However, heat flow at the interface between TMDCs and h-BN, which will play an important role in thermal management of various electronic and optoelectronic devices, is not yet understood. In this paper, for the first time, the interface thermal conductance (G) at the MoS2/h-BN interface is measured by Raman spectroscopy, and the room-temperature value is (17.0 ± 0.4) MW · m−2K−1. For comparison, G between graphene and h-BN is also measured, with a value of (52.2 ± 2.1) MW · m−2K−1. Non-equilibrium Green’s function (NEGF) calculations, from which the phonon transmission spectrum can be obtained, show that the lower G at the MoS2/h-BN interface is due to the weaker cross-plane transmission of phonon modes compared to graphene/h-BN. This study demonstrates that the MoS2/h-BN interface limits cross-plane heat dissipation, and thereby could impact the design and applications of 2D devices while considering critical thermal management.Two-dimensional (2D) materials and their corresponding van der Waals heterostructures have drawn tremendous interest due to their extraordinary electrical and optoelectronic properties. Insulating 2D hexagonal boron nitride ( Compared to the suspended configuration where heat transport only occurs in the basal plane, the difficulty with the supported configuration comes from the three-dimensional (3D) nature of heat transport. Even though the in-plane thermal conductivities of the two materials involved could be obtained individually by suspending each sample across a measurement platform, the thermal conductivity changes when the samples are supported due to the suppression of the flexural mode phonons and interactions between the substrate and superstrate3637Compared to the numerous studies1314222324252614282931323334363738et al. calculated G for metal-semiconductor interfaces such as Al/Si, Al/Ge, Cu/Si and Cu/Ge using exact phonon dispersionset al. measured G for an assembly of dissimilar materials using time-domain thermoreflectance (TDTR), exploring the lower bounds of G and found that such a lower bound actually lies within a narrow range of 8 to 30 MW/m2KFinite interface thermal conductance (G) was first observed by Kapitza at the copper and liquid helium boundaryet al. measured G between epitaxial graphene and 4H-SiC to be 0.0189 MW/m2K, which is five orders of magnitude lower than that expected from their molecular dynamics (MD) simulationset al. measured G between graphene and h-BN to be 7.4 MW/m2K2K2, Taube et al. measured the total G across MoS2/Si including a thin SiO2 layer to be 1.94 MW/m2Ket al. measured G to be 0.44 MW/m2K for MoS2/AuFor supported 2D materials, although heat is dissipated predominantly in the vertical direction as compared to laterally2/h-BN and graphene/h-BN, borrowing the technique from the work of Yue et al.et al.2 (graphene) through Joule heating, and at the same time, monitor the temperatures of MoS2 (graphene) and h-BN via Raman spectroscopy. The thermal conductance across the MoS2 and h-BN interface is (17.0 ± 0.4) MW/m2K, which is ~3 times smaller than G for graphene/h-BN interface measured as (52.2 ± 2.1) MW/m2K. With the expected lower in-plane thermal conductivity of MoS2 compared to graphene, this could indicate that electronic and optoelectronic devices using MoS2/h-BN have a larger propensity for failure due to overheating and therefore this value needs to be considered carefully in the future design of such devices.In this work, we directly measure and compare G between MoS2 sheets and a semi-infinite substrate of h-BN sheets. The graphene layers are stacked in the A-B configuration while the h-BN and MoS2 are stacked in the A-A’ configuration. In each structure, the interfacial area is 2.49 × 2.2 nm2. The interatomic potentials are obtained from literature5354To understand further the physics behind this difference in G, we conduct non-equilibrium Green’s function (NEGF) calculations of these two interfaces, which show a trend consistent with experiments. Our simulation consists of a semi-infinite superstrate of graphene or MoSGint, Gsuper and Gsub are respectively the conductance of the MoS2/h-BN (or graphene/h-BN), MoS2/MoS2 (or graphene/graphene) and h-BN/h-BN interfaces.where 2/h-BN FET device is shown in 2 and h-BN are shown in 2 is shown in 2E′ peak and the −1)2c) between MoS2 and the two inner electrodes, which are used for Joule heating.The fabricated MoS2 and h-BN. Raman spectroscopy has been demonstrated to be an effective method to probe the temperatures of 2D materials3133343740412 and h-BN as a function of temperature. At this stage, the temperatures of the materials are externally controlled via a heater stage, and their corresponding Raman spectra are collected at each temperature point. This calibration ensures that the temperature of each material can be determined by referring to their Raman spectra in the subsequent Joule heating measurement. Next, a series of current bias are applied to the electrically conducting MoS2 to provide Joule heat power given by Po = I2R . The laser spot is located in the middle of the sample. Note here that the relaxation time for electron-phonon interaction is on a short time scale of picoseconds with a relaxation length of the order of nanometersh-BN, while the sample stage is maintained at room temperature. All measurements are performed in a high vacuum environment in a closed-cycle cryostat (<10−5 torr). The Raman spectra of MoS2 and h-BN are monitored at each current bias to determine their temperatures based on the calibration results in the first part of measurement.To determine G, it is necessary to know the temperatures of the MoS2 and h-BN are shown in 2, the in-plane E′ mode is sensitive to strain, while the out-of-plane 2 to avoid interference from strain-induced Raman shifts. The relationship between the Raman shift and the temperature can be expressed as ω is the Raman shift due to temperature change, χT is the first-order temperature coefficient, and ΔT is the temperature change. The χT for the MoS2h-BN are measured to be 0.01212 cm−1 · K−1 and 0.01723 cm−1·K−1 for the first sample, and 0.01243 cm−1 · K−1 and 0.01778 cm−1 · K−1 for the second sample, respectively.The temperature calibration of the Raman spectra of MoS2 Joule heating process, the temperatures of the MoS2 and the underlying h-BN increase, resulting in red shifts of thr Raman peaks as indicated in χP, which is defined as −1 · mW−1 and 0.12697 cm−1 · mW−1 for the MoS2h-BN for the first sample, and 0.09972 cm−1 · mW−1 and 0.12719 cm−1 · mW−1 for the MoS2h-BN for the second sample.During the MoS2 and h-BN is calculated fromThe interface thermal conductance between MoST is the temperature difference beten MoS2 and h-BN.where ΔA is the interface area (including the transfer lengths from contacts) between MoS2 and h-BN, which is 50.50 μm2 and 73.05 μm2 for the two samples we measured. PO = P + Pc is the total Joule heating power, including both the power P from the MoS2 channel and Pc from the contact resistance Rc between MoS2 and the two inner electrodes used for Joule heating. Pc is determined through a 2-probe/4-probe resistance measurement, and T can be obtained based on the Raman measurements described above, i.e.,is the heat flux across the interface. Here, is found . ΔT can −2K−1 and 15.8 MW · m−2K−1 for the two samples.Combining (1) and (2), the interface thermal conductance G is calculated to be 20.1 MW · m2 and the influence from the contact resistance. A three-dimensional finite element method (3D FEM) using COMSOL is used to estimate the local (center of the sample) heat flux. This is realized by treating G as an input parameter and iteratively making it consistent with the local temperature difference observed from experiments  MW · m−2K−1. The G obtained for MoS2/h-BN interface here is larger than that of the MoS2/Au interface which has a value of 0.44 MW · m−2K−1 h-BN surface and/or better thermal coupling between MoS2 and h-BN intrinsically.To improve the accuracy of our measured G, we also consider the non-uniform heat flux within the basal plane of MoSeriments . The G fh-BN interface for comparison. Similarly, the measurement is composed of two parts, that is, the Raman shift calibration against temperature and the Joule heating measurement. The measurement results for the two parts are shown in h-BN interface, with an interface area of 43.55 μm2 and −2K−1 and 47.8 MW · m−2K−1 based on the graphene Raman G band and 2D band respectively. Similar to the MoS2/h-BN case, the local heat flux correction from 3D FEM COMSOL simulation gives 53.7 MW · m−2K−1 and 50.7 MW · m−2K−1 based on the graphene Raman G band and 2D band respectively, resulting in an average effective value of (52.2 ± 2.1) MW · m−2K−1. When observing the heterostructures with AFM, we find bubbles of size ~100 nm randomly dispersed within the interfaces of both MoS2/h-BN and graphene/h-BN h-BN than when it is covalent bonded to SiC, as tensile strain in graphene during the heating process tends to increase the red-shift of Raman peaks, resulting in the overestimation of the temperature of graphene and thereby underestimates G.Compared with the G of 7.4 MW · m2/h-BN interface, graphene/h-BN has a better interface thermal conductance. To better understand the physics behind this difference, we employ the NEGF technique to calculate the thermal conductance and transmittance spectrum. 2/h-BN and the graphene/h-BN interfaces. Our measurement results of G (17.0 ± 0.4) is close to that from NEGF calculation of 26.4 MW · m−2K−1, which serves as the theoretical upper bound of this interface, indicating the high quality of our interface. The measured G for the graphene/h-BN (52.2 ± 2.1) MW · m−2K−1 is larger than that for MoS2/h-BN, the same trend as predicted by NEGF calculations. From the phonon transmittance spectrum across the interfaces [2/h-BN interface is due to the lower phonon transmittance of MoS2, which has a much smaller overlap with the h-BN transmittance spectrum and thus, a smaller effective transmittance spectrum for the interface, as compared to the graphene/h-BN case. Thus, the much lower thermal conductance for the MoS2/h-BN interface is due to the phonon transmission bottleneck caused by the limited number of transmission channels in MoS2, compared to the greater number of transmission channels in graphene.Compared with the MoSterfaces , it can h-BN interface is larger than that of the MoS2/h-BN interface, it is still much smaller than the NEGF-calculated values. This may be due to tensile strain originating from thermal expansion in graphene, which tends to increase the Raman peak red-shift63642/h-BN interface, as we intentionally chose the strain-insensitive 2 as the temperature indicator. Therefore, we observed G for the MoS2/h-BN interface to be closer to theoretical predictions. Another explanation for the discrepancy is the dependence of G on the number of layers which has been shown in molecular dynamics simulations2 or graphene is Joule heated with the temperature higher than 300 K), and G should be temperature-independent above this temperature, based on NEGF calculations  modes whose contributions to the conductance are temperature-independent when T ≫ 100 K.Although G for graphene/ulations . For the2/h-BN and graphene/h-BN heterostructures. This is realized by Joule heating the MoS2 (graphene) superstrate, and monitoring the temperatures of both the superstrate MoS2 (graphene) and h-BN substrate through Raman spectroscopy. The measured thermal conductance of MoS2/h-BN interface is (17.0 ± 0.4) MW · m−2K−1, which is smaller than the value of graphene/h-BN interface (52.2 ± 2.1) MW · m−2K−1, consistent with the trend predicted from our NEGF calculations. We attribute the lower thermal conductance across the MoS2/h-BN interface to the limited cross-plane phonon transmission in MoS2, when compared to graphene. Our measurements and calculations provide basis for the study of in-plane thermal conductivity of MoS2 supported by h-BN, the typical structure for building high-mobility TMDC FET devices. The low interface thermal conductance deserves careful consideration for the design of electronic and optoelectronic devices based on such heterostructures.In summary, we have measured the thermal conductance across the interface of 2D MoS2 flake was exfoliated onto a 285 nm SiO2/Si substrate, and was confirmed to be monolayer from its Raman distance between the E′ and −1) photoluminescence peak (~1.84 eV) (h-BN flake was exfoliated onto another 285 nm SiO2/Si substrate. The MoS2 was then transferred onto the h-BN with polymethyl methacrylate (PMMA) as an assisted layer followed by acetone etching. The electrodes were patterned by standard electron-beam lithography (EBL) process, followed with thermal evaporation of Ti/Au with thickness of 5 nm and 75 nm, respectively. The sample was then annealed in-situ in vacuum at 400 K for 6 hours to release strains produced during the transfer process and ensure that the two materials are fully adheredh-BN heterostructure was prepared following the same procedure.The MoS1.84 eV) . The h-B2/h-BN, bubbles of the size of ~100 nm were found at the interface. We found that these bubbles are almost unavoidable during the transfer process and difficult to remove once formed. Generally, high temperature annealing (350 °C) tends to make small bubbles from larger ones for graphene/h-BNh-BN samples in Ar atmosphere at 350 °C for 3.5 hours, and specifically chose areas with fewer bubbles for device fabrication. For those regions without bubbles, we found the roughness of the surface, which is defined as the root mean square average of height deviations taken from the mean image data plane, is ~1.5 nm, with the main contribution from polymeric residues during the transfer process. These polymeric residues are on the top surface of MoS2, thus won’t affect the interface. We find no obvious difference in roughness for the surface of h-BN and the region of MoS2 without bubbles, indicating the good conformity of the interface.After measurements, the interface quality was characterized with AFM. For MoSHow to cite this article: Liu, Y. et al. Thermal Conductance of the 2D MoS2/h-BN and graphene/h-BN Interfaces. Sci. Rep.7, 43886; doi: 10.1038/srep43886 (2017).Publisher's note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.


Position of the G+ band is discussed in relation with nonuniform doping of graphene with charge impurities. Simultaneous presence in the Raman spectra of heavily irradiated graphene of rather narrow G band and broaden D band is explained by coexistence within the Raman probe of more and less damaged graphene areas. This assumption is additionally confirmed by confocal Raman mapping of the heavily irradiated area.Confocal micro-Raman spectroscopy is used as a sensitive tool to study the nature of laser-induced defects in single-layer graphene. Appearance and drastic intensity increase of D- and D′ modes in the Raman spectra at high power of laser irradiation is related to generation of structural defects. Time- and power-dependent evolution of Raman spectra is studied. The dependence of relative intensity of defective D- and D′ bands is analyzed to relate the certain types of structural defects. The surface density of structural defects is estimated from the intensity ratio of D- and G bands using the D-band activation model. Unusual broadening of the D band and splitting of the G band into G Number of defects and their type is largely dependent on the technology of graphene production  and can + bombardment [Resonance Raman spectroscopy (RRS) plays a major role in investigation of carbon-based materials being quite effective, rapid, and nondestructive method , 20. Thebardment , 23, flubardment , and lasbardment , 25. In bardment .The present paper uses confocal micro-Raman spectroscopy as a sensitive tool to study the nature of defects in single-layer graphene induced by laser irradiation at varied laser power densities.2 for visual observation by means of optical microscopy  process [Raman spectrum of pristine graphene sample measured at laser powers ≤800 kW/cmntensity , 26. App process . It shou−1 is observed in the Raman spectra of the graphene exposed to constant laser power of 800 kW/cm2  for laser power densities of 1400–1600 kW/cm2 agrees with corresponding dependence for grain-boundary defects  the ID to ID′ ratio have a tendency to increase toward the corresponding values found for vacancy-related and sp3-type defects  and El is the energy of exciting laser radiation (eV). Using relation (1), it was found that the average distance between structural defects gradually decrease from ~65 nm down to ~20 nm when laser power density increase from 1400 up to 2000 kW/cm2 at constant exposition time of 1 h %, that is sufficiently lower than those concentrations commonly used in the modeling of structural defects in graphene [where 014/πLD2  correspographene –33.2 are demonstrated in Fig. −1, denoted as G− and G+ in Fig. + to G− band gradually increases with the exposition time, whereas their frequency position remains almost invariable (Fig. −1/%, correspondingly [−1 [+ component of the G band.To study further the evolution of Raman characteristic bands with increase of disorder, we provided time-dependent Raman measurements of the graphene at higher power density and exposition times. The results of these measurements at power density of 1600 kW/cmble Fig. . The obsondingly , but in ngly [−1 , which c2 are demonstrated in Fig. Raman spectra of the single-layer graphene after laser exposure with power density of 2400 kW/cm− component and 2D band (Fig. + component of the G band. It should be noted that the spatial distribution of Raman D-band intensity in Fig. + component, and less damaged area near the edge of the irradiated area, which is reflected by increase in intensity of the G band and relative intensity of the 2D band.As can be seen from AFM image Fig. a, the siand Fig. c, and de2 is related to generation of structural defects. Time- and power-dependent evolution of Raman spectra is studied. The relative intensities of defective D and D′ bands for laser power densities less than 1600 kW/cm2 are shown to agree with corresponding dependence for grain-boundary defects and have a tendency to increase toward the corresponding values typical for vacancy-related and sp3-type defects at higher power densities. The surface density of structural defects is estimated from the intensity ratio of D and G bands using the D-band activation model. Unusual broadening of the D band and splitting of the G band into G− and G+ components with redistribution of their intensities is observed during long-term laser exposure at powers densities higher than 1600 kW/cm2. Position of the additional G+ band is discussed in relation with nonuniform doping of graphene with charge impurities. Simultaneous presence in the Raman spectra of heavily irradiated graphene of rather narrow G band and broaden D band can be explained by coexistence within the Raman probe of more and less damaged graphene areas. This assumption is additionally confirmed by confocal Raman mapping of the heavily irradiated area.Appearance and increase in intensity of D-like modes in the Raman spectrum of single-layer graphene at laser power ≥800 kW/cm


Intradermal injection of treated sarcoidosis spleen tissue resulted in a granuloma response at injection site by 4–6 weeks. Previous work indicates proteins as the possible trigger of this reaction. We aimed to identify Kv-specific proteins and characterise the Kv extracts were separated by 1D-SDS-PAGE and 2D-DIGE and then underwent mass spectrometric analysis for protein identification. Sarcoidosis and control PBMCs were first stimulated with Kv and then with three selected recombinant protein candidates which were identified from the proteomic analysis. PBMC secreted cytokines were subsequently measured by Multiplex Cytokine Assay.p = 0.0018; TNF-α: 2375 pg/mL vs. 42.82 pg/mL, p = 0.0003). Through proteomic approaches we then identified 74 sarcoidosis tissue-specific proteins. Of these, 3 proteins  were identified using both 1D-SDS-PAGE and 2D-DIGE. Data are available via ProteomeXchange with identifier PXD005150. Increased cytokine secretion was subsequently observed with vimentin stimulation of sarcoidosis PBMCs vs. tuberculosis PBMCs . This finding was also observed in vimentin stimulation of sarcoidosis PBMCs compared to PBMCs from healthy controls . No difference was found in cytokine secretion between sarcoidosis and control PBMCs when stimulated with either tubulin or alpha-actinin-4.We observed significantly increased IFN-γ and TNF-α secretion from Kv-stimulated PBMCs of sarcoidosis patients vs. PBMCs from healthy volunteers (IFN-γ: 207.2 pg/mL vs. 3.86 pg/mL, Stimulation with both Kveim reagent and vimentin induces a specific pro-inflammatory cytokine secretion from sarcoidosis PBMCs. Further investigation of cellular immune responses to Kveim-specific proteins may identify novel biomarkers to assist the diagnosis of sarcoidosis. Sarcoidosis is a multi-organ granulomatous disease of unknown cause which occurs in genetically susceptible individuals  but primin vivo skin assay called the Kveim test, was used for diagnosis with sensitivity >70% and specificity >90% [in vivo human studies.Diagnosis of sarcoidosis is complex and relies on a supportive clinical history, radiology and biopsy exhibiting non-caseating granulomas. This approach is resource-heavy and merely suggestive of disease through exclusion of differential diagnoses, rather than specifically diagnosing sarcoidosis . Historiity >90% . Kveim rDespite extensive clinical validation, there has been limited successful research into the triggers of the Kv reaction. A sequential removal of lipids and oligosaccharides did not alter the granuloma-causing capacity of Kv whereas concentration of proteins improved sensitivity, suggesting the cause is likely protein-driven . ImmunolOne previous proteomic analysis of sarcoidosis solid tissue did identify the mycobacterial protein mKatG within Kv . A furthin vivo Kv-induced granuloma at 4–6 weeks would also be detectable ex vivo in peripheral blood. We aimed to define the proteomic signature of Kv itself and to characterise the nature of the ex vivo immune response to both Kv and selected identified Kv-specific proteins.We postulated that early antigen-driven immune responses contributing to the generation of the This study was approved by the St. Mary’s institutional ethics committee (reference: 07/H0712/85) and blood was obtained from participants after providing written informed consent. All sarcoidosis tissue was collected under the same ethical agreement.Sarcoidosis patients were chosen who had recent biopsy-proven pulmonary disease and were not on immunosuppressive therapy; diagnosis was obtained as per ATS guidelines . TubercuSarcoidosis spleen tissue and control spleen was provided by National Disease Research Interchange . Validated Kv was provided by Alvin Teirstein and Porton Down Institute. The method for the preparation of Kv follows the original protocol exactly . For PBM5 PBMCs in a total volume of 100 μL of AIM V media (Sigma-Aldrich) were placed per well in flat bottom plates (Thermo-Fisher). For stimulation, PBMCs were either supplemented with 20 μL of Kv or recombinant protein (20mg/ul of RPMI). Plates were incubated for 36 hrs at 37°C and 5% CO2 and spun at 500 g for 1 min at 21°C. Supernatant was aliquoted for immediate analysis.2.5 x10The multiplex cytokine analysis was performed using pre-coated 96 well plates  according to the manufacturer instructions and as previously described . 25 μL oProtein extracts of Kv were quantified  and 4 μg loaded onto 4–20% polyacrylamide gels , as reported previously . Gels weExperiments were conducted as previously outlined . 50 μg oEach gel lane of a 1D-SDS-PAGE gel was cut into 10 sections and selected gel spots of the 2D-DIGE gels were also excised with a scalpel. 1D-SDS-PAGE gel slices and 2D-DIGE gel spots underwent subsequent in-gel tryptic digest as follows: in brief, gel pieces were washed in 30 μL 50% acetonitrile (ACN) for protein destaining. Gel pieces were dried in a SpeedVac for 10 min, reduced with 15 μL 10mM DTT (Sigma-Aldrich) in 10 mM ammonium bicarbonate pH 8.0 (AmBic) for 45 min at 50°C and then alkylated with 15 μL 50 mM IAM (Bioultra-Sigma-Aldrich) in 10 mM AmBic for 1 h at room temperature in the dark. Gel pieces were washed three times in 30 μL 50% ACN, vacuum-dried before being reswollen with 50 ng of modified trypsin  in 5 μL 10mM AmBic and eventually overlaid with 10 μL 10mM AmBic. After incubation for 16 h at 37°C peptides were extracted twice with 10 μL 5% TFA (Sigma-Aldrich) in 50% ACN, vacuum-dried, resuspended in 8.0 μL of 0.1% (v/v) formic acid and stored at -80°C until further use.m/z 250 to 3,000 and at five spectra per second. Precursor ions were selected for automatic tandem MS (MS/MS) at an absolute threshold of 2,000 and a relative threshold of 0.01, with a maximum of five precursors per cycle.Mass spectrometric identification was conducted as previously described  using elResulting MS/MS data was analysed using MassHunter Mass Profiler software version B.02.01 (Agilent Technologies). Data were then searched with MASCOT Version 2.2.07  against UniProtKB/Swiss-Prot Database Release 2011_11  for the DIGE dataset or against UniProtKB/Swiss-Prot Database Release 2012_01  for the Kveim dataset. Database search was performed with settings as follows: taxonomy filter was set for human entries, MS tolerance (Parent Tolerance) ±20 ppm and the MS/MS tolerance (Fragment Tolerance) ±0.20 Da. Trypsin was set as the digesting enzyme, two missed cleavages of the digest were accepted and carbamidomethylation of cysteines (C) was set as fixed modification. Acetylation of peptide N-termini and lysine residues (K), methionine oxidation (M), deamidated (NQ) and Gln-pyroGlu (N-term Q) as well as amidated C-termini, carbamylated lysine residues and carbamylated N-termini were selected as variable modifications.Search result filters were set to MudPIT scoring, permitted peptides with a score >20 and below the Mascot significance threshold filter of p = 0.01 as valid identification. Resulting dat.-files were then imported into Scaffold Version 3.4.7 (DIGE data) or 3.6.7 (Kveim data)  and an additional database search with the X! Tandem algorithm (Version CYCLONE 2010.12.01.1) was performed with equally stringend search criteria. Besides aforementioned modification pyro-cmC (S-carbamoylmethylcysteine cyclization) on N-termini was added as another variable modification. Here, data are presented in Scaffold that were confidently identified by the two independent search engines i.e. Mascot and X! Tandem and further fulfil the Scaffold implemented criteria of 99% minimum protein probability score, 95% minimum peptide probability score with at least two identified peptides. The mass spectrometry proteomics data have been deposited to the ProteomeXchange Consortium via the PRIDE  partner Monoclonal mouse anti-Vimentin, , was applied as primary antibody at a 1:50 dilution on 8 μm cryosections of snap frozen spleen tissue for 30 min at room temperature. Dako Cytomation Mouse lgG'1, , diluted to the same mouse lgG concentration as the primary mouse anti-Vimentin antibody was used as negative control and run simultaneously with sarcoid spleen sections. For visualization, DAKO LSAB2 kit (Code No. K 0679) and DAKO EnVision+/HRP kits, (Code Nos. K 4004 and K 4006) were used on the Bond autostainer for automated IHC staining.For each direct comparison (e.g. measurement of secreted interferon gamma after vimentin stimulation of sarcoidosis PBMCs vs tuberculosis PBMCs) a p-value was calculated using a Mann-Whitney test. Data analysis was performed using GraphPad Prism version 6 for Macintosh .p = 0.0018) or sarcoidosis PBMCs stimulated with cKv  as seen in PBMCs from sarcoidosis patients and healthy volunteers were stimulated with both sarcoidosis Kv (sKv) and healthy control spleen prepared in a manner similar to Kveim (cKv). The secreted cytokine signature was then analysed. The median IFN-γ concentration was significantly elevated in the cell culture supernatant of PBMCs from patients with active pulmonary sarcoidosis  stimulatp = 0.0003) or sarcoidosis PBMCs stimulated with cKv . This pattern was reproducible using biological replicates of sKv and cKv, derived from different spleens  and healthy volunteer PBMCs   .TNF-α levels were also significantly elevated in sarcoidosis (1139 pg/mL IQR 51.53–2884) vs. tuberculosis  and healthy volunteer PBMCs  . No diffWe identified a pro-inflammatory cytokine response in PBMCs from sarcoidosis patients when stimulated with Kveim reagent (Kv). This response was not observed in PBMCs from healthy volunteers stimulated with Kv nor in sarcoidosis PBMCs stimulated with control Kv made from healthy spleen tissue. Using two complementary proteomic approaches we identified 74 proteins specific to sarcoidosis tissue and found that one, vimentin, induced the similar pattern of cytokine secretion from sarcoidosis PBMCs. This response was not seen with vimentin stimulation of PBMCs from healthy volunteers or those with pulmonary tuberculosis.in vivo experimental model, understanding the effect of Kv in granuloma formation is difficult to assess. However, Kv induced granulomas have histologically been detected at 6–24 hrs and resulted in a CD4+ and CD14+ monocyte assembly at 48 hrs after intradermal inoculation [In the absence of an culation .in vivo granuloma causing capabilities after lengthy storage e.g. >5 years at 4°C [1D-SDS-PAGE coupled to MS/ MS was used to define the Kv proteome and identified 48 proteins specific to sKv and historical vKv. It is noteworthy, that clinical studies have shown that Kv can retain s at 4°C . ComparaThese two complementary proteomic methods found three proteins in common to both approaches: vimentin, tubulin and alpha-actinin-4. Vimentin is a member of the type III intermediate filaments which take part in cell motility and adhesion, subcellular organization and maintenance of cell shape . It is fThe identification of Kv-specific cytoskeletal proteins such as vimentin, tubulin and alpha-actinin-4 might simply reflect increased abundance within granulomatous spleen tissue. However, upon stimulation of sarcoidosis PBMCs with these three candidate proteins, we found that vimentin could induce similarly, a strong and specific IFN-γ and TNF-α secretion as we had observed with Kv itself. A crucial finding is that vimentin did not elicit a pro-inflammatory cytokine response in PBMCs from tuberculosis patients, despite the substantial accumulation of vimentin within tuberculosis granulomas .Rheumatoid arthritis, like sarcoidosis, is a disease where a heightened immune response occurs to a panel of inflammatory stimuli. Intriguingly, a citrullinated form of vimentin is known to be elicit an immune response in rheumatoid arthritis, which correlated with disease activity , 27. NevWe could not demonstrate a vimentin-specific response in all sarcoidosis patients. In this respect, it is also important to note that there was no clinical difference  between vimentin responders and non-responders.Our findings support those from recent studies with a specific regard to a role for vimentin in the disease. One group identified a number of vimentin peptides by mass spectrometric sequencing after affinity-purification of HLA-DR molecules (HLA-DRB1*0301) from bronchoalveolar lavage (BAL) cells of sarcoidosis patients. Vimentin-peptides were amongst the identified peptides, suggesting antigenic properties of vimentin in the sarcoidosis lung . Two sub+ T-cells on a cohort of DRB1*03+ sarcoidosis patients and molecular modelling indicated a specific T-cell receptor–HLA- DRB1*03+- vimentin peptide (Vim429-443) interaction [For the sarcoidosis Kv-reaction, one potential hypothesis is that vimentin merely upregulates cytokine responses from sarcoidosis PBMCs as it is known that other proteins such as mycobacterial-derived antigenic peptides  and serueraction .Clinically validated Kv has been historically used to diagnose sarcoidosis with relatively high sensitivity and specificity. Our data, together with that from the most recent publication , suggestInterestingly, Kv has also recently been shown to stimulate pro-inflammatory cytokine production of PMBCs from HIV positive patients without sarcoidosis . This unOur study has some limitations. Firstly, we have only investigated patients with newly diagnosed pulmonary sarcoidosis. It would be interesting to replicate this work in patients with chronic or quiescent disease as there is some suggestion that Kv skin responses might diminish with time . Althougex vivo Kv stimulation of sarcoidosis PBMCs elicits a cytokine secretion not seen in controls. Perhaps more importantly, measurement of cytokine secretion in response to one of our identified proteins, vimentin, is shown here to be significantly greater from sarcoidosis PBMCS than those from patients with tuberculosis. Both diseases can present in similar ways and a definitive diagnosis often requires histological sampling. Invasive biopsies are time consuming, expensive and associated with significant morbidity. Our data also represents the first time that self proteins have been identified from within Kveim reagent. A growing number of papers have recently hypothesized that autoimmunity may play an important role in sarcoidosis [Current diagnosis of sarcoidosis relies on subjective information with patients often being treated with immunosuppressive therapies without a definitive histological diagnosis. We have shown that an coidosis  and we bS1 Table(DOCX)Click here for additional data file.


Idiopathic pulmonary fibrosis (IPF) is a lethal lung disease of unknown etiology. Patients present loss of lung function, dyspnea and dry cough. Diagnosis requires compatible radiologic imaging and, in undetermined cases, invasive procedures such as bronchoscopy and surgical lung biopsy. The pathophysiological mechanisms of IPF are not completely understood. Lung injury with abnormal alveolar epithelial repair is thought to be a major cause for activation of profibrotic pathways in IPF. Metabolic signatures might indicate pathological pathways involved in disease development and progression. Reliable serum biomarker would help to improve both diagnostic approach and monitoring of drug effects.The global metabolic profiles measured by ultra high-performance liquid chromatography coupled to high-resolution mass spectrometry (UHPLC-HRMS) of ten stable IPF patients were compared to the ones of ten healthy participants. The results were validated in an additional study of eleven IPF patients and ten healthy controls.We discovered 10 discriminative metabolic features using multivariate and univariate statistical analysis. Among them, we identified one metabolite at a retention time of 9.59 min that was two times more abundant in the serum of IPF patients compared to healthy participants. Based on its ion pattern, a lysophosphatidylcholine (LysoPC) was proposed. LysoPC is a precursor of lysophosphatidic acid (LPA) – a known mediator for lung fibrosis with its pathway currently being evaluated as new therapeutic drug target for IPF and other fibrotic diseases.We identified a LysoPC by UHPLC-HRMS as potential biomarker in serum of patients with IPF. Further validation studies in a larger cohort are necessary to determine its role in IPF.NCT02173145 at baseline and from the idiopathic interstitial pneumonia (IIP) cohort study. The study was approved by the Swiss Ethics Committee, Bern (KEK 002/14 and 246/15 or PB_2016–01524).Serum samples from IPF patients have been obtained within the clinical trial The online version of this article (doi: 10.1186/s12931-018-0714-2) contains supplementary material, which is available to authorized users. Idiopathic pulmonary fibrosis (IPF) is a severe and progressive fibrosing interstitial lung disease with sinister prognosis. The median survival is about 3–5 years after onset of symptoms, although there are various disease progression patterns –3. PatieIPF is diagnosed in a multidisciplinary setting with typical patterns in high resolution computed tomography (HRCT) and, if necessary, invasive methods as bronchoscopy and lung biopsy . The diaNowadays, global metabolic profiling using ultra high-performance liquid chromatography coupled to high-resolution mass spectrometry (UHPLC-HRMS) offers the possibility to discover pathologically regulated molecules. These molecules might give information about aberrant metabolic pathways in a disease. Characteristic metabolites might offer the opportunity to understand pathological processes and to discover new potential diagnostic and/or prognostic biomarkers . To our n = 10, n = 11 respectively) and control subjects  were obtained from the Department for Pulmonary Medicine, University Hospital, Inselspital, Bern, Switzerland and stored by −80 °C until analysis. IPF was diagnosed according to the current international guidelines of 2011 +, [M + K]+, [M + H-H2O]+, [M + H-2H2O]+, [M + 2H]2+, [M + 3H]3+, [2M + H]+, [2 M + Na]+ and [2M + K]+ were defined in the positive ion mode, [M-H]−, [M-H-H2O]−, [M + HCOOH-H]−, [M-2H]2−, [M-3H]3− and [2M-H]− in the negative ion mode. The abundances of the features were normalized to all compounds. Features eluting between 0.5–10.5 min with a chromatographic peak width ≥ 0.05 min, having a coefficient of variation of the peak area ≤ 30% in the QC samples, singly charged ions and ions having a maximum abundance across all samples ≥200  or ≥150 , were selected for further evaluation. In addition, the metabolic features having an abundance in blank higher than in QC and which were at least twice more abundant in blank were excluded. Known drug/drug metabolites taken from patient were also excluded from the dataset.After lock-mass correction and chromatographic alignment with Progenesis QI , ion patterns were deconvoluted between 0.45–11.5 min. As possible ions, [M + H]2(cum) was ≥0.5, if the cross validatory ANOVA p-value was ≤0.05 and if the permutation test was passed . Discriminative features were selected based on a VIPPRED score ≥ 1.5 and absolute p (corr) ≥ 0.5. The confidence interval of the selected features was also not including 0. With this approach, only metabolic features with strong model contribution and high reliability were selected. Univariate statistical analysis was performed to consistently evaluate the filtered metabolic features and to screen for additional metabolites. For the pilot study, univariate analysis of variance (ANOVA) was based on arcsinh transformed data from Progenesis QI. For the validation study, ANOVA was performed with SIMCA and was based on log2 transformed and Pareto scaled normalized abundances. The false discovery rate (FDR) due to multiple testing was estimated according to Benjamini Hochberg − ions in the positive and negative ionization mode, respectively.Structure and formula of the isolated discriminative metabolic features were searched against the Human Metabolome Database  and LIPID MAPS Structure Database  with a mass accuracy of 8 ppm. Metabolic features with only one adduct ion were assumed to be [M + H]In the pilot study, 10 IPF patients and 10 age- and sex-matched healthy controls were included. The mean age of IPF patients and of healthy controls was 67.8 respectively 68.7 years. Ninety percent of all participants were male in both groups. Lung functional measurements showed a reduced median forced vital capacity (FVC) of 2.4 L (mean 65% predicted) in IPF patients and a median FVC of 4.1 L (mean 115% predicted) in controls. Average 3-year mortality prediction in IPF patients was in 42.1% according to the currently used GAP staging method (stage II) + at m/z 85.0284 was confirmed by UHPLC-TWIM-MS measurement . Presumably, these co-eluting ions belonged to the same metabolic features. In the negative ESI mode, the detected ion pattern agreed with a phosphatidylcholine (PC) or sphingomyelin [M-CH3]− and [M + CHOO]− ions. These ion patterns were confirmed by the analysis of PC lipid standards (data not shown) and agreed with the literature + 3-hydroxydecanoyl carnitine ion. The detection of the characteristic acylcarnitine fragment [C4H5O2]+ at m/z 85.0284 was confirmed by MS/MS spectra . (PDF 99 kb)MS MS/MS spectra (DDA experiments) of the metabolic features eluting at 9.6 min acquired in the (a) positive and (b) negative ESI mode and potentially assigned to the [M + H]


We here demonstrate a strategy for enhancing the CTE by a two-step optimization of graphene’s surface energy prior to AuCl3 doping. First, exposure to UV ozone modified the hydrophilicity of graphene and was found to decrease the cluster’s geometric capacitance, which had a direct effect on the CTE. Occurrence of lattice defects at high UV exposure, however, deteriorated graphene’s transport characteristics and limited the effectiveness of this pretreatment step. Thus, prior to UV exposure, a functionalized polymer layer was introduced that could further enhance graphene’s surface energy while protecting it from damage. Combination of these treatment steps were found to increase the AuCl3 charge transfer efficiency to 70% and lower the sheet resistance to 106 Ω/γ at 97% transmittance which represents the highest reported performance for doped single layer graphene and is on par with commercially available transparent conductors.Graphene’s attractiveness in many applications is limited by its high resistance. Extrinsic doping has shown promise to overcome this challenge but graphene’s performance remains below industry requirements. This issue is caused by a limited charge transfer efficiency (CTE) between dopant and graphene. Using AuCl Despite significant advances in the synthesis of high quality material, graphene’s performance is still below the industry standards for transparent conducting films (TCFs)4. This shortcoming is caused by graphene’s low intrinsic carrier concentration due to its unique band structure5. A common method to enhance graphene’s conductivity is through extrinsic doping where oxidizing or reducing agents are introduced to remove or add charges to the graphene. This method combines easy scalability and high performance compared to alternative approaches such as graphene/mesh hybrids6 and intercalation compounds7. A wide variety of materials have been employed for extrinsic doping, such as nitric acid8, silver nitrate9, iron chloride7, gold chloride10 and ammonia gas11. While these materials enhance the carrier concentration of graphene, they also decrease graphene’s transparency. To capture the simultaneous variation in transmittance and resistance, a figure of merit (FOM) is commonly employed12 that represents the ratio of graphene’s conductivity at DC and optical frequencies and should be maximized.12 and depends on the dopant’s molecular light extinction coefficient ε and graphene’s carrier mobility μG according to3)17, we found no report of an FOM larger than 50 which does not satisfy requirements for many current-driven applications18. The question arises if there is a fundamental limit for the FOM as previously suggested12 or if optimization of the doping process could increase the FOM enough to compete with traditional TCF materials.Graphene, a two-dimensional carbon allotrope, is anticipated to be an enabling material for flexible and transparent electronics3 charge transfer efficiency on surface energy was found to be due to geometry-induced work function changes of the dopant. To maximize the CTE, we devised a multistep pretreatment that enhances graphene’s surface energy. Exposure to UV-generated ozone showed a clear trend between graphene’s hydrophilicity and the amount of transferred charge which was ascribed to the formation of functional groups in the graphene basal plane. The extent of CTE-increase is limited by the onset of destructive oxidation of graphene at prolonged UV exposure. To minimize the impact of defects on carrier transport, a polymer layer was introduced on top of the graphene and functional groups were produced in this film. Enhanced thermal scission was found to further enhance graphene’s surface energy and thus treated graphene showed an improved CTE of 70% and exhibited the highest reported figure of merit for doped graphene in excess of 110 while retaining a transparency of 97%. These features make the presented approach promising for industrial applications of graphene-based transparent conducting films.Here we demonstrate that control over graphene’s surface energy can enhance the efficiency of charge transfer between dopants and graphene. A strong dependence of the AuCl19. Briefly, copper foil is annealed at 1000 °C under 10 sccm hydrogen gas flow for 30 minutes, then graphene is grown under a flow of 40 sccm methane and 100 sccm hydrogen gas flow for one hour before being cooled down to room temperature under a flow of 10 sccm H2. After growth, graphene is transferred onto quartz or SiO2/Si substrates using established procedures20.Single layer graphene was grown by chemical vapor deposition on copper foil(Alfa Aesar 31882) following previous reports3 solution was produced by dissolving AuCl3 (334049 Sigma-Aldrich) in Nitromethane solution following previous reports17. This solution was cast onto a graphene device at room temperature using a micropipettor and left overnight for drying. Poly(methyl methacrylate) (Microchem A4) was spin coated at 2500 rpm for 1 minute. Spin curves predict a film thickness of approximately 250 nm under these conditions. UV illumination was conducted in a homebuilt chamber using a light source with a power of 13.3 mW.AuClRaman spectroscopy  was carried out using a 532 nm laser source. The sample’s sheet resistance and Hall carrier concentration were measured in van-der-Pauw geometry using an Agilent B2900A source meter and a homebuilt probe station. Work function was measured through Kelvin probe measurements. Haze measurements were carried out in an integrating sphere using an Evolution 220 Spectrophotometer.3 solution on graphene. This process16 is expected to produce a reactive species of AuCl4- ions that remove electrons from the graphene resulting in p-type doped graphene and neutral gold clusters according toWe carry out doping by casting 100 µl of AuCl3 concentration shows a decrease in sheet resistance in agreement with previous reports16 . This observation indicates that at high AuCl3 concentrations charge transfer between graphene and AuCl3 proceeds until graphene has reached the same work function as the dopant. However, for lower AuCl3 concentrations, a lower amount of available charges from AuCl3 will only raise graphene’s work function to an intermediate level.Surprisingly, we find a clear dependence of the achievable sheet resistance on the concentration of AuClion Fig. . To unde4- ions (1 V) that form during the doping process16. Based on this simple picture the highest achievable work function shift upon reduction is 0.78 eV which agrees well with the experimental value. Considering this theoretical reduction potential difference we can estimate a doping concentration of 3 and graphene. We can thus quantify the charge transfer efficiency (CTE) as the fraction of actual chare transfer compared to the complete transfer case.For high concentrations the work function change is found to saturate at approximately 0.7 eV. To identify the physical meaning of this value, we investigate the difference in reduction potential between graphene (0.22 V) and AuCl3 at similar transparencies which highlights the importance of this parameter for enhancing the performance of graphene. In order to identify the origin of the low CTE at low carrier concentrations, we carried out AFM imaging of samples after AuCl3 treatment. We find that higher AuCl3 concentrations form larger clusters on graphene  will exhibit a higher amount of transferred charges Fig. . This re28. Previous reports showed that the UV exposure of graphene will produce functional groups, such as hydroxyl groups29, that enhance its surface energy30 but such functional groups simultaneously increase the carrier scattering31. We indeed observe that graphene’s carrier mobility decreases significantly upon prolonged UV exposure  (PMMA). The polymer is known to preserve graphene’s high carrier mobility32 and can be decomposed into short oligomers by thermal annealing33. For this purpose, annealing of the substrate/graphene/PMMA stack at high temperatures in a forming gas of 600sccm H2 and 400sccm Ar for 150 minutes was carried out. Raman spectroscopy shows the occurrence of a broad peak around the D-band that had been previously identified as a fingerprint of fragmented PMMA after thermal treatment (Fig. 34. AFM images show little residue after annealing (Suppl. Figure 34.The increasing hydrophilicity in graphene had been related to the splitting of water on graphene under UV exposure and this process was found to preferentially occur on the sites of lattice defects in graphene34. Thus, the charge induced in the graphene represents a measure of the radical concentration. We find that the amount of induced charge increases with annealing temperature which supports the model of temperature induced functionalization (Fig. 31. If the annealing temperature exceeded 300 °C, however, the carrier concentration stagnated while the carrier mobility continued decreasing, resulting in a horizontal line. This behavior suggests the formation of lattice defects in graphene which deteriorates carrier transport. This conclusion is supported by measurements of graphene’s resistance change after annealing. Compared to pristine graphene, annealing can reduce graphene’s resistance if temperatures below 300 °C are chosen (Fig. 2 from 70° to 10° and demonstrates the hydrophilic character of the oligomers.To investigate the effect of temperature on bond-scission, graphene’s carrier concentration was analyzed. Upon scission, a charged radical occurs at the end of the remaining chainion Fig. . The enhsen Fig. . AnnealiIn-situ measurements show a qualitatively similar evolution of the graphene mobility under UV exposure with and without the presence of the polymer (Suppl. Figure 3 is deposited on the thus prepared structure. We observe that this sequence is the only one that lowers the sheet resistance at every process step (Fig. We now combine all the introduced pretreatment steps (Suppl. Figure tep Fig.  which suEach step was found to significantly affect the work function of graphene as directly measured by a Kelvin probe. These results suggest that an enhanced CTE is the origin of the conductivity increase Fig. . Under o3 doping without surface pretreatment as shown in Fig. 3 doping result. Furthermore, the achieved FOM represents the highest reported value for any doped single-layer graphene sample and is on par with triple-layer graphene champion devices. We furthermore carried out haze measurements and found that the average increase in haze for samples before and after AuCl3 doping is 1.1% and thus retains the advantageous properties of graphene-based transparent electrodes.The potential of our approach is highlighted by the observed enhancement in the figure of merit. Graphene prepared by the describe sequence of treatment steps exhibits a sheet resistance of 106 Ω/γ at a transmittance of 97% (See Supplementary Figure 3-doped graphene’s figure of merit, and produced doped single layer graphene with the highest reported performance. The presented approach of surface energy control for enhanced doping is applicable to many dopants and opens up a new route to increasing graphene’s potential for optoelectronic applications.In conclusion, we have demonstrated the enhancement of the graphene doping process by optimization of the charge transfer efficiency. This parameter was found to be limited by the geometric capacitance of small clusters and we designed a sequence of pretreatment steps that maximizes the dopant cluster size. Exposure to UV light was shown to increase graphene’s surface energy through introduction of functional groups and directly affect the CTE. To overcome the deterioration of the carrier transport by scatterers in the graphene lattice a sacrificial polymer layer was used as an anchor for functional groups. Combination of these steps increased the CTE to 70%, doubled AuClSupplementary material


Controlling the loss of graphene can be used in the field of transformation optics. We propose a new concept of electromagnetic fence on a monolayer graphene surface plasmon polariton platform. Using a Dot-Density-Renderer quasicrystal metasurface, we can simulate the absorption of gradient index optics structures. Numerical simulations show that the incident waves to our designed electromagnetic fence are trapped toward the central lines and quickly absorbed by the high-loss region. Two basic types of electromagnetic fence and its composite structures have been designed and analyzed, which exhibit excellent broadband absorbing performances at 8 THz–12 THz. Because of its advantages in controlling the soft-boundary effects and easy manufacturing characteristics, the proposed electromagnetic fence seems very promising for THz–frequency-transformation plasmonics applications. For this purpose, graphene has been investigated as a one-atom-thick platform for terahertz devices due to its ability to support and control surface plasmon polariton (SPP) surface waves in the terahertz regime4. Unlike those in noble metals, SPP surface waves in graphene can be tightly confined to the graphene layer because the real part of the effective index in graphene is much larger than that in free space12. In addition, graphene’s most unique and remarkable property is that its conductivity and permittivity can be dynamically tuned by controlling the chemical potential11.Because of the ability to overcome the diffraction limit of light in microchip-sized devices, SPPs are considered as one of the most promising candidates for nanophotonic components14. Moreover, the large effective index of graphene can efficiently constrain SPP waves around graphene sheet, which presents a variety of possibilities in the design of optical transformation devices, such as modulators18, filter19, cloaks20, and absorbers24. Additionally, transformation plasmonics has opened new avenues towards the realization of on-chip nanophotonic devices13. In particular, because of its easy fabrication, the uneven ground plane biased structure can be used to realize various terahertz devices, such as “electromagnetic black holes”25 and “beam-scanning planar lenses”26. The uneven ground plane below the dielectric spacer support can be used to easily tune the chemical potential of graphene to achieve desired permittivity patterns. However, this scheme is limited by soft-boundary effects27, which imply that the designed uneven ground plane pattern cannot be precisely mapped to the required permittivity pattern.A number of schemes have been reported to create desired permittivity patterns on graphene flakes at the sub-wavelength scaleOne way to realize graded index metasurface is to use a various ground-graphene distance approach, as reported in ref. The graphene SPP platform has a finite propagation length, which will limit the device size on this platform. Within the platform, interferences between devices should be carefully controlled by constructing isolating structures. In this paper, a flexible electromagnetic (EM) fence is proposed to trap and absorb the interferences from other devices. Hence, the entire graphene SPP platform can be divided into different operation regions for different devices. Using the designed graded-index structure, both sides of the EM fence can trap and absorb EM waves. Unlike the EM black hole structure reported in ref. Vg and the SPP wave number β can be easily derived based on the Kubo formula, as given in Method. Notably, salient dots with a constant size are used here to relax the requirement for fabrication tolerances21 and minimize soft-boundary effects27. Clearly, the effective index will mainly depend on the dot density, which will be used to realize the graded-index materials and guided SPP waves on the graphene platform. In this paper, an Al2O3 substrate with a relative permittivity of 3.2 is selected as the dielectric material.In this paper, a Dot-Density-Renderer (DDR) structure is used to build a quasicrystal SPP platform. As shown in Fig. neff should satisfy the expression in Eq.  software COMSOL is employed to analyze the scattering. For this purpose, a fitting method is used to describe index f and incident angle θ. As shown, the reflectivity is better than −10 dB between 6 THz and 12 THz for incident angles from 5 deg to 90 deg. Apparently, the residual reflections are from the unwanted scattering effects of the edge dots of the designed black strip. The transmissivity is found to be better than −30 dB between 8 THz and 12 THz for most of incident angles. Theoretically, the absorption increases with frequencies, this is because that the imaginary part of the index for the dots region becomes larger for higher frequencies, as shown in Fig. The simulated transmissivity and reflectivity of the designed black strip are shown in Fig. Figure Ndot(i) can be calculated according to Eq. 33, a choice of τ = 0.3 ps is an appropriate assumption for our device simulation, according to Eq.  SPP surface wave can subsequently be supported on this sheet, and the guided mode wave number β can be expressed by Eq. . The graded-index structure can be designed by modifying the real part of β.Therefore, when the complex conductivity of a d by Eq.  as Δ → 0μc can be determined by the carrier density n using the following equation:For an isolated graphene sheet, the chemical potential Vg on the gate can modulate the carrier density of graphene through the following equation:Cg = ε/t is the geometric capacitance, ε is the permittivity of the dielectric material, and t is the dielectric layer thickness, which may be the key factor for controlling the electrical characteristics of graphene in practical design. It should be noted that the mode area of monolayer graphene sheet waveguides is extremely small34 (approximately 10−7λ02). In other words, the distance between the graphene sheet and the uneven ground is large enough to ensure that the guided mode of the SPP waves will not be affected.Applying a voltage


The two-dimensional (2D) material graphene is highly promising for Hall sensors due to its potential of having high charge carrier mobility and low carrier concentration at room temperature. Here, we report the scalable batch-fabrication of magnetic Hall sensors on graphene encapsulated in hexagonal boron nitride (h-BN) using commercially available large area CVD grown materials. The all-CVD grown h-BN/graphene/h-BN van der Waals heterostructures were prepared by layer transfer technique and Hall sensors were batch-fabricated with 1D edge metal contacts. The current-related Hall sensitivities up to 97 V/AT are measured at room temperature. The Hall sensors showed robust performance over the wafer scale with stable characteristics over six months in ambient environment. This work opens avenues for further development of growth and fabrication technologies of all-CVD 2D material heterostructures and allows further improvements in Hall sensor performance for practical applications. The areas of application cover many fields, such as automotive, consumer electronics, healthcare and defense industry, where magnetic field sensors are used for position detection, current monitoring and angular sensing. Many different magnetic sensors based on a variety of effects have been realized for different applications2. Hall effect-based sensors constitute a significant part of the industry, with an estimated market share to be over 55% in 20143. They are used for magnetic field detection in the field range from 10−7 T to 102 T in a temperature range from −40 °C up to 150 °C. Today, the most ubiquitous sensors utilize an active region made of Si due to the low fabrication cost, highly developed processing technology, good integration into signal processing circuits and reasonable performance properties 7. In comparison, Hall sensors based on III-V compound semiconductors provide better performances11, but are expensive and more difficult to integrate in circuits.Magnetic field sensors today represent a significant growing market, estimated to reach USD 4.16 billion in value terms by the year 2022nD2 and high carrier mobility μ. The previous reports on graphene Hall sensors demonstrated current-related sensitivities  sandwiched between multilayer CVD h-BN (from Graphene Supermarket26). The 2D material heterostructure was prepared on a Si/SiO2 wafer by large area PMMA-supported wet-transfer technique and Ar/H2 annealing for each layer (service of Graphenea). Next, they were patterned to Hall bar structures by using an Al (20 nm) hard mask for etching with CHF3 and O2 gas. The Al hard mask was removed by wet chemical etching in Shipley Microposit MF-319. The 1D edge contacts were fabricated by means of photo-lithography and electron beam evaporation of metals followed by liftoff in acetone. More specifically, we investigated 1D edge metal contacts of Cr (5 nm)/Au (95 nm) and Ti (20 nm)/Au (60 nm) to 2D graphene channels.The Hall sensors were prepared by using a scalable fabrication process. An optical picture of a chip-size batch-fabricated graphene Hall elements is shown in Fig. 2/Si wafer. The Raman spectrum of monolayer CVD graphene27 on SiO2/Si substrate show the G and 2D peaks at 1597 cm−1 and 2652 cm−1 respectively28 with a small D peak . The detected Hall voltage, measured as a function of applied perpendicular magnetic field B⊥ at bias current I = 90 μA, is shown in Fig. nD2 is the sheet charge carrier concentration11, we can extract the Hall mobilities 2/Vs−1 with carrier concentrations around 8 × 1012 cm−2 at zero gate voltage. The high charge doping is most likely due to the wet transfer process that might have trapped impurities at the h-BN/graphene interfaces. The linearity error is found to be around 0.25% and independent of temperature. The Hall sensors also showed stable performance and good response to magnetic field changes in time  in graphene. Figure I at higher bias, which could be due to heating-related effects19. The more prominent bias-induced decrease of SI with Ti/Au contacts compared to Cr/Au indicates more heating-related effects due to higher contact resistances. Comparing several different devices yields a narrow distribution of Hall sensitivity SI (Fig. I values ranging from 65–78 V/AT at room temperature (Fig. From the Hall measurements the current-related sensitivity  SI Fig.  ranging ure Fig. . From th2 substrate with and without exfoliated h-BN capping (see Supplementary Fig. 2 substrate without any encapsulation show SI in the range of 100–200 V/AT at room temperature. However, we observed degradation in sensitivity with time showing values below 65 V/AT after few weeks of fabrication. The Hall sensors with exfoliated h-BN encapsulation show the current-related sensitivity up to 363 V/AT at room temperature, which is 3 to 4 times higher than the all-CVD heterostructure devices. For the control CVD graphene samples, the Hall mobility and carrier concentrations are found to be around 150 cm2/Vs−1 and −2, respectively, which is lower in comparison to all-CVD heterostructures prepared by wet transfer process. The higher carrier concentrations due to trapped impurities at the h-BN/graphene interfaces are supposed to give rise to lower Hall sensitivity in our all-CVD h-BN/graphene/h-BN Hall sensors.To compare the results of all-CVD h-BN/graphene/h-BN with differently prepared graphene Hall sensors, we have fabricated control samples of CVD graphene on SiOI of 53–78 V/AT at room temperature (Fig. 33. Although the high doping of graphene in our devices can contribute to less degradation of Hall sensitivity over time, the utilized h-BN encapsulation is imperative step forward towards protection of graphene for practical applications in ambient environment.In order to verify the stability of our all-CVD h-BN/graphene/h-BN Hall sensors over time in an ambient environment, we have carried out measurements up to 190 days after fabrication. The Hall voltages measured after 1, 22, 49 and 190 days from device fabrication are shown in Fig. ure Fig. , without7. However, sensitivities were at least one order of magnitude below exfoliated graphene/h-BN structures17, unprotected CVD graphene devices on Si/SiO2 substrate12 and on CVD h-BN substrate19. Also, our control experiment with CVD graphene on SiO2 substrate with/without exfoliated h-BN capping showed 2–4 times higher sensitivity. However, the unprotected graphene showed faster degradation of sensitivity with time. The higher sensitivity of devices with exfoliated h-BN encapsulation can be due to introduction of roughness, ripples and wrinkles in the CVD graphene in all-CVDh-BN/graphene/h-BN heterostructures. The CVD grown multilayer h-BN films are known to be polycrystalline in nature and the c-axis of the crystallites points to random directions30, giving rise to defects and roughness in the heterostructures. The significant doping of graphene in the all-CVD heterostructure devices is another major reason behind the reduced Hall sensitivities and it could be due to contaminations introduced during wet-transfer process of large area graphene and h-BN layers. This could explain also the significantly lower minimum magnetic field resolution of Hall sensors based on all-CVD h-BN/graphene/h-BN heterostructures compared to previous reports on graphene Hall sensors22. These are currently the practical challenges for the development of graphene and 2D materials science and technology.Our observed current-related Hall sensitivities in large area all-CVD h-BN encapsulated graphene devices are comparable to those in magnetic Hall sensors based on Siin-situ growth36 of high-quality h-BN/graphene/h-BN van der Waals heterostructures on large areas would by-pass fabrication-related problems and allow for higher Hall sensitivities with lower graphene doping levels. Additionally, the contacts resistance in 1D geometry can also play an important role on device performance. Our reported reduction in contact resistance by switching from Ti to Cr can be further improved by optimizing the fabrication process, in particular the etching angle and 1D edge contact deposition to eliminate incorporated species at the interfaces24.Such Hall sensors would greatly benefit from improved CVD growth methods of graphene and multi-layer h-BN, and large area layer transfer techniques for fabrication of heterostructures. Furthermore, development of methods for In summary, we demonstrated h-BN/graphene/h-BN van der Waals heterostructure Hall sensors using all-CVD 2D materials available on the market. The batch-fabricated Hall sensors with Cr/Au 1D edge contacts showed reproducible contact properties with low resistances. Hall measurements and time-dependent response at different applied magnetic fields revealed current-related Hall sensitivities in the range of 60–97 V/AT at room temperature. Such encapsulated Hall sensors also showed durable operations in ambient environmental conditions over six months. This study demonstrates proof-of-concept batch fabrication of fully encapsulated all-CVD graphene Hall elements allowing for further development towards practical applications.Supplementary Information


Results demonstrated that graphene had an impact on the microstructure and mechanical properties of the nanocomposites. Based on the mechanical properties and microstructure of the nanocomposites, the strengthening and fracture mechanisms of the graphene-reinforced titanium matrix/nano-hydroxyapatite nanocomposites with different graphene content were analyzed.Biomaterial composites made of titanium and hydroxyapatite (HA) powder are among the most important biomedicalmaterials due to their good mechanical properties and biocompatibility. In this work, graphene-reinforced titanium matrix/nano-hydroxyapatite nanocomposites were prepared by vacuum hot-pressing sintering. The microstructure and mechanical properties of graphene-reinforced titanium matrix/nano-hydroxyapatite nanocomposites with different graphene content were systematically investigated. Microstructures of the nanocomposites were examined by X-ray diffraction (XRD), back scattered electron imaging (BSE), scanning electron microscope (SEM) equipped with energy dispersive spectrometer (EDS), electron probe microanalyzer (EPMA), and transmission electron microscope (TEM). The mechanical properties were determined from microhardness, shear strength, and compressive strength. Results showed that during the high-temperature sintering process, complex chemical reactions occurred, resulting in new phases of nucleation such as Ca As a commonly used artificial bone graft material, titanium alloys have excellent specific strength, corrosion resistance, durability, and no inflammation or immunogenic response, but it is a bioinert material ,2. SimulIn recent years, graphene has been widely used as reinforcement to improve the properties of composites ,22.For eSome researchers have shown that lanthanum could form fine oxides in the matrix, which could play the role of dispersion enhancement and improve the plasticity of the composites by reducing the oxidation of the titanium alloy matrix ,28,29. TThe schematic diagram of the experiment is shown in The phase compositions of raw powders, milled powders, and sintered nanocomposites were measured using X-ray diffraction . The microstructure and composition of the milledpowder and sintered samples were analyzed by scanning electron microscope . Element distribution on the surface of the sintered composite was characterized by an electron probe microanalyzer  equipped with both an energy dispersive spectrometer (EDS) and wave length dispersive spectrometer (WDS).Transmission electron microscopy  was used to further analyze the microstructure and phase composition of the sintered nanocomposites. Microhardness of the sintered nanocomposites was tested using a Vickers microhardness tester . During the testing process, the load was 1 kg, holding time was 15 s, and the testing values of repeated measurements were averaged to improve measurement precision. Compressive and shear strength of the samples were tested with a microcomputer control electron universal testing machine  with a movement rate of 0.5 mm/min. The morphology and characteristics of the fracture surface were analyzed by SEM equipped with EDS.4 arising from mechanical alloying of HA during the milling process [The diffraction peaks of Ti, Ta, Nb, Zr, and HA were obviously detected by the XRD spectra of the composite powders . Due to  process .3(PO4)2, TixPy, Ti3O, α-Ti, β-Ti, Ta, Nb, Ca(OH)2, and HA were detected. The titanium element in the sintered composites was in the form of α-Ti, β-Ti, and Ti3O, and α-Ti was the form of titanium in the raw materials. The β-Ti was formed by the phase transition of titanium during the sintering process, and the existence of the β-phase with lower elastic modulus was beneficial in decreasing the elastic modulus of the composites. The presence of Nb, a β-type stabilizing element, made it possible to retain a large amount of β-Ti after sintering. Titanium oxides might be formed by the reaction between titanium and residual oxygen, but Arifin et al. and He et al. reported that titanium oxide (TiO2) was formed by the reaction between the titanium and water vapor produced from the decomposition of HA [The XRD pattern of the sintered composite is shown in on of HA ,33. FurtThe SEM and EDS results of the sintered composites are shown in 3(PO4)2, were distributed in the matrix. Balbinotti et al. found that phases containing Ca and P in the Ti/HA composite were distributed around the titanium matrix [xPy compound in the XRD results. Distributions of the C and La elements were shown in Elements distribution characterized by EPMA on the surface of the 0.5-GNFs composite is shown in m matrix . In addi3P, according to the selected area diffraction pattern of spot 1 in 3, CaZrO3, and so on by analyzing the selected area diffraction spectrum during the high-temperature sintering process . The posd by XRD  was gened by XRD  was gened by XRD  moved led by XRD  showed t3(PO4)2 detected by XRD (3(PO4)2 are shown in Equations (4) and (5) [3(PO4)2 may be formed by the reaction between HA and titanium oxide, as shown in Equations (6) and (7). Furthermore, thermodynamic analysis (∆G < 0) revealed that the reaction (Equation(7)) could be carried out when the temperature was higher than 720 °C [xPy detected by XRD (xPy was formed by the reaction between titanium and phosphorus which was derived from Equation (10) [In addition, the sintered nanocomposite had decomposed to a certain extent at 1000 °C when compared with the relatively stable pure HA due to the existence of Ti and other metal elements . Ye et ad by XRD  was one  and (5) ,28. Howen 720 °C . Based on 720 °C ,36,37,38d by XRD  was formion (10) . In addiion (10) . It was It can be seen that there were some pores, indicated by the white arrows on the surface of the sample a, which As the enrichment regions of Ti, Ta, Nb, and so on in 3, TixPy, and other phases [3, CaZrO3, and so on, which confirmed the presence of the titanium–phosphorus compound previously analyzed. As described above, HA decomposes to produce CaO after sintering, and the free energy of Equation (12) was −90.811 kcal at 1000 °C, so the reaction could happen. CaZrO3 was also detected in the experiment by He et al. and was considered to be a reaction product of the chemical reaction (Equation (10)) [In addition, the lattice constant, crystal structure, and some other parameters (such as interplanar spacing) of β-Ti and Nb were similar, which made it difficult to distinguish them only by TEM. Therefore, it can be concluded that phases of spot 2  also conr phases . In addir phases . It can on (10)) .(12)CaOBased on the mechanical properties and microstructure of the sintered nanocomposites, the strengthening and fracture mechanisms of the graphene-reinforced titanium matrix/nano-hydroxyapatite nanocomposites with different graphene content were analyzed. It was concluded that the strengthening mechanisms of graphene-reinforced titanium matrix/nano-hydroxyapatite nanocomposites with different graphene content were mainly second-phase strengthening, fine grain strengthening, dislocation strengthening, and dispersion strengthening. The diversity of elements added to the composites would improve the performance of the composites due to second-phase strengthening and dispersion strengthening . Nano-le3(PO4)2, TixPy, and Ti3O. The new phases, which easily dropped off under the action of external force due to weak interfacial bonding with the substrate, could hinder the densification of sintering and increase the brittleness of the nanocomposites.Graphene-reinforced titanium matrix/nHA composites were prepared by mechanical alloying and vacuum hot-pressing sintering. Results showed that during the high temperature sintering process, complex chemical reactions occurred, resulting in some new phases of nucleation such as CaThe shear fractures of the sintered composites were brittle fractures, and there were a lot of detritus and pores on the fracture surface of the composite in which their unique pore structure can give them the characteristic of ultralow density and large surface area for human body with structural and functional integration. There were many tiny particles around the holes, and the reaction and decomposition products of HA were distributed in the alloy matrix. It was believed that the weak interfacial bonding was caused by the difference between reaction products of HA and other materials such as the thermal expansion coefficient.The microhardness, shear strength, and compressive strength of the composites had a decreasing trend with graphene contents of 0.5–1.5 wt %. This was caused by the agglomeration of graphene. Furthermore, the strengthening mechanisms of composites were mainly second-phase strengthening, fine grain strengthening, dislocation strengthening, and dispersion strengthening.


The highest degradation rate was found to be 9.66 × 10−6 min−1, achieved by the sample grown at 700 °C for 5 min, which was 62% higher than the sample just treated at that temperature without graphene growth. This work provides evidence that graphene may be grown across pure Titanium metal and opens new avenues in biomedical devices design, tribological or separation applications.The development of 2D nanomaterial coatings across metal surfaces is a challenge due to the mismatch between the metal microstructure and the nanoscale materials. The naturally occurring thin oxidative layer present across all metal surfaces, may lead to low adherence and connectivity. In this paper, graphene/titania/Titanium hybrid films were for the first time fabricated by a single step chemical vapour deposition process across Titanium foils. The presence of graphene as a dopant was found to enhance the photocatalytic performance of the final products, applied to the degradation of organic molecules and to lead to Schottky-like junction formation at the metal/oxide interface. These Schottky junctions, where vacancies are present across the titania material due to the graphene doping and where Ti The potential of graphene nano-composite materials with enhanced surface properties has been widely demonstrated in the areas of super-capacitors, energy storage devices, sensors, membranes, anti-corrosion coatings and electrode-materials to name a few11.Graphene materials, composed of stacked one-atom thick and two dimensional (2D) sp15. Although the CVD method is one of the most versatile techniques to synthesize graphene grains with controllable surface area and number of layers16, challenges to control the substrate/graphene interface and interactions still remain16. The growth mechanism and the impact of the different growth parameters, including the carbon source, the temperature and growth duration on the graphene quality, are highly related to synergistic interactions with the microstructure and roughness of the growth surface17.Graphene fabrication processes include (i) epitaxial growth, (ii) mechanical cleavage of graphite, (iii) reduction of graphene oxide (GO), and (iv) chemical vapour deposition (CVD)18. These metals can perform both as catalysts and as substrates for graphene growth, therefore leading to either direct deposition of graphene or doping of the metal/metal oxide surface grains19. A major challenge in the CVD growth of graphene coatings is the ageing of the metal exposed for extensive periods at high temperatures. This effect may alter the microstructure and grain distribution across the metal as well as affect the bulk electronic or mechanical properties of the carbon-metal interface18.Substrates for graphene growth by CVD primarily include metal materials, such as copper, nickel, aluminium and stainless steel20, novel growth strategies to deposit and atomically bond graphene to Ti are lacking. A recent study demonstrated that graphene can be directly coated on a Ti based alloy (Nitinol (NiTi 50/50)) through CVD, and the graphene layers were found to enhance the bioactivity of the metal surface in terms of cytoskeleton development and of spontaneous osteogenic differentiation of Mesenchymal stem cells21. Although Nitinol is a material more temperature – stable than pure Ti, the potential release of nickel metal from nitinol alloys is however always an issue for medical application22, compared to pure Ti products. Direct surface functionalization or decoration across pure Ti is also seldom reported, which may be primarily attributed to the low catalytic property and high chemical stability of the Ti metal, which strongly restricts surface functionalisation21. Therefore, sol gel pathways and the implementation of versatile 2D nanoscale coatings, such as graphene, across the surface of a metal as an active and functional layer may bring new functionalities and applications to graphene hybrid semiconductors and related Ti/graphene materials in the photocatalysis and photosynthesis areas27.Although Titanium (Ti) and related alloys, as a promising material for their corrosion resistance, biocompatibility, light-weight and photocatalytic properties, have been extensively studied2/Ti materials. The growth of graphene across the Ti template was favoured though the controlled injection of a hydrogen/acetylene mixture, limiting the titania formation over the seeding steps of the growth at low temperature. The impact of temperature, growth duration and surface roughness of the Ti templates were systematically correlated to the quality and homogeneity of the graphene coatings. Furthermore, photocatalytic performance of the combinatorial hybrid materials was investigated through the degradation of a model organic dye molecule. The presence of the graphene doping across the surface of the materials was found to enhance the bare TiO2 photocatalytic properties by simultaneously facilitating electrons mobility from TiO2 to graphene as well as doping the anatase lattice of the surface titania shifting the band gap of the materials, providing Schottky-like behaviours.Herein, a one-step graphene CVD growth process across pure Ti substrates is demonstrated for the first time to generate hybrid graphene/TiO2) (85/15 at.). All chemicals were analytical grade and used without any further purification.Commercial roll formed Ti foils (purity > 99.6%) with a thickness of 0.1 µm were purchased from Goodfellow Cambridge Ltd. Rhodamine 6 G (R6G) were purchased from Sigma Aldrich. Hexane (99%) were purchased from Sigma Aldrich and its vapor was used as carbon source for graphene growth. The other gases used for the CVD growth were a mixture of argon (Ar)/hydrogen  was used to purge and flush the tube and samples. The pressure was then set with these gases at 4 Torr for 15 min with a flow rate of 400 cubic centimetre per minute (ccm). Under the Ar/H2 gas, a series of samples were treated by adjusting the growth temperature  and duration . The hexane flow rate was maintained at 20 ccm for all tests since previously found to lead to the highest quality of graphene in that similar CVD system18. During the growth duration, the pressure in the tube was maintained at 4 Torr and after growth, the samples were naturally cooled under vacuum which is a quenching process to facilitate the growth of TiO2. After the temperature decreased to less than 100 °C, samples were removed from the tube and stored at room temperature for further use.The growth of graphene across the Ti foils was performed in a low pressure CVD set-up with a tube furnace from Lenton (LTF 14/50/450), following a previously demonstrated procedure by our groupThe surface of Ti foils was polished using the following standard mechanical polishing procedure to investigate the impact of surface roughness on the growth of graphene. The first step involved grinding the sample on 1200 grit SiC paper for 1 min, followed by 3 and 1 micron diamond powder polishing for 5 and 7 min, respectively. The growth conditions across the polished samples were kept the same as these were performed across the plain, unpolished, Ti foils.−1 was performed for 10 s with laser power of 25 mW. X-ray photoelectron spectroscopy (XPS) were performed on a K-Alpha X-ray photoelectron spectrometer . A quantitative elemental composition of Ti, oxygen and carbon for a surface depth of 1–10 nm was carried out. The technique could detect elements with a detection limit of 0.1% of the bulk material. An Al Kα (1486.6 eV) X-ray source was used as the excitation source, and the anode was maintained at 250 W, 10 kV and 27 mA at a chamber pressure of (1 ± 0.1) × 10−9 Pa with an oval beam spot size of 400 μm × 400 μm. In a typical setup, the X-ray source was set at 45 μm from the sample surface while the angle between the analyser and the sample surface is 90°. The analysis was performed on 3 different spots on each sample, under the same experimental conditions. The spectra were acquired from 10 cumulative scans, for a pass energy of 152.1 eV and a dwelling time of 0.5 s. The high-resolution spectra for C 1 s, O 1 s and Ti 2p were acquired at a pass energy of 20 eV and 20 cumulative scans. The spectra were then deconvoluted using the CasaXPS software and the high resolution peaks fitted with a mixed Gaussian-Lorentzian function (GL30), as previously reported28, while the asymmetric nature of the sp2 C in the C 1 s spectra was fitted with a Gaussian-Lorentzian product formula modified by an asymmetric form AGL(80) to account for its asymmetric tail29. The asymmetric shape of Ti (0) in the Ti 2p spectra is fitted with a LA30. X-ray diffraction (XRD) patterns were collected on a X’Pert Powder instrument  using a Cu Kα radiation source (λ = 1.54181 Å) operated at 40 kV with a 30 mA current. XRD data were recorded over a range of 6–70°. The ultraviolet visible (UV-Visible) spectra were obtained from USB-2000 ultraviolet visible spectrometer . The R6G absorbance peak at the wavelength of 526 nm under UV-Vis was used to determine the concentration over time. The UV-Visible Diffuse Reflectance Spectroscopy (DRS) was conducted on a LAMBDA 1050 UV-Vis spectrophotometer  equipped with an integrating sphere. Prior to the analysis, the reflectance port was removed to take off the contribution of the specular reflection from the total reflectance, only the diffuse reflectance was consequently collected. The reflection was calibrated on a Spectralon standard, and the calculated %R is referred to this standard. The data was acquired with a UV WinLab software at a fixed scan speed of 485.33 nm/s, and the Kubelka Munk function subsequently determined.Scanning electron micrographs (SEM) were captured on a JSM 7800 F FEG-Scanning Electron Microscope . The tests were performed on un-coated samples as provided with a 5 keV accelerating voltage at a distance of 10 mm without prior coating. Raman spectra were obtained on an inVia Raman microscope  at the laser wavelength of 514 nm. An extended scan ranging from 100 to 4000 cm2 was placed on the bottom of the cuvette, which was exposed to UV light generated from an OmniCure S2000 UV lamp  to induce the degradation of R6G. The intensity of UV light is 250 mJ/cm2 with the wavelength, set with a grated plate filter, between 320 and 480 nm and the sample to lamp distance is 10 cm. The degradation of R6G was determined through measuring the peak absorbance of R6G at the wavelength of 500 nm at different time intervals.As a model molecule, R6G was selected to test the photocatalytic performance of Ti foils with or without graphene growth. A stock solution of R6G was prepared at a concentration of 7.5 mg/L. For each test, 2 mL of R6G solution was transferred to a quartz cuvette and the sample with an area of around 25 mm2/Ti hybrid materials were investigated in terms of their morphology, structure, and photocatalytic properties, to not only understand the growth mechanism, but also evaluate synergistic properties and the role of graphene doping on the band gap of the semiconductor. A schematic of the growth process is shown in Figure Graphene/TiO31. The hexane gas flow during graphene growth at the experimental temperature was also found to affect the surface morphology. Without any gas flow, there was noticeable large crystal-like structure formed in the heat-treated samples  of graphene, which is related to the number of graphene layers36. The Raman results from the time series grown at 900 °C indicate that graphene can be clearly detected from 5 min of growth duration. For shorter growth durations, it is likely that the carbon atoms surface segregation leads a predominant role in the graphene growth mechanisms. As previously shown for catalytic surfaces similar to Ti/TiO2, such as Ni surfaces, carbon atoms shall first diffuse from the surface into the first few nanometers of the bulk of the growth support during the annealing stage and predominantly at high temperature. After the diffusion step, the carbon atoms shall precipitate and recombine onto the activated Ti catalyst during the cooling period37. Further improvements in the growth conditions to limit H2 uptake may include a sharper control of the moment when H2 gas is introduced, particularly upon reaching a set reactor temperature higher than 700 °C.The graphene component of hybrid graphene/TiOopy Fig. . The blaopy Fig. . For the39. The 2D-band is clearly shown in the sample grown at 900 °C and the broad 2D-band indicates a thick graphene layer21. Similar trends and graphene quality were obtained for the polished Ti foils  intensities can be used to assess the presence of defects across the graphene42. As can be seen from Table D/IG ratio was achieved for the sample grown at 700 °C for 5 min, correlated with the overall carbon content, which was also the highest compared for this sample across the series, as indicated by the XPS results  3/2p and 1/2p can be clearly shown at 458.6 eV (FWHM = 1.22 eV) and 464.4 eV (FWHM = 2.03 eV), respectively, in good accordance with the literature30. Two more oxidation states are however present, as Ti3+ and Ti(0), classified in Table 3+ species can be observed in Figure 3+ 2p 3/2, as shown in Table 55, although these species are highly unstable and their generation and control is particularly challenging. The presence of these species supports the design of the Schottky-like diode by lower energy barrier to activation. This behaviour was attributed to the presence of a multi-oxide interfacial layer in the top layer, available to UV excitation, of the hybrid material, as previously demonstrated for Ti ant TiOx alloys56.In order to gain more insights into the potential generation of O vacancies, and consequently of Ti suboxides that are known to be beneficial in photocatalytic phenomena2 layer was further investigated by XRD . The potential of graphene nanomaterials and as dopant was recently investigated in the form of quantum dots58 but also low dimensional semi-conductors60. The absorption spectra show that the absorbance peaks of R6G for all samples decreased with respect to exposure duration (Figure 61:C0 and C are the initial concentration of R6G and the concentration of R6G over time, k is the kinetic constant, a is a constant, and t is the reaction time. The fitting results for the UV-visible spectra are shown in Table 2/Ti hybrid materials are 62%, 110% and 205% higher than those achieved by the control heat-treated samples at 700 °C, 900 °C and 1000 °C, respectively. This trend therefore indicate that graphene has a positive influence on the photodegradation properties of the Shottky-like surfaces.The photocatalytic performance of the graphene/TiOn Figure . The kin16, combining graphene with TiO2 offers synergistic effects whereby electrons generated form the dye molecules may flow more freely from the TiO2 anatase to the graphene domains, thus generating Schottky-like barriers at the interface between TiO2 and conductive layers such as metal or graphene and due to the presence of the vacancies in the form of Ti3+ ions, leading to an enhanced charge separation. A Schottky barrier interface electric field between the Ti-metal and the Ti-oxides shall confine vacancies to the surface of the semiconductor as also observed across metal doped TiO2 catalysts62. Such interfaces may generate photoexcited electron-hole pairs and an electrostatic depletion layer at the junction between the metal-like graphene and the titania semiconductor, which causes the hybrid material to act as an electrical rectifier63. This process may thereby delay the charge recombination upon catalytic reactions65. The incorporation of graphene therefore offers outstanding doping capabilities for TiO2 photocatalytic properties enhancement. In addition, Ti left in the hybrid materials can also perform as a conductor to conduct electrons generated by TiO2 during the photocatalytic process.Due to the high electron mobility of 2-dimensional graphene sheets68. However, these results are comparable with similar TiO2 based film system with the complete dye decomposition time between 3 and 4 h70. The presence of graphene doping the TiO2 matrix as well as vacancies as shown by the Ti3+ content in this hybrid materials facilitated the photocatalysis, despite of low anatase TiO2 content that is regarded to be more efficient than pure rutile while being applicable onto large surface areas of materials in a facile manner.The photocatalytic efficiency reported in this study is not as good as that of slurry system due to the lower exposed surface area that is available for photocatalytic reactions and low porosity of the supported catalyst layer2/Ti hybrid samples, the spectra in Fig. 2, respectively, during the transition from anatase to rutile shown on the XRD results, which is in good agreement with other studies72. One would assume that plain TiO2 would absorb only in the UV-region as it has no activity under visible light due to the high band-gap, 3.2 for anatase and 3 for rutile, respectively. The presence of clear absorption features in the visible range is therefore inconsistent with the literature74. However, considering that the photocatalytic tests were conducted under a UV light with wavelength up to 500 nm, the outcomes of the analysis are in good accordance with these spectra demonstrating that the plain foils absorb radiation in that range. The sample where graphene was grown at 700 °C is offers the strongest absorption which corresponded to the best catalytic performance.The impact of the doping was therefore further evaluated by DRS to evaluate the change in the band gap of the material. Figure 2/Ti hybrid film was demonstrated on a pure Ti foils through CVD method. Issues related to H2 absorption by the Ti matrix were partly alleviated by controlling the temperature ramping rate during graphene growth. By manipulating the growth conditions, the content of graphene and TiO2 can be controlled. The introduction of graphene to TiO2 and the formation of a high yield of vacancies in the form of Ti3+ ions within the matrix has strongly reinforced the photocatalytic properties of the hybrid material. Worth to mention, the use of film based photocatalyst is advantageous to the nanoparticle based system, as films can be easily applied and removed from the dye solution, minimising the concern of any surplus nanomaterials left in solution after photocatalytic treatment. The materials fabricated in this paper has provided practical application of graphene grown on pure Ti metals, which potentially expand the CVD graphene to a new metal member and explore the related applications.The fabrication of graphene/TiOSupplementary Information