)]·BF4·1.5(CH3CN)·2H2O 5, while similar reaction of L1 with CuII(BF4)2 in N-methylpyrrolidone (NMP) gives complex [Cu2(L1)2(NMP)(H2O)]·4BF4·12NMP·1.5H2O 6.The coordinating anion cyanide plays a distinct structural role in complex 5 was determined in space group Pbca from data collected with synchrotron radiation. The asymmetric unit is composed of one CuI, one L1 and an acetonitrile ligand, a disordered BF4–, and solvent CH3CN and water sites. The CuI has approximate tetrahedral coordination and is coordinated by one acetonitrile and three symmetry related L1 ligands with Cu–N distances between 2.005(5) and 2.049(5) Å. The L1 ligand has approximate molecular C3 symmetry and each pyridyl group coordinates, bridging between three symmetry-related CuI centres.The crystal structure of complex + complex of 5 is a 2D coordination polymer with a hexagonal network of 63 topology also referred to as hex which forms in the bc plane of the crystal lattice, where both the CuI and L1 components act as 3-connecting centres, I positions are roughly coplanar but the network shows considerable puckering due to the bowl-shape of L1. Whether the L1 molecular cavities are oriented up or down alternates throughout the network, and the network is racemic, with the L1 enantiomer alternating along the c axis. The 2D networks pack along the a axis such that there is bowl-in-bowl stacking of the L1 molecular bowls of adjacent coordination polymers throughout the lattice, ca. 4.7 Å, indicating that there are no π–π stacking interactions between them. The a unit cell length of 9.459(2) Å is indicative of this type of bowl-in-bowl stacking of CTV-type scaffolds, akin to that observed in α-phase CTV.30 In the previously reported example of a coordination polymer involving L1 with a tetrahedral metal centre, namely complex [Ag(L1)2][Co(C2B9H11)2], a doubly-bridged chain structure is formed and a head-to-head motif is observed between L1 ligands of adjacent chains.15 In the overall crystal lattice of 5 channels are formed which are occupied by ligated and solvent acetonitrile, water and disordered BF4– counter-anions.The [Cu(L1)(NCMe)]2(L1)2(NMP)(H2O)]·4BF4·12NMP·1.5H2O 6 was determined in space group P21/c and the given formula represents the asymmetric unit, and the cationic [Cu2(L1)2(NMP)(H2O)]4+ complex forms a 3D coordination polymer. Both CuII centres have similar geometries, each being coordinated by pyridyl groups from three L1 ligands, one aquo and one NMP ligand in a distorted square pyramidal fashion, II centres the NMP ligand occupies the apical position and Cu–N distances range from 1.985(7) to 2.036(5) Å, Cu–OH2 distances 1.959(6) and 2.003(4) Å and apical Cu–O distances 2.169(7) and 2.282(5) Å. If the terminal aquo and NMP ligands are disregarded then the CuII cations have near T-shaped connectivity, with trans angles 159.5(3)° around Cu1 and 178.8(2)° around Cu2, and cis N–Cu–N angles between 87.0(2) and 90.4(2)°. Both crystallographically distinct ligand L1 types bridge between three Cu centres, however they are not structurally similar. One has all ester groups arranged with carbonyls pointed towards the centre of the molecular bowl, while the other has only one ester directed inwards. Both contain a guest NMP molecule within their molecular bowls, The crystal structure of complex ·Br·3.5DMSO 7. Complex 7 was obtained from vapour diffusion of ethyl acetate into a DMSO solution containing equimolar amounts of L2 and CuBr2. The crystal structure was solved in space group P1 and the given formula represents the asymmetric unit although an uncoordinated Br– counter-anion and additional solvent molecules were not located in the structure. The two L2 ligands within the asymmetric unit are of different ligand enantiomers. One methyl pyridyl group and one methyl group were both modelled as disordered across two positions, see ESI.II centres, shown in – ligands at Cu–Br distances 2.463(1) and 2.572(1) Å, and three pyridyl ligands from three L2 ligands at Cu–N distances ranging from 2.060(11) to 2.107(5) Å. Cu2 is square pyramidal with an axial Br– ligand (Cu–Br 2.814(2) Å), three L2 pyridyl ligands (Cu–N distances from 1.996(7) to 2.089(7) Å), and a coordinated DMSO at Cu–O distance 2.134(7) Å.The use of a coordinating anion with L2 results in the formation of a very different type of 2D coordination polymer structure in the material cyclononatriene core of ligands L1–L3 mean that the coordinating 4-pyridyl groups are often orthogonal to one another and this leads to 2D and chain networks which are significantly distorted from planarity and leads to formation of tubular like arrays such as in complex 4 or promotes formation of head-to-head linking of the ligands to form cage (complex 1) and capsule metallo-cryptophane (complexes 2 and 3) motifs within the framework. The cage and capsule-embedded frameworks in particular have molecular guest binding sites that are potentially available to new guest molecules, along with large channels throughout the ordered crystal lattice. Although these materials do not withstand complete evacuation of mother liquor, they exhibit new and highly promising structural types for the formation of porous materials with hierarchical pore spaces, and point to an important design principle – the networking of cage or capsule motifs.Host–guest relationships may form within or between networks that render embedded molecular recognition site of the ligands inaccessible, and this is illustrated here by complex


We propose direct synthesis of ultra-thin graphitic films on a dielectric substrate using sacrificial Ni catalyst layer, which significantly increases the crystallinity of the photoresist pyrolyzed at the temperature of 800 °C and above. A considerable amount of multilayer graphene in the photoresist film pyrolyzed in the presence of the Ni catalyst gives rise to an enhancement of the Raman signal of dye Sudan III molecules deposited on the substrate. We demonstrate comparable enhancement of the Raman signal from Sudan III molecules deposited on the fabricated graphitic substrate and those deposited on graphene, which was conventionally transferred to the silica substrate. However, the direct synthesis of these highly ordered atomically thin graphitic films on dielectric and semiconductor substrates remains a difficult task. This is partially because the chemical vapor deposition (CVD), a conventional graphene synthesis technique, relies on a metal (e.g. Cu and Ni) catalyst3, while applications usually require graphene deposited on a dielectric/semiconductor substrate. The transfer of an atomically thin film essentially involves manual processing, which is the major bottleneck in the implementation of graphene in practical devices. The direct and scalable graphene synthesis on insulating substrates is an eagerly awaiting breakthrough that should open pathways towards an industrial scale exploitation of graphene in sensing, electronics and photonics4.Ever since graphene isolation 15 years ago, a variety of applications have been proposed and demonstrated for both single and multilayered graphene8. In the framework of such an approach, one needs just to remove the remains of the metal film from the substrate in order to obtain graphene deposited on the dielectric substrate. Although this technique is relatively simple and straightforward, the electronic properties of the obtained polycrystalline single- and multilayer graphene are usually poorer than those of the transferred one5. Nevertheless directly deposited graphitic films are suitable for a number of applications in optics and photonics4.In the search for the methods of the direct graphene deposition there have been proposed CVD based techniques that exploit sacrificial ultrathin copper or nickel catalyst films pre-deposited on the target dielectric substrate rather than bulky metal foils9. This enhancement mechanism is well balanced and delicate system affected by the graphene Fermi level, highest occupied molecular orbital (HOMO) and lowest unoccupied molecular orbital (LUMO) energy levels of the analyte, excitation laser wavelength and the structure of the analyte molecule10. Together with ability of graphene to protect the analyte from photobleaching and suppress fluorescence such a specificity may be highly beneficial for molecular spectroscopy11.Graphene enhanced Raman spectroscopy (GERS) manifests itself as the enhancement of the Raman signal from an analyte deposited onto the graphene sheet. In contrast to the conventional surface-enhanced Raman spectroscopy (SERS), which exploits local field enhancement in the vicinity of metal nanoparticles deposited on the dielectric substrate, GERS is driven by the charge transfer between the analyte and graphene12. We demonstrate in particular that fabricated GERS-active substrates provide the enhancement of the Sudan III Raman lines comparable to that can be obtained by using graphene conventionally grown on a copper foil and transferred onto a dielectric substrate.In this paper, we demonstrate wafer scale direct deposition of multilayer graphene on a dielectric substrate by using pyrolysis of photoresist film, which is graphitized at 800 °C in presence of a 10 nm thick nickel nanocatalyst layer. We show that fabricated films can be used for GERS of Sudan III dye molecules, an illicit, carcinogenic food and fuel marker dyeIn order to fabricate the graphitic film, we evaporated a 10 nm ±2 nm) thick nickel catalyst layer on a 0.5 mm thick silica substrate, spin coated it with nLOF resist , and then baked the structure in hydrogen atmosphere  at temperature 600 °C–900 °C. Schematic illustration of the graphitic film preparation is shown in Fig.  nm thickThe graphitic film surface morphology and thickness were studied by atomic force microscope (AFM - Thermo Microscopes Explorer 4400–11) and thickness by stylus profiler (Weeko Dektak 150), respectively. Figure 13.To reveal how the presence of the Ni layer affects the graphitization, we measured the Raman spectra of the fabricated samples using 514 nm excitation wavelength. In order to avoid effects of the graphene self-heating the excitation beam power was fixed at 0.25 mW−1 and ~1580 cm−1, respectively)14. The 2D peak at ~2700 cm−1, a widely recognized signature of the crystalline graphene15, is barely visible for the films pyrolyzed both with and without Ni catalyst.The Raman spectra of samples processed at temperatures of 600 °C and 700 °C shown in Fig. 2 hybridized domains increases14.When the process temperature exceeds 800 °C the Raman spectrum of the PPF remains almost unchanged, while the Raman spectrum of the PPF+Ni changes drastically. Figure Next, we performed GERS experiments with films fabricated at different process temperatures using aromatic azo dye (Sudan III) as a target analyte . In the experiment, droplets of 0.1 mM Sudan III diluted in acetone were placed on the surface of the PPF and PPF+Ni samples and dried out.Measured Raman spectra of the prepared samples are presented in Fig. −1 , rinsed the sample in water and compared Raman spectra from the PPF with and without nickel particles. One can see from Supplementary information Fig. 17, i.e. it takes place well above melting temperature of inexpensive and flexible glass substrates18. In contrast, the PPF+Ni with improved crystallinity can be fabricated at a considerably lower temperature, which can open a way for a scalable fabrication of the GERS substrates by using inexpensive boron silica glass, which is widely used in chemical industry.Typically, CVD graphene is formed in the temperature range 900–1000 °C19. In this temperature range, Ni catalyzes transformation of the amorphous carbon into graphitic crystal while the Ni film will melt and recede. This process resembles multilayer graphene synthesis on the surface of the Ni foil, when dissolved carbon atoms leaves the foil as soon as the temperature in the CVD chamber starts decreasing20. In our experiment situation is similar - Ni particles in the resist film recede leaving graphitic crystal behind.We believe that one major reason for the formation of the graphitic carbon film is a high carbon solubility in nickel. One may expect that above 700 °C, the carbon solubility increases drastically giving rise to merging Ni and the resist films21. However, the strong 2D-peak, suppressed D-peak and sharp G-peak indicate that the adding Ni catalyst results in higher degree of crystallization in comparison to the PPF on a bare silica substrate21. That is despite the fact that PPF+Ni is multilayered and defected film, it performs well in terms of GERS. This experimental finding well corresponds to earlier observations that multilayer graphene can have as good GERS properties as monolayer graphene and that defects can even improve the GERS signal23. Therefore, the proposed nickel nano-catalyst assisted growth of multilayer graphene can be employed for scalable fabrication GERS substrates capable of provide an increase of the Raman signal comparable to that of conventional transferred graphene. Moreover, the proposed technique allows one to either remove Ni particles if required but otherwise can provide a nice way to obtain metal-graphene composite material, which may lead into useful applications25.Since melting and receding of Ni film are very random processes, one may expected that the graphitic film, which is formed between Ni particles, is rather defected. The relative intensities of the D- and G-peaks in the Raman spectrum see Fig.  indicate26. (ii) The structure rule, according to which the GERS enhancement can occur if the analyte molecule is structurally “compatible” with graphene. For example, it may have Dnh symmetry or it may consist of parallel connected conjugated benzene rings, which match the structure of graphene enabling strong coupling (and charge transfer)26.Overall, GERS is a highly selective process and it has been demonstrated that the GERS effect may occur when one of two molecular selectivity rules is satisfied. (i) The energy rule, which states that in order to generate significant enhancement, HOMO/LUMO energy levels of the analyte molecule need to be close enough to the Fermi level of graphene27, while the Fermi energy level of the graphene and graphite is at −4.6 eV. Therefore, for the 514 nm laser, the energy level rule does not give a justification for GERS26. Thus, we believe that in case of Sudan III, the enhancement is related to the structure rule rather than the energy rule. Although, Sudan III is non-symmetrical molecule, it has a parallel conjugated carbon ring structure that enables coupling to the graphitic PPF+Ni substrate, which in turn increases π-π interaction and thus enhances the Raman signal26. On the contrary, benzene rings of Sudan III cannot couple to the sp2 orbitals of strongly amorphous PPF implying that no enhancement can take place.The HOMO/LUMO energy levels of Sudan III are −5.8 eV and −3.1 eV, respectivelyIt is also worth noting that the Raman enhancement is slightly stronger in PPF+Ni in comparison to a monolayer graphene fabricated by a conventional transfer technique (see Supplementary information). Here the difference between PPF+Ni and transferred graphene most likely originates from the rough surface of the PPF+Ni sample, as the roughness increases the surface area, which can then consequently increase the number of analyte molecules coupled to the PPF+Ni substrate surface. Thus, the presented technique, which results in rougher graphitic films, can actually be beneficial to GERS.Finally, in order to confirm that presence of the graphene flakes on the rough surface of the PPF+Ni film enhances Raman scattering of Sudan III, we fabricated also a rough PPF described in Supplementary information. However, no enhancement of the Raman signal was observed when Sudan III molecules were deposited on such a rough PPF substrate. This experimental finding clearly indicates that the enhancement of the Raman signal with PPF+Ni substrate, indeed, originates from the GERS rather than simply the increase of the surface area.28. However, rather high surface roughness of the fabricated PPF+Ni film improved slightly Raman scattering in comparison to the transferred, flat graphene substrate.In summary, we have proposed a scalable and inexpensive technique for a GERS substrate fabrication by pyrolysis of the photoresist film deposited on a silica surface coated with a 10 nm thick Ni layer. We observed that at temperature of about 800 °C the Ni film melts and recedes to submicron size Ni particles. At this same temperature range the catalysis of Ni begins and the amorphous PPF undergoes transformation from amorphous carbon to polycrystalline graphitic film. This transformation was observed with Raman measurements. Furthermore, the Raman enhancement of Sudan III molecule on this graphitic film was observed to be about one order of magnitude. It is well-known that GERS increases the Raman signal much less in comparison to those metallic SERS platforms−5 mBar vacuum using 99,99% pure Ni pellets. Sample was next coated with a resist  spin coating at 2000 rpms. Spin coating resulted in about 230 nm thick resist film on the Ni+silica substrate.Nickel thin film was evaporated on a 0.5 mm thick, 25 mm in diameter fused silica substrate. Thermal evaporation was done by Leybold Univex 300 in 102 flow in order to remove oxygen from the chamber. Next the chamber temperature was risen to the actual process temperature by 20 °C/min. After 10 min baking in the chosen process temperature the sample was cooled down to room temperature (overnight) in static 8 mBar H2 atmosphere.Samples were pyrolyzed in a chemical vapor deposition (CVD) apparatus. The CVD chamber was pumped for one hour in 5 sccm HGERS measurements were performed using Renishaw inVia Raman microscope with 514 nm excitation and Sudan III as the target molecule. The samples were prepared for the measurements by depositing a 2 µl droplet of 0.1 mM Sudan III  diluted in acetone on the surface of the samples. The temperature dependence measurements were carried out using 0.25 mW power, 10 s integration time and 100X objective. In total, 9 measurements were recorded for each sample  and the spectra were then averaged. The Raman mapping measurements were collected from a 200 µm × 600 µm area at 10 µm intervals using 0.5 mW power, 5 s integration time and 20X objective.Supplementary information


CADASIL) is a familial fatal progressive degenerative disorder. One of the pathological hallmarks of CADASIL is a dramatic reduction of vascular smooth muscle cells (VSMCs) in cerebral arteries. Using VSMCs from the vasculature of the human umbilical cord, placenta and cerebrum of CADASIL patients, we found that CADASIL VSMCs had a lower proliferation rate compared to control VSMCs. Exposure of control VSMCs and endothelial cells (ECs) to media derived from CADASIL VSMCs lowered the proliferation rate of all cells examined. By quantitative RT‐PCR analysis, we observed increased Transforming growth factor‐β (TGFβ) gene expression in CADASIL VSMCs. Adding TGFβ‐neutralizing antibody restored the proliferation rate of CADASIL VSMCs. We assessed proliferation differences in the presence or absence of TGFβ‐neutralizing antibody in ECs co‐cultured with VSMCs. ECs co‐cultured with CADASIL VSMCs exhibited a lower proliferation rate than those co‐cultured with control VSMCs, and neutralization of TGFβ normalized the proliferation rate of ECs co‐cultured with CADASIL VSMCs. We suggest that increased TGFβ expression in CADASIL VSMCs is involved in the reduced VSMC proliferation in CADASIL and may play a role in situ in altered proliferation of neighbouring cells in the vasculature.Cerebral autosomal‐dominant arteriopathy with subcortical infarcts and leukoencephalopathy ( NOTCH3 is a type I transmembrane receptor belonging to the Notch signalling family, one of a group of “elite” intracellular signalling pathways.NOTCH3 gene have been detected,NOTCH3. These mutations affect the number of cysteine residues, leading to an unpaired cysteine residue.CADASIL is caused by mutations in the In CADASIL, pathological observations have shown abnormality in numerous cells within neurovascular units, including VSMCs and endothelial cells (ECs).NOTCH3 have a lower proliferation rate than control VSMCs.It has been proposed that the degeneration of the smooth muscle layer in arteries/arterioles of patients with CADASIL is the result of increased VSMC death.Transforming growth factor‐β (TGFβ) is a cytokine that has multicellular function in the cells and has a great involvement in the fibrosis of several diseases. The role of TGFβ in cell proliferation is well‐established.NOTCH3, we report a role of TGFβ in VSMC and EC proliferation rates, suggesting it is a critical mechanism involved in CADASIL pathophysiology.This study aimed to investigate why VSMC proliferation is impaired in CADASIL. Using immortalized human umbilical, placental and cerebral VSMCs derived from individuals carrying the R133C mutation or control 22.1Patient‐derived VSMC cell lines: umbilical cord (Umb)‐ and placental (Pla) VSMCs were established from blood vessels of genetically verified patients with CADASIL and control subjects as previously described.2.2Patient‐derived VSMC cell lines: Umb‐,2.3™ Cytometry, and data were analysed using the CellQuest software (BD Biosciences). The proliferation rate of VSMCs was also investigated using immunofluorescent staining of the proliferation‐associated Ki67 protein as previously described.CellTrace carboxyfluorescein succinimidyl ester (CFSE) (Invitrogen) was used to enable the measurement of VSMC proliferation. Carboxyfluorescein succinimidyl ester is a membrane permeant fluorescein‐based dye that can be used to track cell division due to the progressive halving of the fluorescence intensity of the dye in cells after each division. Briefly, CADASIL and control VSMC were grown to 90% confluence and harvested after washing twice with PBS. The cells were centrifuged and re‐suspended with 1 mL pre‐warmed PBS followed by labelling with 1 μmol/L CFSE for 10 minutes at 37°C. Equal numbers of CFSE‐labelled cells were incubated for either 3 or 7 days. Unlabelled cells were used as negative controls. Flow cytometry was performed using a FACSCalibur2.4Human fibroblasts and UASMC were exposed to media that CADASIL PlaVSMC cells were grown in conditioned medium for 24 hour, and 3 and 7 days. The PlaVSMC‐conditioned medium was applied either undiluted or diluted 1:1 with complete medium. The proliferation rate of VSMCs was investigated using immunofluorescence staining of the proliferation‐associated Ki67 protein as described previously.2.5RNA was isolated from Pla‐ and CerVSMC using the RNeasy Mini kit (QIAGEN). The quality of RNA was determined with RIN (RNA Integrity Number) of 10 Figure B. cDNA w2.64 cells/mL in 24 well plates and cultured overnight. The medium was then changed and cells were incubated in the presence or absence of 12 μg/mL TGFβ‐neutralizing antibody for 24 hour. Vascular smooth muscle cells or ECs were fixed and stained with Ki67 antibody as described earlier.Transforming growth factor‐β pan‐specific neutralizing antibody was derived from recombinant human TGFβ1, porcine platelet‐derived TGFβ1.2, porcine platelet‐derived TGFβ2 and recombinant amphibian TGFβ5 porcine platelet‐derived TGFβ1 and 2, which detects TGFβ1,TGFβ1.2, TGFβ2, TGFβ3 and TGFβ5 . Vascular smooth muscle cells was seeded at 5 × 102.74 cells/mL in 6‐ or 24‐well plates. Pla‐ and CerVSMCs were seeded at 5 × 104 cells/mL onto the membrane of transwell cell culture inserts and allowed to grow overnight. After 24 hour, the transwell insert membrane containing the VSMCs was placed into plates containing ECs. Vascular smooth muscle cells were treated with a TGFβ‐neutralizing antibody (12 μg/mL).To examine the effect of TGFβ produced by VSMC on EC proliferation rates, Pla‐ and CerVSMC were co‐cultured with ECs in a non‐contacting co‐culture transwell system . Endothelial cells were plated at 5 × 102.8t‐test was used for two‐group comparisons. P‐values <.05 were considered significant. The results are representative of three independent biological replicates expressed as mean ± SEM.Statistical comparison of values obtained for VSMC cell lines was analysed by one‐way ANOVA followed by Bonferroni's post hoc test. Student 33.1right and left panel), confirming that the control PlaVSMCs were actively proliferating as compared to CADASIL PlaVSMCs after 3 and 7 days . Reduced CADASIL UmbVSMCs proliferation rate (divided) compared to control UmbVSMCs was also confirmed P < .01).To examine the proliferation rate in PlaVSMC and CerVSMC, we labelled cells with CFSE dye for flow cytometry or stained with Ki67 antibody for confocal microscopy Figure A‐D. A res Figure A. Analyss Figure B. The deP < .01).Proliferation was further evaluated by anti‐Ki67 staining and confocal microscopy Figure C. A lowe3.2P < .01). A decreased proliferation rate was also observed when two different conditioned medium dilutions (1:1 and 1:2) from CADASIL PlaVSMCs were added to the commercial UASMC cell line  in CADASIL VSMCs compared to control VSMCs as quantified by qRT‐PCR. However, the increased level of TGFβ3 was more obvious as compared to the TGFβ1, ‐2. Notably, the expression of TGFβ3 was significantly higher in both CADASIL Pla‐ and CerVSMCs compared to control VSMCs . Next, we investigated whether removing TGFβ would increase the proliferation rate of CADASIL VSMCs to the level of control VSMCs. Quantification analysis demonstrated a significantly lower proliferation rate of the CADASIL PlaVSMCs (P < .01), which addition of TGFβ‐neutralizing antibody restored to control levels . Confocal microscopy analysis showed a significant decrease in Ki67‐positive cells in CADASIL CerVSMCs (P < .05), and the same effect of TGFβ‐neutralizing antibody was noted in CADASIL CerVSMCs .To identify candidate factors secreted by CADASIL VSMC that could result in a lower proliferation rate of neighbouring cells, we monitored the expression level of several genes in the cell proliferation pathway by qRT‐PCR. We identified a number of genes with altered expression levels in CADASIL Pla‐ and CerVSMCs as compared to their control counterparts . The proliferation rate of ECs did not change in the presence of TGFβ‐neutralizing antibody as shown by confocal microscopy  in Pla‐ and CerVSMCs, where the increased level of TGFβ3 was more obvious as compared to the TGFβ1 or TGFβ2. TGFβ isoforms are expressed in the central nervous system and among them, TGFβ1 has been widely considered as an injury‐related cytokine which is a crucial regulator of nervous system physiology and its vasculature. TGFβ2 and TGFβ3 are also localized in radial glial cells, neuronal cell bodies in the telencephalic cortex and cerebellum, suggesting a role in the regulation of neuronal migration and differentiation as well as glial cell proliferation and differentiation. The expression of TGFβ2 and TGFβ3 persists in the entire adult CNS areas including cortex, hippocampus, striatum, brainstem and cerebellum. In the CNS, TGFβ2 mRNA and protein are predominantly expressed in astrocytes of white matter.TGFβ3 gene expression in CADASIL VSMCs. Other TGFβ‐isoforms genes were also expressed higher in CADASIL VSMCs compared to control VSMCs. Variability in expression of the TGFβ3 gene could result in enhanced myofibroblast activity, proliferation and induction of ECM synthesis.Transforming growth factor‐β has been shown to be involved in differentiation/proliferation of VSMCs.TGFβ receptor expression may have an impact on EC proliferation in ECs co‐cultured with CADASIL VSMCs. Our qRT‐PCR data showed down‐regulation of TGFβR3 in ECs co‐cultured with CADASIL VSMCs (Table TGFβR1 and TGFβR2 in ECs were unaffected (data not shown). Interestingly, the presence of TGFβ‐neutralizing antibody significantly changed the TGFβR3 gene expression level, which suggests that TGFβ exerts its action on EC proliferation through the TGFβR3‐receptor. This data suggest that several members of the TGFβ superfamily are potentially involved in proliferation rate changes in our system.Degenerative changes in capillary vessels are also involving ECs in CADASIL and showing endothelial abnormalities which suggests a secondary character of endothelium damage in CADASIL.Cs Table , while TCDC34, CHEK1, etc.).Several studies have demonstrated that during NOTCH target gene activation, there is an interaction between NOTCH and TGFβ signalling pathways. Both NOTCH and TGFβ are known to be involved in many regulatory functions of VSMC and other cells, such as Wnt and MAP kinase signalling, cell cycle‐related genes (NOTCH3 gene, the synergy between TGFβ and NOTCH3 might be compromised. In this study, increased TGFβ expression was observed in CADASIL NOTCH3‐mutated VSMCs. It could therefore be expected that also other factors involved in cell proliferation, differentiation and ECM family gene expression in VSMCs are affected. Further, changes to NOTCH3 function due to mutations and increased TGFβ in CADASIL can also have an impact on TGFβ receptor expression in ECs, as shown in our study. The interplay between TGFβ and NOTCH3 in CADASIL needs further elucidation for the understanding of the VSMCs degeneration.So far, few studies have addressed gene expression and signalling pathways in CADASIL. However, it has been shown that both TGFβ and NOTCH are involved in differentiation and proliferation of VSMCs. In CADASIL, due to mutations in We believe that increased TGFβ3 reflects an inflammatory condition and even an involvement of TGFβ in fibrosis in CADASIL. Support for our hypothesis is the outcome of several studies that showed NOTCH3 and TGFβ1 signalling play a key role in the pathogenesis and progression of chronic cardiovascular disease.In CADASIL, the cells in the vasculature can be exposed to complex changes, such as aggregated NOTCH3, and GOM. In this study, we have implicated TGFβ and TGFβ receptors as additional players in CADASIL disease. Therefore, more detailed studies are needed to elucidate the interplay of TGFβ signalling in the vasculature in CADASIL. In summary, our data provide evidence for direct involvement of TGFβ in proliferation of VSMCs and ECs. We believe that our findings provide a better understanding of the pathogenesis of CADASIL disease.The authors have no conflicts to disclose.MP, LF, NM, KC, CG and EVS performed the research; MP, HK, MV, BW, ES and HB designed the research study; MP, NM, LF and HB analysed the data; MP, KC, HK, ES and HB wrote the manuscript. Click here for additional data file. Click here for additional data file. Click here for additional data file. Click here for additional data file. Click here for additional data file. Click here for additional data file. Click here for additional data file. Click here for additional data file.


Samples with diverse urea/GO mass ratios were treated at 800 °C in inert atmosphere to generate graphene with diverse microstructural characteristics and levels of nitrogen doping. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were used to study the microstructural features of the products. The effects of doping on the samples structure and surface area were studied by X-ray diffraction (XRD), Raman Spectroscopy, and Brunauer Emmet Teller (BET). The GO and urea decomposition-reduction process as well as nitrogen-doped graphene stability were studied by thermogravimetric analysis (TGA) coupled with mass spectroscopy (MS) analysis of the evolved gases. Results show that the proposed method offers a high level of control over the amount of nitrogen inserted in the graphene and may be used alternatively to control its surface area. To demonstrate the practical relevance of these findings, as-produced samples were used as electrodes in supercapacitor and battery devices and compared with conventional, thermally exfoliated graphene.Graphene sheets doped with nitrogen were produced by the reduction-expansion (RES) method utilizing graphite oxide (GO) and urea as precursor materials. The simultaneous graphene generation and nitrogen insertion reactions are based on the fact that urea decomposes upon heating to release reducing gases. The volatile byproducts perform two primary functions: (i) promoting the reduction of the GO and (ii) providing the nitrogen to be inserted  Graphene, the two dimensional material consisting of an atom thick layer of carbon in a spangement , and itsangement ,4. GO caangement ,6. Once angement , plasma angement , sonicatangement ,9, electangement  and use angement ,13,14,15While graphene produced from GO has been reported as a material with a promising electrochemical performance when used for the manufacture of supercapacitor and lithium ion battery electrodes, its practical energy storage capacity is still well below the theoretical values ,17. TherThe intial work, published in 2010 , was folThe microstructural characteristics of the nitrogen-doped graphene samples were analyzed by scanning electron microscopy (SEM) and transmission electron microscopy (TEM)  and compThe first evidence of nitrogen insertion in the doped samples was encountered in the spectroscopic data generated. The energy dispersive X-ray spectroscopy (EDS) spectra  of sampld002 peak toward higher angles  as the urea/GO ratio in the precursor increases. This suggest that the presence of urea not only affects the exfoliation process, but that the amount of urea may be used to control the level of exfoliation and thus the interlayer spacing in graphene samples. As shown previously, the decomposition of urea generates additional volatile species that aid the reduction of the GO (more complete reduction), but also results in an overall less effective expansion process [To obtain further insight into potential structural changes in the graphene with increasing urea content, we analyzed all samples using X-ray diffraction (XRD). The obtained XRD patterns a reveal  process . Consequ process  or to th process . However process ,34,36,372-carbons, as well as the resulting second order and combination modes. The most notable spectral change is the increase in the D band-to-G band intensity ratio (ID/IG) with increasing urea content. Although, according to XRD data  than most urea/GO-derived samples, even without any nitrogen inclusions. This discrepancy results from differences in the exfoliation mechanism as previously discussed [ID/IG therefore provides a fast and convenient tool to estimate the concentration of nitrogen atoms in doped-graphene samples. For a more detailed description of the Raman spectra of urea/GO-derived graphene samples, the reader is referred to work by Mowry et al. [XRD data , both iniscussed . Given ty et al. . While Raman measurements provided indirect evidence for a direct relationship between urea content and level of N-doping, we conducted thermogravimetric analysis (TGA) coupled with mass spectroscopy (MS) to directly evaluate the N-content of the samples. TGA was also conducted in order to identify the differences in thermal stability of the diverse samples. The burn off temperatures of the specimens were studied under an oxygen-containing atmosphere and the mass spectral signal of the byproducts generated as the carbon structure burns off. 2 (hence similar onset temperatures for oxidation), the burn off rate is higher for samples containing less urea. These samples are more exfoliated, and thus more porous, as compared to graphene with higher stacking order. At temperatures around 500 °C, mass transport becomes the rate-limiting reaction step in carbon nanomaterial oxidation, leading to reduced oxidation rate in low porosity materials. In addition, the lower concentration of two dimensional defects (due to nitrogen repair of lattice defects during exfoliation with urea in a nitrogen atmosphere) also contributes to the higher thermal stability with increasing urea content. The TGA-MS results are thus in good agreement with XRD and Raman data, providing further evidence for the increment in structural order with increasing urea content. While all samples consist of graphitic sp2 as main byproduct for pure graphene and CO2 and NO for doped graphene under the atmosphere and conditions used. As concentrations of nitrogen in the precursor increase, the temperature at which maximum NO and CO2 leave the structure/off gas increases  surface area measurements for samples with diverse doping levels show an exponential decay as the amount of urea used in the precursor increase (correlated to the nitrogen doping level). Such change is also evident when weighing the same amount of the different powders: As the amount on nitrogen introduced increase, the level of exfoliation decreases. See 2  demonstrated superior capacitive characteristics to the un-doped graphene proving that the doping has a positive benefit to the electrical properties of the material. The 0.33 sample performed nearly identically to the bare thermally exfoliated graphene. The 0.2 and 0.16 samples demonstrated worse electrical characteristics than others, further emphasizing the presence of an optimal peak doping concentration in the vicinity of 0.25.Nitrogen-doped graphene has also been reported to provide superior electrochemical performance and cycle stability in Li-ion batteries, as compared to pristine graphene ,45. The reversible/Cirreversible capacity ratio. Under this consideration, N-doped graphene would be the more favorable anode material (0.52 vs. 1.49); (3) Given the lower surface area values of N-doped graphene (<10 m2/g), we can conclude that the addition of urea is an effective way to control the extent of GO exfoliation, allowing for the optimization of Creversible/Cirreversible in graphene materials. As expected, the thermally exfoliated graphene is subject to a high irreversible capacity loss (>1600 mAh/g), exhibiting a reversible specific capacity of ~860 mAh/g at the end of the 1st cycle. However, in contrast to previous studies ,45, the 2 units, there seems to be an optimal value (sample with 0.25), demonstrating that despite the reduction in surface area, there must be an addition in charge mobility that improves sample performance.The resulting capacitance of the un-doped graphene prepared by thermal exfoliation in units of F/g was remarkably higher when compared with doped samples. Specific capacitance has been strongly correlated to the surface area of the graphene or doped graphene employed in the electrode material. Such phenomenon could be explained when relating the capacitance to the number of charges interacting at the electric double layer: the higher the surface area, the larger the sites where charges in the interface could be created. However, once the capacitance values get normalized by surface area, using F/mreversible/Cirreversible ratio in lithium ion batteries. High control over surface area and interlayer spacing may also provide opportunities for other advanced battery technologies, such as sodium ion batteries, where conventional graphitic materials are unsuitable. Electrochemical testing demonstrated that addition of urea offers a unique pathway to optimize the critical CWhile an in-depth discussion of the electrochemical performance of urea/GO-derived graphene is outside the scope of this study and will be published elsewhere, the presented data demonstrates the unique characteristics of urea/GO-derived graphene and the opportunities in energy storage applications. GO was produced from synthetic graphite powder  following a variation of the Hummer’s method . A mixtuin situ via the thermal reduction of graphite oxide in the presence of urea under inert atmosphere (reduction expansion synthesis). A sample in which no urea was employed, along mixtures with urea/GO ratios of 0.16, 0.2, 0.25, 0.33, 0.4, 0.6, 0.8, and 1 were used as precursors. The urea/GO mixtures were ground by hand in a mortar, placed in an alumina boat, and introduced into a quartz tube provided with airtight end adapters for atmosphere control. The quartz tube was flushed with nitrogen and inserted for 10 min into a Lindberg tubular furnace preheated to 800 °C. Previous studies demonstrated that the urea completely decomposes by 600 °C, however, by heating to 800 °C the mixtures seem to achieve a greater degree of expansion and oxygen groups loss, producing a higher quality product [Doped graphene was produced  product . The resScanning and transmission electron microscopy were employed to study the microstructure of the samples. A Zeiss Neon 40  operating between 2–20 kV was employed for SEM studies. A JEOL 2010F FASTEM  field emission gun STEM/TEM equipped with Gatan image filtering (GIF)  system was used to acquire the TEM and EELS data. Energy dispersive X-ray spectroscopy was performed utilizing the TEM. −1. The Cu X rays generated had a λ of 1.5418 Å. The diffraction data was analyzed using Xpert Highscore Software  using the ICDD PDF2  database.Powder samples were analyzed by X-ray diffraction using a Philips PW1830 X-ray diffractometer  operating in a continuous scan mode from a 2θ of 5° to 70°. Two second steps were utilized with a step size of 0.020° at a speed of 0.01° sRaman spectroscopy studies were carried out using an inVia confocal raman microspectrometer  with 514-nm laser excitation under ambient conditions. Spectra were recorded using a total accumulation time of 20 s. Nitrogen isotherms were used to determine total surface area and micropore volume distribution employing Brunauer Emmet Teller (BET) and Horvath-Kawazoe (HK) methodologies, respectively. A Quantachrome Nova 4200e surface area and pore size analyzer  using 9 mm large bulb long cell was employed for this analysis. Samples were degassed following an increasing temperature profile, starting at 100 °C for 30 min, followed by 300 °C for 150 min.Thermogravimetric and evolved gas mass spectral analyses were performed to identify the various off gassing temperatures and constituents of heated samples. A Netzsch STA 449 F3 Jupiter connected to a MS403C Aëolos mass spectrometer  was used. The temperature programmed oxidation of the doped samples generated in the furnace was conducted in an atmosphere that contained 80% inert gas and 20% oxygen (simulated air) with the intent of promoting the burn off of both the un-doped graphene and the doped samples.N-Methyl-2-pyrrolidone (NMP) and polyvinylidene fluoride (PVDF) . For supercapacitor testing, the resultant slurry was spread onto 25 μm thick roughened nickel foil to a height of 350 μm using a film applicator, placed into a Barnstead Lab-Line L-C oven  at 100 °C for 1 h and then punched into 15 mm disks using a MTI Corporation precision disc cutter. The completed electrodes were assembled in pairs between two acrylic non-conductive polymeric blocks, separated by a celgrad membrane  soaked in electrolyte KOH solution. For battery electrode preparation, slurries were cast on a copper foil using a blade-applicator. The dried electrodes (thickness ~0.2 mm) were cut to 12 mm-diameter disks, transferred to an Ar-filled glovebox, and assembled in CR2032 coin cells in half-cell configuration . Electrochemical testing was conducted in an electrolyte containing 1M LiPF6 in EC:DMC:DEC (1:1:1). Capacitor and battery electrodes were fabricated using a mixture of the as-produced graphene or doped graphene, acetylene black, Battery and capacitor testing was conducted using a Maccor Model 4200M tester . Capacitance was evaluated by applying three different current sources for a set number of cycles and measuring the resultant voltage trends. The currents and number of cycles utilized were 150, 350, and 500 μA for 50, 50, and 20 cycles, respectively. Battery testing was conducted using a charge/discharge rate current of C/10. We were able to successfully dope graphene with nitrogen functional groups with a high level of control over the degree of doping utilizing the reduction expansion synthesis (RES) method. SEM and TEM analysis revealed that the graphene produced resulted in thin sheets of disordered graphene, as is typical for samples made by reduction of graphite oxide. EELS analysis provided unequivocal evidence that nitrogen was embedded in the structure of graphene. Similarly, EDS shows nitrogen peaks in the doped samples further demonstrating the effectiveness of RES in doping the graphene. The BET surface area measurements for samples with diverse doping levels show an exponential decay as the amount of urea used in the precursor increase. These values can be controlled by modifying the amount of urea in the precursor samples.d” of the (002) planes as the concentration of nitrogen increased. The Raman intensity ratio between D and G band, ID/IG, increases with increasing urea content while the G band position remains nearly unchanged. These results confirmed the trends observed by BET analysis and indicated that there is a pronounced surface area loss as the amount of doping increases. XRD analysis of various levels of doped graphene revealed peak shifts indicating lower interplanar spacing “All doped samples demonstrated to have high thermal stability up to 550 °C when studied by TGA and mass spectral analysis performed under temperature programmed oxidation conditions. The MS data corroborated that the amount of urea in the precursor samples can be directly linked to the amount of nitrogen species lost during temperature programmed oxidation/burn off of the samples, thus providing a path to control the amount of nitrogen species included in the structure and the surface area of the products.Indeed, the introduction of nitrogen impurities by RES method in graphene produced from GO had a large effect in the graphene structure and improved the performance of electrodes based on such nitrogen doped graphene. Doping graphene with nitrogen clearly represents a step ahead in the development of electrode materials for new energy storage devices and the method presented herein provides a route to introduce a controlled amount of nitrogen and imprint a specific surface area in the products.


Photoresponse studies reveal that photoresponsivity in one-layer MoS2 is surprisingly higher than that in seven-layer MoS2 by seven times. Spectral-dependent studies further show that the internal quantum efficiency in one-layer MoS2 can reach a maximum of 65%, far higher than the 7% in seven-layer MoS2. Our theoretical modelling shows that asymmetric potential barriers in the top and bottom interfaces of the graphene/one-layer MoS2/graphene heterojunction enable asymmetric carrier tunnelling, to generate usually high photoresponsivity in one-layer MoS2 device.Two-dimensional layered transition-metal dichalcogenides have attracted considerable interest for their unique layer-number-dependent properties. In particular, vertical integration of these two-dimensional crystals to form van der Waals heterostructures can open up a new dimension for the design of functional electronic and optoelectronic devices. Here we report the layer-number-dependent photocurrent generation in graphene/MoS Vertical integration of two-dimensional materials can open unprecedented possibilities towards design of efficient optoelectronic devices. Here, the authors investigate the photoresponse properties of a graphene/MoS2/graphene heterostructure, revealing promising quantum efficiency performances. However, Shockly–Read–Hall recombination and Langevin recombination in the WSe2/MoS2p–n junction could compromise the carrier extraction performanceThe two-dimensional 2D) layered materials such as graphene and semiconducting transition-metal dichalcogenides (TMDs) can exhibit peculiar electronic properties depending on their exact composition, thickness and geometry1D layered34678910111271572/graphene vdWHs by creating a device with two distinct regions containing one-layer (1L) and seven-layer (7L)-MoS2 to exclude extrinsic device factors. Significantly, we discover a surprisingly higher photoresponsivity and internal quantum efficiency (IQE) by going from ML MoS2 device to a monolayer device. The discretized electrostatic potential barriers were introduced to interpret the photocarrier tunneLling and extracting in ultrathin vdWHs.Here we report a layer-number-dependent photocurrent generation in graphene/MoS2/graphene heterostructure device on Si/SiO2 substrate. A focused laser was used to generate electron-hole pairs in the MoS2 layer, which can be separated by the asymmetric potential between the top graphene (GrT)/MoS2 and the bottom graphene (GrB)/MoS2 junction to produce photocurrent . The GrT layer is located directly on the MoS2 flake  to overlap with MoS2 flake and the GrB.The photocurrent generation in our devices was mapped by scanning photocurrent microscopy, where a focused laser beam was raster-scanned over the sample, while the photocurrent was being measured (Ids–Vds) data obtained in the dark  through ultrathin tunnelling barrier of 1L-MoS2ocurrent . Figure measured . The spaAeff, ϕB, q, m, m, d and h are effective contact area, barrier height, electron charge, free electron mass, effective electron mass, barrier width (MoS2 thickness) and Plank's constant, respectivelywhere −1) of the device. The photoresponsivity of 1L-MoS2 increases with decreasing the laser power could be attributed partly to absorption saturation in MoS2 and partly to the screening of a built-in electric field by the excited electrons in the conduction band of MoS2 (ref. 2 increases with increasing the laser wavelength and a maximum photoresponsivity of 68 mA W−1 in 1L-MoS2 vdWHs was achieved at the wavelength of 633 nm and the laser power of 100 nW (2/graphene vdWHs device (∼30 mA W−1 without plasmonic enhancement)72) and previous reports (5∼30 nm MoS2), which will be further discussed by comparison of external quantum efficiency (EQE) and IQE between 1L-MoS2 and 7L-MoS2 vdWHs in Based on the photocurrent response and input laser power, we can determine the photoresponsivity  and the GrT layer is located directly above the MoS2 flake (inside the solid line), to overlap with MoS2 flake and the GrB. The number of layers of MoS2 was then confirmed by Raman spectra measurements 152 (band gap ∼1.3 eV)To further unambiguously illustrate the difference between monolayer and ML MoSe device . The 8 μurements  and atom2 region exhibited much more pronounced photocurrent than the 7L-MoS2 region (Ids–Vds) data obtained in the dark . This is a rather surprising and counter-intuitive discovery, considering that the 7L-MoS2 should have significantly higher optical absorption than 1L-MoS2 , the second factor  dominates in our device.The enhanced photoresponsivity of 1L-MoSc) is thus given (in 2D) byIn general, the spectral dependence of the photocarrier generation rate can be estimated from first principles, using the Fermi golden rule. The probability of transition from a valence band state (υ) to a conduction band state . It is also noted that the IQE in 1L-MoS2 region steadily increases when the photon energy is reduced to approach the bandgap of 1L-MoS2 (1.82 eV (ref. 2 (1.82 eV). It is noteworthy that in this region k2 is nearly independent on the thickness. According to the theoretical calculations, 1L-MoS2 has a much larger density of states near the band edges, at the valence band top due to the heavier hole masses, and near the conduction band edge, due to the presence of the 2D van-Hove singularities . The hilarities , increas2 is driven by diffusion and drift process, which is dictated by different band bending at GrT and GrB MoS2 interface (originated from the difference of GrT and GrB doping densities)72 having a continuous band bending  and GrT (2.7) are determined by the standard approximation of ɛGrB=(ɛ21L-MoS+ɛ2SiO)/2=4.2 and ɛGrT=(ɛ21L-MoS+ɛair)/2=2.7, respectively22ɛ21L-MoS=4.5  can be determined by the difference of electron tunnelling probabilities between bottom and top junction barriers, which is normalized by the sum of tunnelling probability to top (TT) and bottom junction barrier (TB) (I)=(TT–TB)/(TT+TB)). The potential barrier was sliced to thin square barriers with width of 0.01 nm and then the tunnelling probability was calculated through overall square barriers  strongly relies on the excitation photon energy  is not altered appreciably in response to the energy of excited electrons (or absorbed photon energy), because effective barrier height (U–E) of the bottom and top junction changes by the similar magnitude  because of the decrease of effective barrier height (U–E), whereas the tunnelling probability to the TT remains almost at 1 regardless of the electron energy. Consequently, the photocarrier extraction rate (I) decreases in proportion to electron excitation energy  decreases in proportion to the electron excitation energy  were calculated in terms of the electrostatic potentials.The electrostatic potentials were calculated in terms of the number of MoS2 layers . The reg2) is shorter than that of indirect gap (7L-MoS2), which should be taken into account in recombination process during carrier drift. We have measured the carrier lifetime in direct gap (1L-MoS2) and indirect gap (7L-MoS2) based on transmission change in pump–probe measurement (δnx) by initial excited carriers (δ0) for 1L-MoS2  and ML-MoS2  along with carrier drift distance  rate into graphene layer, much shorter than carrier lifetime in 1L-MoS2 (60 ps), resulting in highly efficient transfer of photo-excited carriers in MoS2 into graphene before recombination. Therefore, we multiplied photocarrier extraction rate (I) in 1L-MoS2 by 100% of normalized excited carrier density. I) reaches around 65% when the number of MoS2 layers is reduced from 7L-MoS2 to 1L-MoS2 and the photon energy approaches the bandgap (1.82 eV). The calculated I are in good agreements with measured IQE of 1L-MoS2  and open circuit voltage (Voc) by an external gate voltage  of the GrB due to the relatively strong electrostatic screening by ML-MoS2  is approximately twice of that in single layer (MoS2), the IQE in our device is about twice higher. The lower IQE in in graphene/1L-WSe2/1L-MoS2/graphene p–n heterojunctions may be partially attributed to Shockly–Read–Hall recombination and Langevin recombination, which is not present in graphene/1L-MoS2/graphene devices.In summary, we have reported an unusually efficient photocarrier generation from monolayer MoS2/graphene device, the graphene was grown with a chemical vapour deposition process272 (300 nm SiO2) substrate, and patterned into 8 × 30 μm strips as the bottom electrode using a photolithography and oxygen plasma etching process. The MoS2 flakes were then transferred onto the graphene strips through a dry transfer approachT electrode was transferred and patterned on the MoS2 flake and GrB. Directly overlapping graphene area was patterned and etched away, to avoid short circuit between the GrT and GrB. The metal electrode, for probe contact or wire bonding purposes, was patterned on the bottom and GrT electrodes by e-beam lithography followed by e-beam deposition of Ti/Au (50/50 nm).For the fabrication of the vertical heterostructures of graphene/MoS2 heterostructures were excited with a cw solid-state laser at a wavelength of 405 and 532 nm. A low laser power of ∼100 μW was used to avoid heating and PL saturation. The PL of MoS2 samples were calibrated by using rhodamine 6G molecules as the standard. Raman spectroscopy  was used to characterize the MoS2 and graphene with a wavelength of 514 nm (2.41 eV) and a Rayleigh line rejection filter. As Raman scattering efficiency has usually little layer thickness dependence, the PL spectra normalized by Raman intensity reflects directly the luminescence efficiencyElectrical transport measurements were conducted with a probe station and a computer-controlled analogue-to-digital converter at room temperature. The scanning photocurrent measurements were conducted with the same electrical measurement system under a SP2 MP Film confocal microscope coupled with Ar/ArKr laser  and HeNe laser . All optical measurements were carried out under ambient conditions at room temperature by using an inverted microscope coupled to a grating spectrometer with a charge-coupled device camera. The optical beams were focused on the sample with a spot diameter of ∼1 μm. MoSet al.The DFT calculations were performed using the Quantum ESPRESSO codea) used was optimized for the primitive unit cell. The supercell length along the direction perpendicular to the plane c is taken to be twice the lattice parameter, except for preliminary bandstructure calculations performed using a smaller c, which are identified in the text. The Brillouin zone was sampled using a 2 × 2 × 1 Monkhorst–Pack gridThe monolayer was modelled using a supercell consisting of 4 × 4 unit cells of the monolayer material, comprising a total number of 48 atoms. The lattice parameter  and the delay between the two pulses is controlled by a mechanical delay stage (Newport M-IMS300PP).The femtosecond laser source is based on a 250 kHz Ti:sapphire regenerative amplifier (Coherent RegA 9050), which provides a 1.55 eV photon energy, 50 fs pulse width and 6 μJ per laser pulse. The Reg A output pulse is separated into two pulses by a ratio of 7:3. The 70 % laser pulse is used to generate 3.1 eV pump–photon energy by second harmonic generation in a 1 mm-thick beta barium borate crystal. The other 30% laser beam is focused in an ultra-clean sapphire disk to generate a white-light super-continuum, which serves as a probe pulse. Both the pump and the probe pulses are focused on the 1L-MoSThe data that support the findings of this study are available from the corresponding author upon request.How to cite this article: Yu, W. J. et al. Unusually efficient photocurrent extraction in monolayer van der Waals heterostructure by tunnelling through discretized barriers. Nat. Commun.7, 13278 doi: 10.1038/ncomms13278 (2016).Publisher's note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.Supplementary Figures 1-9, Supplementary Notes 1-6 and Supplementary References