Epitaxial graphene has proven itself to be the best candidate for quantum electrical resistance standards due to its wide quantum Hall plateaus with exceptionally high breakdown currents. However one key underlying mechanism, a magnetic field dependent charge transfer process, is yet to be fully understood. Here we report measurements of the quantum Hall effect in epitaxial graphene showing the widest quantum Hall plateau observed to date extending over 50 T, attributed to an almost linear increase in carrier density with magnetic field. This behaviour is strong evidence for field dependent charge transfer from charge reservoirs with exceptionally high densities of states in close proximity to the graphene. Using a realistic framework of broadened Landau levels we model the densities of donor states and predict the field dependence of charge transfer in excellent agreement with experimental results, thus providing a guide towards engineering epitaxial graphene for applications such as quantum metrology. In these models the linear magnetic field dependence of n is predicted to saturate at high magnetic fields. However, further work has suggested that the carrier density continues to increase to much higher magnetic fieldsδ-function Landau levels in graphene and constant densities of states within the charge reservoirs. Experimental data in which the full extent of the magnetic field dependence can be measured is therefore crucial in obtaining a complete understanding of this process which underpins the benefits of graphene in quantum metrology. Using a combination of pulsed high magnetic fields up to 57 T and very low carrier density samples we present here the first detailed study of the full magnetic field dependence of charge transfer in epitaxial graphene through a ν = 2 plateau which remarkably extends from B = 3 T to over 50 T. We show that by using the occupancy dependence of the quantum Hall breakdown current we can demonstrate that the carrier density increases almost linearly throughout the ν = 2 plateau. This method is validated using a fit to the temperature dependence of the magneto-conductivity which provides a direct measure of the position of the Fermi energy with respect to the nearest (N = 0) Landau level and hence provides an independent measure of the carrier density. We also demonstrate that by extending the existing models1112In epitaxial graphene grown on the Si-terminated face of SiC, charge transfer from the underlying substrate to the graphene leads to a strongly magnetic field dependent carrier density and an exceptionally wide quantised 10 cm−2, as determined from the low field Hall coefficient ν = 2 quantum Hall state is observed to begin well below 1 T and the longitudinal resistivity, Rxx is undetectably small from 0.7 T up to 14 T, as shown in jc = Ic/W a sudden onset of longitudinal resistance indicates the breakdown of the quantum Hall effect. The value of jc is defined as the current density above which Vxx > Vc, where the critical voltage Vc is determined by the noise limit of the measurement. It can be extracted from the I − Vxx characteristics in the quantum Hall state, as shown in Vc = 0.5 μV. By measuring the breakdown current as a function of magnetic field throughout the ν = 2 plateau we are able to extract a wealth of data, in particular about the magnetic field dependent carrier density.We begin with the sample having its lowest density, where the zero field carrier density is 1.5 × 10jc centred exactly at an integer filling factor (e.g. ν = 2) falling linearly to zero at the edges of the plateau (approximately ν = 2 ± δν)2324In traditional semiconductor quantum Hall systems, where the carrier density is constant, the breakdown current has been shown to have a triangular dependence on both filling factor and magnetic field, where a sharp peak is observed in δν is typically in the range 0.2 < δν < 0.32324N = 0 or N = 1 Landau level become energetically easier to access, dissipation occurs at progressively lower measurement currents until jc decreases to zero. In order to explain the extremely wide plateaus observed in graphene both here and previously312ν = 2 state changes only slowly with magnetic field thus creating an asymmetric and very broad jc profile along the plateau (where  plateau .σxx) of this sample. The longitudinal conductivities as functions of inverse temperature from 1.4 to 200 K for a series of magnetic fields are shown in σxx is linear with 1/T. In contrast, the non-linearity of the log σxx − 1/T plot at low temperatures shows the importance of variable range hopping (VRH) which causes significant deviations from the conventional Arrhenius activation behaviour. Excellent fits to the total longitudinal conductivity can be achieved however by combining thermal activation between extended states in adjacent Landau levels and VRH. The total conductivity is calculated as,In order to produce a quantitative measurement of the magnetic field dependent carrier density, we have studied the temperature dependence of the magneto-conductivity  which comes mainly from σTA, dominant at high temperatures. No significant difference in EF is found by using either the Mottwhere T0 is a fitting parameter. In terms of σTA, obtaining a simple analytic formula can be difficult, especially when characteristics of Landau level broadening are taken into account2829where the characteristic temperature gv = gs = 2 are the valley and spin degeneracies, EN is the energy of the Nth Landau level, and the standard deviation s represents the Gaussian broadening of a Landau level. In our case the thermal energy is small compared with the Landau level separation and the only contributions are from the extended states within the mobility edges (EN ± Eμ) at the centres of the N = 0 and N = 1 Landau levels. Since the Fermi level EF is also well within the region of localised states, from the Kubo formulawhere k = γβ(EN − EF) − (βs/2)2, where N = 0 and N = 1 Landau levels have fixed Gaussian broadening, i.e., s and Eμ are constants independent of magnetic field and temperature. The value of s is chosen to be ~12 meV in order to match the energy scale of the low-field Gaussian random disorder potentials in these systemsEμ is determined to be 7.8 meV  by fitting the data for B = 1.2 T, where EF is known to be exactly at the mid-point between the two Landau levels since at elevated temperatures the minimum in conductivity and Rxx is observed at this field. We also note that the dimensionless parameter EF is a slowly decreasing quantity which approaches the mobility edge Eμ gradually. This suggests that the magnetic field dependent charge transfer rate is slightly lower than the increase in the total Landau level DOS. EF as a function of magnetic field is given by G(E) is given by Employing the above formulation, the temperature-dependent magneto-conductivities are very well fitted across a broad temperature range from 1.4 up to 200 K, as shown in In order to provide a quantitative relation to describe the charge transfer which can be compared to the magnetic field dependence of the breakdown current, we introduce the following phenomenological expression for the field dependent carrier density,p, B0 and Bc are fitting parameters. This expression is chosen as, for large values of p, it rapidly swaps over from a constant carrier density n0 at low field, where ν = 2 would occur if there were no charge transfer at this low field carrier density, to a slightly sublinear increase with field Bc is the magnetic field corresponding to ν = 2 including the additional charge transfer. The carrier density and occupancy predicted by this equation for this sample are shown in p = 13. Using the analytical form for n given by p = 5, 10, 15 and 20, which demonstrate the appearance of a strong asymmetry as a function of field and a very considerable broadening of the breakdown current field dependence. jc to the data deduced from p = 13 which gives an excellent description of the data. We use p = 13 in all subsequent fitting. The magnetic field at which jc goes to zero (14 T) corresponds to an occupancy as deduced from the temperature dependence of σxx of 1.74  current of 3 μA in a 57 T long pulse magnet. At both positive and negative magnetic fields the magnetotransport is dominated by the extended ν = 2 plateau. Using the low field Hall coefficient we extract an initial carrier density of 1.5 × 1011 cm−2. Also at low magnetic fields we observe a clear weak localisation peak in the resistivityν = 6 and ν = 10  from ref. and,ϵ1(2), γ1(2), and d1(2) are the absolute permittivity, the assumed constant DOS of surface donor states, and the distance between graphene and the donor states for the SiC/graphene (graphene/PMMA) interface, respectively  =  (ϵ0 = vacuum permittivity), d1 = d2 = 0.3 nm, we plot the relation between γ1 and γ2 in γ1) decreases from 30% to 0% of the total DOS in the two charge reservoirs, γ2 will rapidly increase to infinity. It is not physical to have such a large DOS in the PMMA interface (γ2), so it is unlikely that γ1 + γ2 remains roughly constant over a large range, γ1 ~ 0.8 − 1.7 × 1014 cm−2 eV−1, γ2 ~ 0.3 − 1.3 × 1014 cm−2 eV−1, with γ1 + γ2 ~ 2 × 1014 cm−2 eV−1. Our estimation suggests that the surface states in SiC play a more significant role in the charge transfer process. The surface states responsible are closely related to the large number of atomic defects (vacancies and adatoms) predominantly located in the first underlying atomic layer, which consists of a non-conductive surface reconstruction of carbon atoms covalently bonded to the SiC which is formed during the growth process3738394041424344EF3740434415 cm−2 for the reconstructed layer leads to DOS in the range of 0.8 − 1.5 × 1014 cm−2 eV−1. Despite the variations in experimental conditions, the reported values are consistent with our estimate for γ1. For the graphene/PMMA interface, we attribute the less significant but still relatively high γ2 to the increased number of charge traps introduced by the intensive UV treatment2021Bc (the magnetic field corresponding to the occupancy ν = 2 where the minimum Rxx and the maximum jc occur) as a function of the zero field carrier density. Based on Another important consequence from our model is that we can deduce vF is the Fermi velocity. The left hand side and the right hand side of the equation represent the charge equilibrium at B = 0 and B = Bc, respectively, with EF being B = 0, and at the mid-point between the N = 0 and N = 1 Landau levels Bc. The solution to B0 (red lines) in Bc = B0, which represents the case of a constant carrier density without the field dependent charge transfer. Combining ν = 2 plateau occurs, using B0 as shown in the where nmax that can be transferred into graphene, corresponding to the limit B → ∞ and EF → 0 when other high magnetic field effects  are neglected. From Finally, we calculate the maximum carrier density or,nmax and n0 are both in units of 1011 cm−2. This relation is shown in nmax represents the upper limit to the possible validity of B → ∞. In practice, the highest carrier densities reached in our measurements, nmax due to the finite magnetic field strength we could apply. Therefore, in this regime, n(B).where Bc, nmax) of the breakdown of the quantum Hall effect and the field dependent charge transfer, using just the zero field carrier density (n0 or B0). These results can thus provide realistic references for the optimum operating conditions for a quantum Hall resistance standard using epitaxial graphene. We emphasise that our model can ultimately provide a complete n(B) dependence, if the exact details of the Landau level broadening are known. Such dependences are simulated in n0 listed, by numerically solving the equation set,Hence, with our model using a realistic framework of broadened Landau levels, we have been able to relate and accurately predict some of the most important characteristics (G(E) is the DOS of Landau levels. In this simulation, we continue to assume Gaussian Landau level broadening, such that G(E) is given by s = 12 meV of the Gaussian broadening for the same sample (n0 = 1.5 × 1010 cm−2) on which the temperature dependent measurements were made. For the samples with n0 = 1.5 and 8.7 × 1011 cm−2 in which s is unknown, our model suggests values of about 25 and 30 meV, respectively, both well within the normal range among similar samplesn(B) dependences given by our charge transfer model (dashed lines), and the those given by p = 13 (solid lines). At high enough magnetic fields, the inset to nmax as given by where v = 2 plateau can be observed in this system for low enough starting densities, n0, and demonstrate the very large extent of the magnetic field dependent carrier density in epitaxial graphene. We have shown that the quantum Hall effect breakdown current can be used to accurately measure this increase in carrier density with field, which is found to be over an order of magnitude in some cases. Using the models that we propose, we have been able to accurately describe and predict some of the most important features of the field dependent charge transfer process and its effects on the quantum Hall breakdown. Our models and results are widely applicable towards a broader and deeper understanding of the high magnetic field transport properties of graphene and are crucial for engineering epitaxial graphene devices for applications such as quantum Hall metrology.In summary, we have studied the breakdown and the temperature dependence of the quantum Hall effect using high magnetic fields and very low density epitaxial graphene grown on SiC. Our new measurements show that the full width of the W = 20 μm and length L = 20 μm. Electrical contacting was made using large area electron beam evaporated Ti-Au. A non-volatile dual-polymer gating technique using PMMA/MMA and ZEP520A was applied to tune the carrier density in the epitaxial graphene using a combination of UV illuminationT = 2 K unless otherwise stated.The devices studied were prepared from epitaxially grown graphene on the Si-terminated face of SiC. Each device was patterned using electron beam lithography and oxygen plasma etching into an eight-leg Hall bar geometry with a width How to cite this article: Alexander-Webber, J. A. et al. Giant quantum Hall plateaus generated by charge transfer in epitaxial graphene. Sci. Rep.6, 30296; doi: 10.1038/srep30296 (2016).


Based on the experimental data, the most probable structure of the complexes is suggested. In the present study, we have examined cytotoxic activity of (9′-fluorene)-spiro-5-hydantoin (L1) and (9′-fluorene)-spiro-5-(2-thiohydantoin) (L2) and their Pt(II) complexes on the retinoblastoma cell line WERI-Rb-1.This paper presents a method for synthesis and cytotoxicity of new platinum(II) complexes of (9′-fluorene)-spiro-5-hydantoin (L1) and (9′-fluorene)-spiro-5-(2-thiohydantoin) (L2). The new obtained complexes were studied by elemental analysis: ultraviolet–visible, attenuated total reflection Fourier transform infrared (ATR-FTIR), and  The profound interest on different hydantoin derivatives stems from the well-established medical applications of some hydantoin derivatives as antiepileptic drugs. Recently1H- and 13C-nuclear magnetic resonance (NMR) spectroscopy) .. Furthermore, considering literature review and biological activity of fluorenylspirohydantoins, the study of the possible effects that these compounds may have on the viability of human tumour cells (WERI-Rb-1) is particularly interesting.The aim of this study is to present a method for synthesis of new platinum(II) complexes of (9′-fluorene)-spiro-5-hydantoin (L1) and (9′-fluorene)-spiro-5-(2-thiohydantoin) (L2), its isolation in pure form and its characterization using various spectral methods (ultraviolet–visible (UV–vis), attenuated total reflection Fourier transform infrared (ATR-FTIR), and 1H- and 13C-NMR for Pt(II) compounds and additionally Raman spectroscopy for free ligands .In the present work, we studied the complexation properties of (9′-fluorene)-spiro-5-hydantoins and its 2-thio derivative see . The rea4)2[PtCl4] – Sigma-Aldrich) and solvents used for synthesis of the complexes were with a p.a. qualification. The two organic ligands L1 and L2 were synthesized according to methods described in Marinov et al. [A metal salt  of (9′-fluorene)-spiro-5-hydantoin (L1) in 10 cm4)2[PtCl4] in 10 cm3 H2O;0.0002 mol (0.0746 g)  of (9′-fluorene)-spiro-5-(2-thiohydantoin) (L2) in 10 cm4)2[PtCl4] in 10 cm3 H2O;0.0002 mol (0.0746 g)  and the stirred crystals of L1, L2, Pt(II)L1 and Pt(II)L2 were pressed by an anvil to the reflection element. The Raman spectra of L1 and L2  were measured on a RAM II (Bruker Optics) with a focused laser beam of 500 mW power of Nd:YAG laser (1064 nm) from 4000 to 100 cm−1 at a resolution of 2 cm−1 with 25 scans. The NMR spectra were taken on a Bruker Avance II+ 600MHz NMR spectrometer operating at 600.130 and 150.903 MHz for 1H and 13C, respectively, using the standard Bruker software. Chemical shifts were referenced to tetramethylsilane (TMS). Measurements were carried out at ambient temperature.Electronic spectra of absorption were registered on a Lambda 9 Perkin-Elmer UV/vis/NIR(Near infrared) Spectrophotometer from 200 nm to 1000 nm. ATR-FTIR spectra were registered on a VERTEX 70 spectrometer (Bruker Optics) from 4000 to 600 cmThe physicochemical and spectral characterizations of the obtained compounds were as follows:2.The retinoblastoma cell line (WERI-Rb1) was purchased from ATCC-HTB-169. It was cultured in RPMI 1640, containing 10% FCS, streptomycin and penicillin at 37 ºC in a humidified atmosphere containing 5% CO4 cells/well. After a cultivation period of 24 h, the compounds were added at a concentration of 100 μM and incubated for 24, 48 and 72 h, respectively. WST-1 was added to the cells at these time points and incubated for 4h. WST-1 is a colorimetric assay for the non-radioactive quantification of cell proliferation, viability and cytotoxicity. The highest achievable concentration at which we were able to test it was only 100 μM.The cytotoxic effect of L1, L2, Pt(II)L1 and Pt(II)L2 was assessed on a retinoblastoma cell line using WST-1 assay . All compounds were first dissolved in DMSO and then diluted in the culture medium. The concentration of DMSO in the wells did not exceed 1%. Cells were seeded in triplicates in 96-well plates at a density of 6.5 × 10Absorbance was measured on ELISA SUNRISE Reader. The wavelength for measuring the absorbance of the formazan product is 450 nm and the reference filter was set at 620 nm in accordance with the WST-1 manual. Cells grown in culture medium alone and in appropriate concentrations of DMSO were used as controls. The percentage of viable cells was calculated as a ratio of the OD value of the sample to the OD value of the control. The data are presented as mean ± standard error of the mean.4)2[PtCl4] at molar ratio M:L:OH− = 2:4:1 for PtL1 and 1:2:3 for PtL2. Neutral complexes were synthesized and isolated as brown and yellow precipitates, respectively. All complexes were investigated by elemental analysis: UV–vis, ATR-FTIR, and 1H- and 13C-NMR spectroscopy.Complexation with Pt(II) was conducted under alkaline conditions using a metal salt namely L1 and Pt(II)L2 complexes were observed at 227 nm, 276 nm and 222 nm, 230 nm, 273 nm, respectively. Solid ATR-FTIR spectra of the free ligand and its complexes were recorded in order to clarify the structure of the formed metal complexes and to determine the coordination modes of the ligand. All vibrational frequencies observed in the ATR-FTIR spectra of the complexes and the free ligands were given in “Materials and Methods” section. The bands at 3360 and 3189 cm−1 of the free ligand L1 may refer to the stretching vibrations of two N–H groups of the hydantoin ring. In the ATR spectrum of the PtL1 complex, the same bands were observed at 3232 and 3136 сm−1 which shifts to lower frequency by 128 and 53 сm−1 as compared to the corresponding free ligand bands. The bands of free ligand L1 at 1775 and 1715 сm−1 can be attributed to the stretching vibration of C4 = O and C2 = O groups of the hydantoin ring. The same bands in ATR spectrum of PtL1 complex were at 1774 and 1705 сm−1, i.e. a slight change is observed. This fact shows that the two carbonyl groups of the ligand hydantoin ring did not participate in the coordination.The elemental analyses data showed metal-to-ligand ratio 1:6 for PtL1 and 1:2 for PtL2. The UV–vis spectra were registered in DMSO. Maxima in the UV–vis spectra of the free ligands L1 and L2 were observed at λ−1 (shoulder) and 3153 cm−1 were observed which may refer to the stretching vibrations of N–H groups of the hydantoin ring. These bands disappeared in the spectrum of the PtL2 complex; there was a very wide, low-intensity band between 3500 and 2500 cm−1. On the whole, all bands in the ATR spectrum of PtL2 complex were wider than those in the free ligand. In the spectrum of L2, there were bands at 1751 and 1729 сm−1 and 1530 сm−1; the first two bands can be attributed to the split stretching vibration of C4 = O, and the second to stretching vibration of C2 = S group. For the PtL2 complex, the first two bands were shifted by 21 and 23 сm−1 to 1772 and 1752 сm−1, respectively. On the other hand, the thiocarbonyl stretching was shifted by only 3–1527 сm−1. This fact shows that the thiocarbonyl group did not participate in the coordination with metal ion. Based on the experimental data, the most probable structure for the PtL1 complex was suggested with two deprotonated NH groups of ligand L1 and four uncharged organic ligands molecules. For the PtL2 complex square planar geometry was suggested with two ligand molecules coordinated in a bidentate fashion.In the ATR spectrum of the free ligand L2, the bands at 3244 cmOur results revealed that L1, Pt(II)L1 and Pt(II)L2 had a cytotoxic effect on the human retinoblastoma cell line WERI-Rb1  and  3. in vitro experiments are required to study the possible applications of hydantoin derivatives and their mechanism of action.Based on these results further 1H- and 13C-NMR spectroscopy, confirmed their structures. The preliminary results of our study showed that L1, Pt(II)L1 and Pt(II)L2 exhibit anticancer activity in vitro.The synthesis of two Pt(II) complexes with (9′-fluorene)-spiro-5-hydantoin (L1) and (9′-fluorene)-spiro-5-(2-thiohydantoin) (L2) were described and the various spectral data, UV–vis, ATR-FTIR, and


A proportion of Alzheimer's disease cases displays inclusions of the RNA-binding protein, TDP-43. Considering the pathogenic role of tau mis-splicing, we compared tau isoform expression between Alzheimer's disease cases with or without TDP-43 inclusions. The average ratio of tau isoforms containing or lacking exon 10 (4R/3R ratio) or the total level of tau mRNA was not significantly different between cases with or without TDP-43 pathology in any of the brain regions examined. Although TDP-43 functions may be affected, TDP-43 does not critically regulate expression or splicing of tau in Alzheimer's disease suggesting that TDP-43 contributes to Alzheimer's disease through mechanisms independent of tau. In addition, 20%–30% of AD cases display cytoplasmic inclusions of transactive response DNA-binding protein of 43 kDa (TDP-43). The presence of TDP-43 inclusions in the hippocampus and entorhinal cortex correlates with neuronal loss in late onset dementia and, furthermore, a strong correlation exists between the presence of TDP-43 pathology and specific clinical features of AD, especially impaired cognition and amnesia . TDP-43 2t test.Frozen brain tissue was obtained from individuals free from neurological disease and individuals with AD and controls. Cases were assessed for TDP-43 pathology by immunocytochemistry. RNA was extracted using the Qiagen RNeasy/lipid kit (Qiagen). Tau E10 splicing was analyzed by end-point reverse transcription polymerase chain reaction using a forward primer 5′-labelled with an infrared fluorescent dye, yielding a signal independent of the length of the individual products. Gels were imaged, and polymerase chain reaction products were quantified using an Odyssey Infrared Imaging System (LI-COR Biosciences) . Statist3We analyzed tau splicing in a cohort of 14 AD cases with (ADTDP+) and 15 AD cases without (ADTDP−) TDP-43 inclusions and 15 age-matched nondemented healthy controls . The amy4TARDBP gene, encoding TDP-43, have been linked to familial forms of amyotrophic lateral sclerosis (Mutations in the clerosis , thus arclerosis . Howeverclerosis . Whetherclerosis . ConsideThe authors declare that they have no conflict of interest.


It is found that the microwave reaction time has a significant effect on the microstructure of graphene/MnO2 composites, and consequently, the electrochemical properties of the supercapacitors based on graphene/MnO2 composites are strongly microstructure dependent. An appropriately longer microwave reaction time, namely, 15 min, facilitates a very dense and homogeneous microstructure of the graphene/MnO2 composites, and thus, excellent electrochemical performance is achieved in the supercapacitor device, including a high specific capacitance of 296 F/g and a high capacitance retention of 93% after 3,000 times of charging/discharging cycles.Graphene has attracted much attention since it was firstly stripped from graphite by two physicists in 2004, and the supercapacitor based on graphene has obtained wide attention and much investment as well. For practical applications of graphene-based supercapacitors, however, there are still many challenges to solve, for instance, to simplify the technological process, to lower the fabrication cost, and to improve the electrochemical performance. In this work, graphene/MnO81.05.ue; 78.67.Sc; 88.80.fh However, metal oxides usually have a high electrical resistance, thereby leading to a low power density, and the high cost of RuO2 also limits its wide applications. For the second group, they are usually used in double-layer capacitors because of their electrochemical stability and high accessible surface area. For the third group, polyaniline and polypyrrole [2 composite is such an alternative material.In recent years, with the deterioration of the environment and the scarcity of natural resources, more and more researchers have turned their attention to the field of energy. Supercapacitor, serving as a novel energy storage device, is one of the mostly focused topics. The materials for supercapacitors that have been intensively studied so far can be divided into three groups: transition metal oxides, carbon materials, and conductive polymers . In the ypyrrole  for instypyrrole . Due to 2 can form a pseudocapacitance, thereby increasing the capacitance of the capacitor. To prepare graphene more simply and more safely, and finally to make the mass production come true, various methods have been developed [4 in a pH-neutral solution is expressed as follows [Owing to its high specific surface area, remarkable chemical stability, and superior electron mobility ,7,8, graeveloped , includieveloped , chemicaeveloped ,15, preseveloped , epitaxieveloped , and cheeveloped -20. In teveloped . The red follows :2 composite solution directly for a period of time to coat the active material on a current collector. Moreover, the process-structure-property relationships were systematically investigated for the graphene/MnO2 composites. Interestingly, it was found that the microwave reaction time has a significant effect on the microstructure of graphene/MnO2 composites and the electrochemical properties of the supercapacitors based on graphene/MnO2 composites are therefore strongly microstructure dependent.It is noteworthy that a simple method was used here for the supercapacitor packaging, i.e., dipping Ni-foam into a graphene/MnOGraphite oxide (GO) was synthesized firstly from natural graphite according to a modified Hummers method . A GO so2 composites were prepared by a redox reaction between graphene and potassium permanganate under microwave irradiation [4 powder (0.95 g) was added into the graphene suspension and stirred for about 10 min. Subsequently, the resulted suspension was heated using a household microwave oven  for several minutes, and then cooled to room temperature naturally. Finally, the black deposit was filtered and washed several times with distilled water and alcohol, and then dried at 80°C for 6 h in a vacuum oven. Different microwave reaction times, namely, 5, 10, and 15 min, were selectively used to study its effect on the microstructure and the electrochemical properties of the graphene/MnO2 composites. Due to the lab condition restriction and, more importantly, in order to avoid the overflowing of the reaction solution, we did not extend the microwave reaction time further.Graphene/MnOadiation . In the 2 composites were measured by X-ray diffraction  using Cu Kα radiation (λ = 0.154056 nm). The microstructure was characterized by scanning electron microscopy  and transmission electron microscopy . Note that the samples were dispersed in alcohol and dropped on a holey copper grid for TEM observations. The electrochemical measurements  were measured by using an electrochemical station  with a two-electrode system, which consists of two identical working electrodes in 6 M KOH alkaline electrolytes.The crystallographic structures of the graphene/MnO2 composite, carbon black, and polytetrafluoroethylene (PTFE), were mixed in a mass ratio of 75:20:5 and dispersed in ethanol [The fabrication of working electrodes was carried out in the following way. Briefly, the materials, including graphene/MnO ethanol . Then th2 composites with different microwave reaction times . It can be seen that the as-prepared graphene shows a weak and broad diffraction peak at 2θ = 23°, corresponding to the diffraction of the (002) plane. For the three graphene/MnO2 composite samples, their XRD patterns are almost the same, exhibiting a weak and broad peak of the (002) plane centered at 2θ = 23° (indexed for graphene) and two weak peaks indexed for α-MnO2 at 2θ = 37° (211) and 2θ = 66° (002) [θ ~ 43°, corresponding to the (100) crystal plane of graphene, cannot be observed after the deposition of MnO2, indicating that the surfaces of graphene were fully covered by nanoscale MnO2 and thus a lower degree of graphitization was induced, which was similarly reported by Yan et al. [2 composite samples have similar chemical compositions. To know more about the influence of microwave reaction time on the graphene/MnO2 composites, their microstructure and electrochemical properties have been systematically investigated, as will be discussed below.Figure Figure 6° 002) . The pea2θ = 37° 1 and 2θ  . The pe2 composites prepared with different microwave reaction times. Judging from Figure 2[2 nanoneedles grew almost on the whole graphene surface. Especially from the high-magnification TEM images , galvanostatic charge/discharge (CD), and electrochemical impedance spectroscopy (EIS) tests using a two-electrode system. To avoid complex steps, we adopted a simple method of dipping directly the as-prepared Ni-foam into the suspension of graphene/MnO2 composite, carbon black, and PTFE to prepare electrode slices. This method is simple to operate, and the active materials are not easy to fall off once they have been coated on Ni-foam. Furthermore, the active materials are relatively well distributed. However, there is also a disadvantage, that is, it is hard to control the mass of the active materials. Hence, we cannot get the difference in specific capacitance directly just by looking at the CV and CD diagrams since the mass of the active materials is different in those three samples. Nonetheless, the mass of the active materials can be calculated out by comparatively weighing the Ni-foam before and after coating the active materials, and the specific capacitance can be therefore obtained for each supercapacitor device.To characterize the electrochemical behavior of the button-type supercapacitor based on graphene/MnO2 composites with various scan rates in the range of 5 to 100 mV/s within a 0 to 0.8-V voltage window. It is well known that the more similar the shape of the CV cycles of a supercapacitor to a rectangle, the lower the contact resistance of the supercapacitor [2 gets firstly a little larger and then smaller again, but always tends to distribute more uniformly on almost the whole graphene sheet, and thus the relatively less stacking of graphene might be achieved as well. This could increase the specific surface area, thereby enhancing the electrical conductivity of the device.Figure apacitor ,26. As tapacitor ,28. The 2 composites. All the CD curves are linear and symmetric in the voltage range 0 to 1 V, indicating very good electrochemical properties of the graphene/MnO2 composite electrodes [Figure ectrodes . Using tI is the change/discharge current, t is the discharging time, m is the mass of the active material of two electrodes, and V is the voltage window after the deduction of the IR drop. Accordingly, as the microwave reaction time increases from 5 to 15 min, the specific capacitance at a charging current of 2 mA reaches 246, 260, and 296 F/g, respectively. As aforementioned, though the mass of the active material in each device is different, we can get the specific capacitance values similarly through calculating the data in CV curves. The enhancement of specific capacitance with increasing the microwave reaction time is believed to stem from the improved microstructure, as mentioned above. The results obtained in this work are comparable to or even higher than those reported in the literature for similar graphene-based material systems, where the specific capacitance was reported to be about 200 to 300 F/g [where  300 F/g .2 electrodes with different microwave reaction times, measured in a 6 M KOH alkaline electrolyte solution. The horizontal axis intercepts at high frequency in the Nyquist plots for the devices with microwave times 5, 10, and 15 min are, respectively, 0.70, 0.65, and 1.13 Ω, indicating that with the change in microwave reaction time, the electronic resistance (Rs), including the ionic resistance of the electrolyte, the intrinsic resistance of the substrate, and the contact resistance at the interface of the active material/current collector [Rct) at the electrode/electrolyte interface [Electrochemical impedance spectroscopy is a very significant measure to evaluate the quality of supercapacitors. Figure ollector , changesnterface -36. The nterface ,37-39. A2 composite with improved microstructure.Another critically important factor to evaluate the quality of a supercapacitor is the cycling stability. It is shown in Figure 2 composites have been prepared by a relatively simple microwave sintering method, and a simple method, i.e., dipping Ni-foam into a graphene/MnO2 composite solution directly for a period of time to coat the active material on a current collector, has been proposed for the supercapacitor packaging. It is demonstrated that the microwave reaction time has a significant effect on the microstructure of graphene/MnO2 composites and the electrochemical properties of the supercapacitors based on graphene/MnO2 composites are strongly microstructure dependent. An appropriately longer microwave reaction time, namely, 15 min, facilitates a very dense and homogeneous microstructure of the graphene/MnO2 composites, and thus, excellent electrochemical performance has been achieved in the supercapacitor device based on the graphene/MnO2 composite, comprising a high specific capacitance of 296 F/g and a high capacitance retention of 93% after 3,000 times of charging/discharging cycles. The results obtained in this work pave the way for facile synthesis and optimization of graphene-based composite materials for practical supercapacitor applications.High-quality graphene/MnOThe authors declare that they have no competing interests.XZ designed the study. CZ, ZW, and PS performed the preparation and characterization of materials and supercapacitor devices. YR, JZ (Jiliang), JZ (Jianguo), and DX contributed to the data analysis and discussion. CZ and XZ analyzed the data completely and co-wrote the manuscript. All authors reviewed the manuscript. All authors read and approved the final manuscript.CZ is a postgraduate student at Sichuan University (SCU). XZ has a PhD degree and is a professor at SCU. ZW and PS have a master's degree. YR is a postgraduate student at SCU. JZ (Jiliang) has a PhD degree and is a professor at SCU. JZ (Jianguo) has a PhD degree and is a professor and head of a school at SCU. DX is a professor and head of a research laboratory at SCU.


Radiation induced hardening, evaluated via nanopillar compressions before and after He+ irradiation, is significantly reduced with the inclusion of graphene layers; the flow stresses of V-graphene nanolayers with 110 nm repeat layer spacing showed an increase of 25% while pure V showed an increase of 88% after He+ dosage of 13.5 dpa. The molecular dynamics simulations confirmed that the graphene interface can spontaneously absorb the nearby crystalline defects that are produced from a collision cascade, thereby enhancing the lifetime of the V-graphene nanolayers via this self-healing effect. In addition, the impermeability of He gas through the graphene resulted in suppression of He bubble agglomerations that in turn reduced embrittlement. In-situ SEM compression also showed the ability of graphene to hinder crack propagation that suppressed the failure.Ultra high strength V-graphene nanolayers were developed for the first time that was demonstrated to have an excellent radiation tolerance as revealed by the He Advanced structural material for next generation nuclear facilities require enhanced radiation resistance under harsh irradiation environments to ensure an extended lifetime and therefore the safety of nuclear energy systems24i.e. grain boundary) or heterogeneous (secondary phases) can significantly enhance the radiation tolerance and has been actively researched over the last few decades. One example of using both homogeneous and heterogeneous interfaces in bulk alloys is polycrystalline ferritic/martensitic (F/M) steels with oxide-dispersions (ODS), which is regarded as an effective nuclear structural material due to the ability to absorb crystalline defects at the grain boundaries  as well as the interfaces formed with the oxide phases that are acting as sinks for radiation induced crystalline defects141516Current approaches in enhancing the radiation tolerance include selection of material with reduced scattering cross-section for reduced swelling and engineering of its microstructure to reduce the radiation induced crystalline defects. There are limited metal choices available for nuclear structural applications, and steel is the most widely used material in nuclear facilities. Ni-based alloys are used in applications requiring stability at high temperatures, while Vanadium (V)-based alloys are known to have lower neutron cross section, and reduced radiation induced swelling, and moderate high temperature strength891011et al.22A particularly interesting class of material with enhanced radiation resistance is the metallic nanolayered composite that has been reported to have ultra-high strength181920222325+ irradiation. To further understand the deformation mechanism of V-graphene nanolayers after ion irradiation, in-situ SEM nanopillar compression testing, transmission electron microscopy (TEM) microstructure analysis of the irradiated specimens, and molecular dynamics simulations were performed to understand the enhanced radiation resistance.In our previous work, we reported that graphene with its 2D geometry, outstanding mechanical properties of 130 GPa in strength and Young’s modulus of 1 TPa can be an effective strength enhancer when incorporated in the form of metal-graphene nanolayered composite313334V-graphene nanolayers with 110 nm and 300 nm repeat layer spacings as well as pure V thin film were prepared as outlined in the schematic in The synthesized V-graphene nanolayers with 300 nm repeat layer spacing were analyzed for the initial microstructure using the TEM. Both the micrograph and the selected area diffraction (SAD) shown in + implantation at 120 KeV was used to irradiate pure V, and V-graphene with different repeat layer spacings. The irradiation condition was simulated using Stopping Range of Ions in Material (SRIM)+ ions have sufficient energy to reach down close to the substrate interface (with the penetration depth of ~600 nm on vanadium thin film), thereby fully irradiating the V layers that are to be tested for mechanical properties. The irradiated specimens were analyzed for the presence of crystalline defects and any changes in the microstructure, and detailed cross-sectional TEM image with its SAD patterns with irradiated V-graphene nanolayers with 300 nm repeat layer spacing are shown in 38To analyze the radiation tolerance of V-graphene nanolayers, He+ irradiated pure V and V-graphene nanopillars were then evaluated for radiation induced hardening by comparing the nanopillar compression results before and after the irradiation as shown in et al. for the case of bulk V upon irradiationet al., the He doped bulk V showed hardening of up to twice the original strength and the corresponding brittle failure was reported, where the ductility was reduced from 26% to 4%. For the case of V-graphene nanolayers, however, a significantly smaller degree of embrittlement and radiation induced hardening was observed with the inclusion of graphene, especially with smaller repeat layer spacings or higher density of graphene layers. The radiation induced hardening resulted in increase in flow stress at 5% plastic strain from 4.8 GPa to 6 GPa (25% increase) for the V-graphene with 110 nm repeat layer spacing, from 3.1 GPa to 5.0 GPa (61% increase) for the V-graphene with 300 nm, and from 2.5 GPa to 4.8 GPa (88% increase) for the case of pure V.In the presence of radiation induced defects and He bubble formations, metals can become harder and more brittle, as known as radiation induced hardening. The HeIn-situ SEM nanopillar compression tests were performed for direct observations of the deformation process for irradiated pure V and V-graphene nanolayers with 110 nm repeat layer spacing. For the case of irradiated pure V, brittle failure was observed at 20% strain as shown in As outlined above, He bubble formations are of concern in maintaining the mechanical stability in nuclear structural materials. He is known to rapidly combine with vacancy clusters to form bubbles that can migrate and agglomerate to result in intergranular failure by embrittlement and crck formation4142294445+ irradiation were increased by 40% and 25%, respectively while pure V after irradiation resulted in increase in flow stress of 88%. Significant reduction in radiation hardening in V-graphene nanolayers is expected to be due to reduction in He bubble formation, as well as the ability of the V-graphene interface to absorb the crystalline defects. The total radiation hardening can be expressed asThe increase in yield strength after He implantation is commonly accepted to be due to the formation of radiation induced crystalline defects as well as He bubbles22232447where Radiation induced hardening due to presence of He bubbles  simulations on a model system of pure V and V-graphene nanolayers. We model the radiation event as a collision of a primary knock-on atom (PKA) with 2 keV See . BecauseOne of the main reasons for the reduction in crystalline defects in the presence of the graphene interface is due to the reduction in the vacancy-interstitial formation. The knock-on energy from the collision cascade is typically relaxed through local heating of the crystal lattice as well as by forming vacancy-interstitial pairs, but a large portion of the energy in the case of V-graphene nanolayers is consumed to break the covalent bonds in graphene due to the high atomization energy (7.4 eV) of the graphene. Thus, significant damage of graphene is observed when the knock-on event occurs close to graphene layers see , and thiet al.t) for the vacancy to arrive at the interface asSecond reason for less defect formations is due to the V-graphene interface acting as an efficient sink for the interstitial defects or crowdions. Crowdions move along the [111] directions until they encounter the V-graphene interface, and form adatoms between the graphene and V. Vacancies also diffuse to the V-graphene interface and the arrived vacancies are then trapped at the interface because the vacancy formation energy is smaller at the graphene interface (1.2 eV) than in the bulk lattice (2.5 eV) see . The vac13 Hz, 4 s and 2.43 × 105 s, respectively, at room temperature. However, during the ion irradiation testing and FIB processing to fabricate pillars for compression testing, the actual temperature can increase up to ~200 °C that could then enable vacancies to reach the interface in just a 0.26 seconds for V-graphene with 110 nm repeated layer spacing. Therefore it is expected that the interstitials in the form of crowdion and vacancies gathered at the interface to annihilate near the graphene interfaces. Higher density of the V-graphene interfaces  results in more effective annihilation of interstitials and vacancies, thereby increasing the self-healing ability of the nanolayers. Our analysis is consistent with the experimental observation in which the V-graphene with 110 nm repeat layer spacing showed significantly smaller radiation induced hardening than in 300 nm repeat layer spacing specimen. The self-healing ability as confirmed from MD simulations is expected to be one of the major causes for enhanced radiation tolerance in V-graphene nanolayers, while suppression of He bubbles at the graphene interface also suppress the brittle failure.where + irradiation with dosage of 13.5 dpa revealed that the V-graphene nanolayers had reduced formations of radiation induced crystalline defects compared with pure V, which agreed with the nanopillar compression results of irradiated specimen V-graphene showing reduction in radiation induced hardening and suppression of brittle failure. In-situ SEM compression tests showed that the graphene interface can hinder the crack propagation thus helping to suppress the brittle failure. The structural material degradation due to He bubble formations was shown to be significantly reduced in V-graphene nanolayers due to the impermeability of He through the graphene layer that prevents agglomeration of He gas into large bubbles. Orowan’s model was then used to determine the approximate length scale for the pinning points from the presence of radiation induced defects that indicated less accrual of defects in the V-graphene layers. Finally, MD simulations confirmed that the cause for less crystalline defects in V-graphene layers is due to the graphene interface being able to absorb the crystalline defects that are introduced from collision cascade. Therefore, inclusion of graphene in the form of V-graphene nanolayers can not only result in initially high strength material, but the graphene can self-heal the crystalline defects that are introduced during irradiation as well as terminating migration of He bubbles to result in extremely radiation resistance material.In this study, we developed and analyzed the radiation resistance V-graphene nanolayered composites with varying repeat layer spacings of 110 nm, 300 nm and compared the results to those of pure V. He2 (625 μm/300 nm) substrate was chosen for deposition of the V-graphene nanolayers. Vanadium nanolayers were first deposited using radio frequency (RF) sputtering under vacuum at 5 × 10−6 Torr with Ar gas of 5 mTorr at 10 sccm, and RF power of 200 W was used. After deposition of the V layer, graphene was fabricated by chemical vapor deposition (CVD) and then transferred on to Si/SiO2/V substrate. By repeating the above steps, nanocrystalline V-graphene nanolayers composites were synthesized with 110 nm and 300 nm repeat layer spacings, while keeping the total thickness to be the same at ~600 nm.The Si/SiO−1. The load-displacement data from compression test were analyzed with the conventional constant volume and homogeneous deformation assumption model to calculate the true stress-true strain curvesIn-situ SEM nanopillar compression tests were also conducted to confirm the deformation behavior of the nanopillars using a Hysitron Picoindenter (PI-95) in a Quanta 3D FEG FIB. Same sized nanopillars were compressed under in-situ SEM at a nominal constant strain rate of 0.002 s−1.To test the mechanical properties of the V-graphene nanolayers composites, the Quanta 3D FEG Focused ion beam (FIB) milling was utilized to fabricate V-graphene nanopillars with diameter of 250 nm and height of 600 nm. The pillar compression tests were performed by the Hysitron Tbi-750  nanoindentation system with a 10 μm diameter flat-ended cube corner tip. All pillars were compressed over 15% total strain in a displacement control feedback mode (DC mode) using the Hysitron quasi-static indentation transducer with nominal engineering strain rate of 0.002 s+ ion irradiation was performed at Kyungjoo Korea Multi-Purpose Accelerator Complex (KAERI) facility at 120 keV with beam current of 1.2 mA and dosage of 1 × 1018/cm2 at room temperature. The SRIM calculations indicate (state the simulation conditions and assumptions) that this irradiation condition is sufficient to reach through thickness of the V layers down close to the substrate interface. The benefits of ion irradiation in probing the radiation resistance of materials rather than using neutron irradiations were previously reported in the work by Kiener et al.To evaluate the radiation resistance of V-graphene nanolayers, HeHow to cite this article: Kim, Y. et al. Radiation Resistant Vanadium-Graphene Nanolayered Composite. Sci. Rep.6, 24785; doi: 10.1038/srep24785 (2016).


Similar bandgap has been observed in graphene layers on metal substrates by local probe measurements; however, this phenomenon has not been observed in graphene layers on a near-insulating substrate. Here, we present bulk bandgap-like features in a graphene layer epitaxially grown on an h-BN substrate using scanning tunneling spectroscopy. We observed edge states at zigzag edges, edge resonances at armchair edges, and bandgap-like features in the bulk.Securing a semiconducting bandgap is essential for applying graphene layers in switching devices. Theoretical studies have suggested a created bulk bandgap in a graphene layer by introducing an asymmetry between the A and B sub-lattice sites. A recent transport measurement demonstrated the presence of a bandgap in a graphene layer where the asymmetry was introduced by placing a graphene layer on a hexagonal boron nitride ( In a recent transport study, the presence of a minigap at the neutral point in a graphene layer on an h-BN substrate was confirmedSince its first successful isolation, graphene has been considered as a material suitable for future application in signal-switching devices with spin-sensitive transport1341h-BN, the valley degree of freedom was found to remain intact58911131415In a graphene layer mechanically placed on et al.2223h-BN may better mimic those for a freestanding graphene layer than those for a graphene layer on a metal surface. One may notice that most theoretical predictions were done for a freestanding graphene layer.Scanning tunneling microscopy (STM) is an ideal tool for both imaging local geometric structures and observing local electronic structures. While the above theoretical predictions were derived for a free-standing graphene layer, most previous STM studies were performed with graphene flakes epitaxially grown or drop-cast on metal substrates. Energy gaps of 0.2–0.3 eV were reported in a GNR that was drop-cast on Au(111)h-BN ML on a Cu(111) surface. Edge states of the graphene layer and bulk gap-like features were observed in this pure and well-defined heterostructure.In this study, we performed atomically resolved electronic structure measurements in epitaxially grown GNIs on an epitaxially grown h-BN ML grown on a Cu(111) substrate. After growing approximately half of the h-BN ML on a Cu(111) surface using borazine (B3H6N3), the incoming graphene precursor gas (ethylene) was decomposed on the exposed half of the Cu(111) surface. Most graphene layers were grown on the exposed Cu substrate, while some were grown as GNIs on the h-BN layer. h-BN layer, which showed moiré patterns with many different periods, similar to earlier resultsh-BN surface is ≈1/1000 times than that achieved on a Cu(111) surface. A large dosage (>108 L) of precursor gas was used to grow graphene layers on h-BN263–104 L of ethylene, small GNIs were nucleated on the h-BN surface. The produced GNIs on the h-BN surface were scarce, but they were pure and contained no traces of impurities. The GNIs grown on h-BN were usually located near the h-BN or Cu step edges, as shown in h-BN is usually featureless at a sample bias of −3–3 eV. h-BN at a sample bias of 0.7 V, revealing a featureless flat island. As the bias was increased to 4.0 V, moiré patterns under the top-layer graphene became visible, as shown in We produced a graphene layer on an h-BN/Cu(111), as demonstrated by the different shifts of surface peaks due to different charge transfer from Cu substrate to overlayers: ≈−0.3 eV for graphene, and ≈−0.2 eV for h-BN. A weak dip at ≈−0.4 eV originated from the Dirac point in the graphene layer on Cu(111). In addition to the shifted Cu surface states, the h-BN spectrum had no distinctive features other than ≈4 eV bandgap of the h-BN layer (not shown in h-BN was rather different from that on Cu(111), as shown in 8The identity of the grown layers was confirmed by the STS spectra. The spectra in 2829h-BN, as illustrated in h-BN layer, many defects are present, unlike for free-standing or epitaxial graphene layers on a metal surface. Considering this observation, the zigzag edges seem to be energetically more stable than armchair edges in the graphene layer on h-BN. The spectra change to the bulk spectra with a short decay length near the zigzag edges  modulates within a period of 10–16 lattice constants  and S6 b31h-BN using a tight-binding calculation; the simulated image is very similar to the observed STM image. In a GNR with zigzag edge terminations, the tight-binding calculation predicted a flat band at the Fermi level between 2/3kXΓ and kXΓ in momentum space where kXΓ = 3/4kXΓ and kXΓ is the crystal momentum in the direction from Γ to X in the first Brillouin zone. Unlike GNR, in the case of a GNI with a long zigzag edge, the states at the Dirac point are expected to be reduced but still presentdI/dV at the Fermi level is higher at the edge than in the bulk near the Fermi level), suggesting presence of the edge state. Therefore, the difference between the edge and the bulk is more meaningful than small features at ≈0.05 eV, ≈0.10 eV and ≈0.15 eV at the edge (location 1). dI/dV) subtracted from the bulk dI/dV. These difference spectra clearly show the presence of the edge state at this zigzag edge. The measured edge state at the Fermi level decreases with increasing distance from the edge. The height, and therefore, the strength of the edge state decreases with increasing distance from the zigzag edge. A gap-like feature is visible near the Fermi level, and the size of this gap does not vary with distance from the edge of the bulk stripe. The peak-to-peak energy gap was estimated to be ≈(0.09 ± 0.01) eV, as shown in dI/dV subtracted from the bulk dI/dV, is plotted in dI/dV versus distance. An ab initio calculation was performed to calculate the edge state at five equal distances from the zigzag edge with a periodic boundary of 4 nm ; left side of k = 2/3kXΓ), showing a semiconducting character  single crystal was cleaned by repeated Ar+ sputtering and annealing cycles. Sub-monolayer h-BN was grown using borazine (B3H6N3) at 10‒5 Pa while the sample was maintained at a temperature of 1000 K. Ethylene (C2H4) at 10‒3 Pa was introduced to form graphene on the half-monolayer h-BN-covered Cu(111) surface at 1100 K. The sample was maintained at that temperature for 2 h and then cooled at a rate of 30 K/s. All STM and STS measurements were conducted at approximately 4.9 K. STM imaging was performed under constant current mode, whereas STS was performed under constant height mode.The graphene on the t in the model Hamiltonian was fitted using the measured energy gap and a hopping parameter of t = −2.88 eV. The effective wavevector k adapted to the observed decay length of the edge state was obtained from the analytical solution for the states localized at the zigzag edge. Certain commercial equipment, instruments, or materials are identified in this paper in order to specify the experimental and theoretical procedures adequately. Such identification is not intended to imply recommendation or endorsement by the National Institute of Standards and Technology, nor is it intended to imply that the materials or equipment identified are necessarily the best available for the purpose.Density functional theory (DFT) calculations were performed within the generalized gradient approximation (GGA) using the Vienna ab initio simulation package (VASP). The projector-augmented wave potentials, as implemented in VASP, were employed to describe DOS at atom centers. The energy cutoff for the plane-wave basis set was adjusted to 400 eV in the GGA. The model structures were optimized until the Hellman–Feynman forces acting on the atoms became lesser than 0.01 eV/Å. To include the weak van der Waals interactions between adsorbates and graphene, we adopted Grimme’s DFT-D2 correction based on a semi-empirical GGA-type theory. For Brillouin zone sampling, we used a 27 × 27 × 1 grid in the Monkhorst-Pack special k-point scheme. In addition to the DFT calculation, we also performed a tight-binding calculation with the Hubbard−model Hamiltonian for zigzag GNRs with various widths (3.4–34 nm). The ratio of the on-site repulsive potential U to the hopping parameter How to cite this article: Hwang, B. et al. Energy Bandgap and Edge States in an Epitaxially Grown Graphene/h-BN Heterostructure. Sci. Rep.6, 31160; doi: 10.1038/srep31160 (2016).


However, the low active material utilization, low sulphur loading and poor cycling stability restrict their practical applications. Herein, we present an effective strategy to obtain Li/polysulphide batteries with high-energy density and long-cyclic life using three-dimensional nitrogen/sulphur codoped graphene sponge electrodes. The nitrogen/sulphur codoped graphene sponge electrode provides enough space for a high sulphur loading, facilitates fast charge transfer and better immobilization of polysulphide ions. The hetero-doped nitrogen/sulphur sites are demonstrated to show strong binding energy and be capable of anchoring polysulphides based on first-principles calculations. As a result, a high specific capacity of 1,200 mAh g There is intensive research underway into the cathode development of lithium–sulphur batteries. Here, the authors report a lithium–sulphur battery using nitrogen/sulphur codoped graphene structure which displays excellent electrochemical performance with high sulphur loading. Desphrinkage 0% during2S5 or Cs+ to suppress the polysulphide shuttle and passivate the lithium-metal surface17192021−2 and the sulphur content in the electrode is <70 wt. %, which are not enough to satisfy the demands for high-energy density batteries5122324Much effort has been devoted to solving these issues: improving the electrical conductivity of the sulphur electrode by integrating with carbon materials or conductive polymers891011121315et al.et al.2829Moreover, flat current collectors face the challenge of increasing the electrode thickness due to the easy delamination of active materials and limitation in lithium diffusion kinetics. Therefore, three-dimensional (3D) battery electrode design has recently been explored as an effective way to enhance the energy per surface area of batteries26−2 when the graphene sponge is employed as a 3D current collector. The interconnected graphene network facilitates fast electron and ion transfer. The sulphur and/or nitrogen doping could facilitate fast charge transfer, increase the affinity between the polysulphide species and the carbon-based framework, facilitate better immobilization of the polysulphide ions, and promote the Li2S/polysulphide/S reversible conversion to improve the electrochemical performance of Li–S batteries192122−1 at 0.2C rate, a high-rate capacity of 430 mAh g−1 at 2C rate and excellent cycling stability for 500 cycles with nearly 100% Coulombic efficiency for the N,S-codoped graphene electrode.Instead of adopting solid sulphur as the precursor followed by melt-diffusion or liquid infiltration methods, the dissolved lithium polysulphide systems have recently been proved to make active materials distribute more uniformly, improve sulphur utilization, alleviate electrode structure variation and facilitate the kinetics of the Li–S redox reaction573233342S, GO/urea and GO/thiourea as precursors followed by freeze-drying. Na2S, urea and thiourea were employed as sulphur, nitrogen and nitrogen/sulphur sources, as well as reducing agents for GO during the hydrothermal assembly process. After freeze-drying, the morphology of the samples was retained and no obvious shrinkage/expansion is observed. The size and shape of the graphene-based sponges can be adjusted by controlling the concentration of GO, reaction temperature and the volume/shape of the used autoclaves. The freeze-dried N,S-codoped graphene sponge is light and can stand stably on the top of a dandelion without deforming it at all . The N2 adsorption–desorption isotherms of all the graphene-based sponges show a main adsorption at a relative pressure >0.9 with a small hysteresis loop  is used to evaluate the degree of defects in carbon materials, and a higher ratio indicates an increased amount of defects that is proportional to the extent of reductionID/IG=0.93), the ID/IG ratio increases to 0.99, 1.03, 1.07 and 1.16 for rGO, S-doped graphene, N-doped graphene and N,S-codoped graphene sponges, respectively, suggesting the removal of oxygen-containing groups and the insertion of heteroatoms into the graphene network with more structural defectsID/IG ratio is indicative of the enhanced reduction, which is in agreement with the electrical conductivity tests of these sponges  and Raman spectroscopy . GO show sponges .2 carbon increases and becomes narrower , as shown in narrower , indicatrrower 22, which crrower 22, 46, 47.rrower 22, the thr2S6 solution was used as a reference, the rGO, S-doped graphene, N-doped graphene and N,S-codoped graphene sponges were immersed into the Li2S6 solution for 2 h and the concentration variation of Li2S6 solution was analysed. The typical peaks of the polysulphide solution located at 260, 280, 300 and 340 nm , and the semicircles located in the higher and lower frequency regions, respectively, correspond to the surface film resistance (Rf) and the charge-transfer resistance (Rct) of the batteriesZw) and space charge capacitance (CPE′)Rct of the N,S-codoped graphene electrode is much smaller than those of the rGO, S-doped graphene and N-doped graphene electrodes (2122−1). Compared with the rGO, S-doped graphene and N-doped graphene electrodes, the discharge/charge profiles of the N,S-codoped graphene electrode have an obvious higher discharge plateau at ∼2.32 V (reduction of sulphur to long-chain lithium polysulphides) and a longer plateau at ∼2.10 V (formation of short-chain lithium polysulphides) with corresponding charge plateaus at ∼2.30 and ∼2.45 V (transformation from Li2S2/Li2S to long-chain lithium polysulphides and then to sulphur)To show whether the doped heteroatoms of S and/or N on the graphene have a positive effect on the performance of Li–S batteries, these graphene sponges were cut into slices, compressed and directly used as free-standing electrodes for a series of electrochemical measurements. Different from two-dimensional graphene materials, 3D interconnected porous graphene sponges can provide multidimensional electron transport pathways and easily accessible active sites. Therefore, electrochemical impedance spectroscopy was used to analyse the kinetics of the electrochemical reactions. ectrodes , which c−1, and the capacity is retained at 822 mAh g−1 after 100 cycles with nearly 100% Coulombic efficiency, which is much higher than that of the rGO, S-doped graphene and N-doped graphene electrodes with values of 495 mAh g−1, 646 mAh g−1 and 705 mAh g−1, respectively. Lithium nitrate (LiNO3) is added to promote the formation of a stable passivation film, and the capacity decay at the initial few cycles is partially due to the LiNO3 decomposition and the formation of a SEI film on the lithium anode surface54−1, and the voltage plateau at ∼1.6 V is attributed to the reduction of LiNO3 on the carbon and lithium surface−1 after 100 cycles, demonstrating its little contribution to the overall capacity. To demonstrate the rate capabilities of these electrodes, the current density was changed from 0.2 to 2C rate, as shown in −1, 912 mAh g−1, 675 mAh g−1 and 430 mAh g−1, indicating a high charge/discharge capability. However, the rGO electrode shows almost no capacity at 2C rate, demonstrating poor rate performance. When the current density was reduced back to 0.5C rate, the discharge capacity can be recovered to 878 mAh g−1 for the N,S-codoped graphene electrode, indicating that the electrode structure remains stable after the high-current-density test. Even after the rate capability test, the N,S-codoped graphene electrode still shows excellent long-cyclic performance with higher capacity (550 mAh g−1) and retention (63%) compared with the rGO , S-doped graphene  and N-doped graphene  electrodes at a current density of 0.5C rate up to 500 cycles. The capacity decay is only 0.078% per cycle for the N,S-codoped graphene electrode, much better than those of the rGO, S-doped graphene and N-doped graphene electrodes with decay rates of 0.179%, 0.124% and 0.121% per cycle, respectively. These results imply that the nitrogen- and sulphur-doped 3D porous graphene sponge could not only facilitate fast electron/ion transfer, but also interact strongly with lithium polysulphides to significantly enhance the rate performance and long-term cyclic stability.The cycle performances of these electrodes were first tested at a small current density of 0.2C rate, as shown in −2 by adding more Li2S6 catholyte into the N,S-codoped graphene electrode with the help of a graphene-coated separator, which has been confirmed as a very useful approach to obtain high-performance Li–S batteries2829−1 at 0.2C rate, and the reversible discharge capacity could reach ca. 500 mAh g−1 at a large current density of 2C rate  the 3D graphene network and graphene layer can accommodate the volume change of the active material during cycling; and (iv) the chemical binding between the sulphur/nitrogen heteroatoms and lithium polysulphide/Li2S reduces the active material loss, realizing long cycle life and high-energy/power density Li–S batteries.To further improve the energy density of the whole battery system and satisfy the demands of high-energy batteries, the active material loading was increased from 4.6 to 8.5 mg cm 2C rate . In addi0 cycles . The adv2S) deposition is present on the surface of rGO and the corresponding metallic lithium surface is rough with many coarse lithium agglomerations  is influenced by the incorporation of N and/or S atoms into the graphene lattice. Here we used a small LiSH molecule to model the Li–S end. This model is simple but should be sufficient enough to demonstrate the influence of doping particularly on the binding interaction of the terminal Li+ cation with a doped graphene sheet, although lithium polysulphides may also exist as complex clusters5758+, we also considered Li2S adsorption at a few selected dopant sites, as shown in 2S and LiSH, implying that it would be reasonable and feasible to use LiSH for the purpose of supporting our experimental observations. For a reference, we first considered the binding of LiSH to a pristine graphene sheet and a 1,3-dioxolane (DOL) molecule, which is widely used for the electrolyte in a Li–S cell. As shown in Eb=0.78 eV, in good agreement with the previous DFT resultsEb is given by ELiSH+EGr−ELiSH/Gr, where ELiSH, EGr and ELiSH/Gr represent the total energies of an isolated LiSH, pristine (or doped) graphene and LiSH adsorbed graphene, respectively. Our calculations also predict that the Li of LiSH can be strongly bound to the O site of DOL with Eb=0.93 eV (Eb than that for the graphene case (0.78 eV) may suggest that lithium polysulphides would dissolve into the DOL-based electrolyte instead of adsorbing on the graphene surface, consistent with the existing experimental observations10To better understand the doping effect, density functional theory (DFT) calculations were performed to examine how the binding strength of the Li–S end of linear lithium polysulphides . On the other hand, the quaternary N doping  eV, as compared with the separately doped cases  between 10° and 70° at a scan rate of 0.04° s−1. Raman spectra were collected with a 488 nm laser under ambient conditions with a WITEC Alpha300 S micro-Raman system at room temperature. XPS analysis was performed on a Kratos Analytical spectrometer at room temperature with monochromatic AlKα  radiation. Ultraviolet–visible absorption spectroscopy analysis was performed to evaluate the polysulphide adsorption capability of rGO, S-doped graphene, N-doped graphene and N,S-codoped graphene sponges. About 4 mg each of rGO, S-doped graphene, N-doped graphene and N,S-codoped graphene sponges were separately added into four sealed vials of Li2S6 solution , and the Li2S6 solution without adding anything was used as a control. After absorption for 2 h, the ultraviolet–visible absorption spectra of these solutions were collected with an Evolution 300 UV–vis spectrophotometer with baseline correction. The electrical conductivity of the electrodes was measured by a standard four-point probe resistivity measurement system . Three measurements were taken at different positions on the sample, and the average value was taken.SEM observations were carried out on a FEI Quanta 650 SEM operated at 20 kV. STEM was performed with a Hitachi S-5500 SEM, and EDS was used for collecting elemental signals and mapping. XRD patterns were collected with a Philips X-ray diffractometer with Cu Kα radiation  and LiNO3  in DME  and DOL  (1:1 by volume). The polysulphide catholyte was prepared by chemically reacting sublimed sulphur  and an appropriate amount of Li2S  in the blank electrolyte to form Li2S6 (1.0 M) in the solution. The solution was then stirred at 50 °C in an Ar-filled glove box overnight to produce a brownish-red Li2S6 catholyte solution.The blank electrolyte was prepared by dissolving an appropriate amount of lithium trifluoromethanesulfonate (LiCF2 and a mass of 0.9–1.4 mg. About 12 μl of Li2S6 catholyte (1.0 M) was added into each of the rGO, S-doped graphene, N-doped graphene and N,S-codoped graphene electrodes, corresponding to a sulphur loading of 4.6 mg cm−2. The Celgard 2500 separator was then placed on top of the electrode, followed by adding 30 μl of the blank electrolyte. Finally, the lithium-metal foil anode was placed on the separator as the anode. An Arbin battery cycler was used to perform the galvanostatic cycling measurements at 1.5–2.8 V (versus Li/Li+) at room temperature. The current density set for tests was referred to the mass of sulphur in the cathode and was varied from 0.2C to 2C rate. For the high sulphur loading cathode, 22 μl of Li2S6 catholyte (1.0 M) was added into the N,S-codoped graphene electrodes, corresponding to a sulphur loading of 8.5 mg cm−2. The graphene-coated separator was used to place on the top of the sulphur cathode followed by adding 50 μl of the blank electrolyte. Electrochemical impedance spectroscopy measurements were performed with a Solartron Impedance Analyzer (Solartron 1260A) in the frequency range of 1 MHz to 0.1 Hz with an AC voltage amplitude of 5 mV at the open-circuit potential. For the cycled samples, the cycled cells were disassembled inside an Ar-filled glove box, and the electrodes were rinsed with 1,2-dimethoxyethane solvent to remove the lithium salt and dried inside the glove box at room temperature before analysis.The Li/dissolved polysulphide cells (CR2032 coin cells) were assembled in an Ar-filled glove box. The rGO, S-doped graphene, N-doped graphene and N,S-codoped graphene sponges were dried at 120 °C under vacuum overnight before using. These samples were cut, compressed and shaped into rectangular plates to be used as the current collectors with an area of 0.5 cmAb initio Simulation Package (VASP)z) direction to avoid interactions between graphene and its periodic images. To model N and/or S dopant atoms at graphene edges and vacancies, we employed a graphene flake consisting of 37C atoms in a periodic simulation box of dimensions 20 × 20 × 15 Å; dangling C bonds were passivated by H atoms. A gamma-centred 2 × 2 × 1 Monkhorst–Pack mesh of k points was used for the Brillouin zone integration. All atoms were fully relaxed until the residual forces on constituent atoms became smaller than 0.02 eV Å−1. We used the semi-empirical approach proposed by GrimmeThe atomic configurations and binding energies reported herein were calculated using DFT within the Perdew–Berke–Ernzerhof generalized gradient approximation (GGA-PBE)How to cite this article: Zhou, G. et al. Long-life Li/polysulphide batteries with high sulphur loading enabled by lightweight three-dimensional nitrogen/sulphur codoped graphene sponge. Nat. Commun. 6:7760 doi: 10.1038/ncomms8760 (2015).Supplementary Figures 1-13 and Supplementary Tables 1-4


Sn/Te and the evaporation conditions. It is shown that as-deposited layers with RSn/Te ranging from 0.4 to 0.9 exhibit excellent characteristics as humidity sensors operating at room temperature—very high sensitivity, good selectivity, fast response and short recovery period. Ageing tests have shown that the layers possess good long-term stability. Results obtained regarding the type of the water adsorption on the layers' surface help better understand the relation between preparation conditions, structure, composition and humidity sensing properties.In this paper the humidity sensing properties of layers prepared by a new method for obtaining doped tin oxide are studied. Different techniques—SEM, EDS in SEM, TEM, SAED, AES and electrical measurements—are used for detailed characterization of the thin layers. The as-deposited layers are amorphous with great specific area and low density. They are built up of a fine grained matrix, consisting of Sn- and Te-oxides, and a nanosized dispersed phase of Te, Sn and/or SnTe. The chemical composition of both the matrix and the nanosized particles depends on the ratio R The fir2 sensing layers are their insufficient long-term stability and poor selectivity, i.e., the ability to detect a gas among a mixture of many gases. One possibility to improve these parameters is doping with suitable dopants—metals or other oxides  was measured at intervals of 2 s. The hysteresis of the sensor was measured in the humidity chamber at 25 °C. The resistance variation of the sensor was followed with increasing humidity from 30% to 75% and then with decreasing humidity from 75% to 30%, with steps of 10% RH.As a measure for the sensitivity of the sensing layers the change in their electrical resistance (R [Ω]) as a function of the relative humidity (RH [%]) was used. The resistance was measured within the interval RH = 30%–100% at temperatures of 25 to 80 °C in a test humidity chamber equipped with temperature and RH controller . The measurement of R [Ω] of the layers was performed with a multichannel ohmmeter (MCOM), product of National Instruments . The data acquisition and processing, as well as the control of the ohmmeter, were computerized using LabView software. The experiments concerning the response time and recovery period (kinetic measurements) were performed using saturated solutions of NaOH and . They contained twenty parts of oxygen, ten parts of tellurium and eight parts of tin . The solid state reaction between Sn and TeO2 during the co-evaporation resulted in the formation of eight parts of SnO2, two parts of TeO2 and eight parts of elemental Te. After removal of the elemental Te by the ion beam, the O:Te:Sn-ratio changed to 66.7:6.7:26.7%. As seen in As mentioned above, Auger-electron spectroscopy was used to investigate the homogeneity of the Sn-O-Te layers. However, the direct application of the method is not possible because of the high mobility of the elemental Te present in the layers according to 3.2.3.2.1.3 is much lower than the respective values for the bulk materials TeO2, Te, Sn and SnO2—between 5.7 and 7.3 g/cm3 [Sn/Te ranging from 0.4 to 1.3. As can be seen, the electrical resistance of layers with RSn/Te = 0.4–0.9 decreases with increasing relative environmental humidity spanning 3–4 decades of resistance, the response curve reveals a close exponential relationship between resistance and RH and can be linearized by taking the logarithm of the resistance. Layers with RSn/Te = 1.3 do not show the same dependence, i.e., only layers with RSn/Te < 1 possess high sensitivity to humidity.As the Sn-O-Te layers obtained have a nanosized structure, low density and high specific area respectively  have unambiguously confirmed that with increasing atomic ratio RSn/Te (up to 1.3) the amount of TeO2 decreases and the amount of Sn-oxides and elemental Te increases. At RSn/Te < 1 the layers consist of a mixture of Sn oxides (SnO2 and some SnO) and TeO2, doped with very finely dispersed elemental Te. Obviously, there exists an optimum ratio between the products of the redox reaction (RSn/Te = 0.7–0.9), which causes the high sensitivity towards humidity at room temperature. It is known that the presence of a second oxide at a level of 2–10 wt.% and of other dopants in SnO2-based sensors leads to improved gas response and shifts the maximum of sensitivity to lower operation temperature [Sn/Te in the range of 0.7–0.9 is probably related also to the non-stoichiometry of SnO2, which favors the adsorption of the specific oxygen adsorbate As we have already mentioned, during the vacuum co-deposition of TeOn oxides ,5,9. InvThe release of electrons leads to decreased resistance with increasing humidity. It should be stressed however, that, as noted by other authors too, the mechanism of gas impact on the electrical properties of mixed oxides is much more complicated as compared with homogenous oxides and its detailed understanding needs a more comprehensive study.etc. [The lack of sensitivity to humidity of thermally treated layers is due to significant changes in the in-depth distribution of the layers' components and in the structure .10 Ω). One of the ways to overcome this problem is the additional doping with some metals (in our case with Pt). The doping with 2–3 at.% Pt leads to a resistance decrease of one order of magnitude (shown in the same figure) compared with the undoped layers, without any changes in sensitivity [It is important to note that the high sensitivity to humidity is observed only at room temperature (25 °C). As seen in sitivity . Therefo3, H2SO4, HCl, bases—NH4OH and ethanol) at room temperature was checked. No change in the resistance was observed, i.e., the layers exhibit a very high selectivity.To study the selectivity of the humidity sensitive Sn-O-Te layers their sensitivity towards different vapors  in the presence of moisture [The results presented in this paper demonstrate the applicability of the method developed at our Institute for preparing nanosized doped SnOmoisture . A shortThe layer preparation method allows the introduction of finely dispersed dopants in order to obtain materials sensitive to other reduction gases. It is compatible with conventional microelectronics technologies and after optimization the sensors can be integrated on one chip with field-effect transistors or with entire subsystems for data processing.


Two dimensional atomically thin crystals of graphene and its insulating isomorph hexagonal boron nitride (h-BN) are promising materials for spintronic applications. While graphene is an ideal medium for long distance spin transport, h-BN is an insulating tunnel barrier that has potential for efficient spin polarized tunneling from ferromagnets. Here, we demonstrate the spin filtering effect in cobalt|few layer h-BN|graphene junctions leading to a large negative spin polarization in graphene at room temperature. Through nonlocal pure spin transport and Hanle precession measurements performed on devices with different interface barrier conditions, we associate the negative spin polarization with high resistance few layer h-BN|ferromagnet contacts. Detailed bias and gate dependent measurements reinforce the robustness of the effect in our devices. These spintronic effects in two-dimensional van der Waals heterostructures hold promise for future spin based logic and memory applications. Two dimensional 2D) crystals of graphene (Gr) and hexagonal boron nitride (h-BN) have gained immense attention in recent years because of their exciting properties for nanoelectronics1D crystal456789429151617Recently h-BN has been shown to be a promising 2D material for tunnel barrier applications591966781023In this article we report an inversion of spin signal with large magnitude of spin polarization in non-local graphene spin valves with h-BN|cobalt (Co) high resistance (HR)–low resistance (LR) injector-detector (with different h-BN thickness) combination tunnel contacts. Through systematic nonlocal pure spin transport and Hanle spin precession measurements on devices having different interface barrier conditions, we associate the negative spin polarization with HR h-BN|Co contacts. The characteristic polarization and spatial spin filtering across such contacts leads to relatively large magnitude of spin signal in these devices. We further establish the robustness of the sign of polarization through bias and gate dependent measurements. Our results show the possibility of manipulating the sign and magnitude of spin polarization due to different interface conditions with thickness of h-BN barriers. The large and negative spin polarization using such high resistance (HR) h-BN|Co contacts observed in graphene non-local spin valves indicates spin filtering effects and opens avenues for efficient 2D vdWH spintronic devices.2 substrates. A layer of CVD h-BN |h-BN vdWH devices with ferromagnetic Co contacts. A schematic representation of realizing injection and detection contacts with different h-BN thicknesses along with an optical microscope image of an actual device is presented in CVD h-BN  was tranchniques . DetailsRi = 170 kΩ μm2) and a low resistance (LR) detector (Rd = 25 kΩ μm2). Such a difference in contact resistance is a characteristic feature of CVD h-BN tunnel barriers due to a spatial thickness distribution, which is mostly 1–2 layers with some patches of higher thickness (~3 layers) as reported previously92 substrates  characterization of the injection and detection contacts in a three-terminal configurationbstrates . This thbstrates , which ibstrates . The chator Rd =  kΩ μm2. NLV) for parallel and anti-parallel configurations of the injection and detection ferromagnetic electrodes. This gives a measure of the induced spin splitting (Δμ) in graphene and hence the spin polarization of the Co|h-BN tunnel contacts. The different shape anisotropies of the ferromagnetic contacts ensure distinct switching fields for each of them. In contrast to the normally observed NL spin signal in graphene spin valves  configuration as depicted in the inset of shown in , we obseRNL) by spin precession in a perpendicular magnetic field . We have also found similar spin lifetime in graphene devices prepared with TiO2 tunnel barrier and in other studies4262V−1s−1. The actual tunnel resistances (not Resistance × contact Area) in our device iR and dR >> iR and dR are the actual contact resistances of injector and detector contacts and the channel spin resistance28L and resistance where tal data , we obtaP is considered to be quite large for graphene spin valve devices32NLΔRand P with injection bias current and voltages, exhibiting a decreasing trend at higher bias. It is worth mentioning that the spin lifetime and diffusion constant were seen to be fairly independent of the bias  due to an effective increase in the thickness of h-BN by one layer29Pd ≈ 15%  is responsible for the negative sign of the spin polarization, we performed measurements on a series of devices with different combinations of tunnel resistances. In 2. The normal sign for a spin valve signal in LR-TR combinations 2 also show a normal spin signal without any inversionBased on the sign of the observed spin signal, we group these devices into two categories. First kind, with inverted spin valve signal that has been observed exclusively in devices with a HR-LR contact combination , which pTo understand the robustness of the spin signal inversion, we further studied the influence of the applied injection current bias and gate voltage on the spin signal of our devices at room temperature. gV), which controls the carrier density in graphene, on the sign of the spin signal is examined. In NLΔR for both HR-LR and LR-LR devices in the range gV = ±40 V. In both cases we observe a decrease in the magnitude of NLΔR by 10–15% for higher gV (higher carrier concentration). This behavior is in agreement with theoryTR is much higher than the graphene channel resistance ChR. In such regime, a decreasing trend in the magnitude of NLΔR at higher gate voltage gV is usually observed in lateral spin valve devices. Throughout the gV sweep range, we observe a negative NLΔR for the HR-LR device, which shows that the observed negative spin polarization in a HR Co|h-BN|Gr tunnel contact is stable against variation of carrier density in graphene.Next, the effect of a gate voltage  interfaces4647et al. predicted82O3, MgO have been observed to bear a positive spin polarization2O3-Ta2O3, SrRuO3 and Fe3O4 barriers have also shown negative spin polarization4350515253Spin resolved conductance across the junctions of 2D crystals with ferromagnets has been the subject of recent theoretical and experimental interest due to their atomically thin nature. Karpan In conclusion, we have observed an inversion of spin signal in non-local graphene spin valves with ferromagnetic tunnel contacts with h-BN barrier. Through controlled experiments on several devices having different injector-detector contact resistances, we confirm that the high resistance Co|h-BN junctions with negative spin polarization lead to the observed inversion of spin signal. The effect is found to be robust as established through systematic bias and gate voltage dependent measurements. The enhanced polarization with thicker h-BN junctions indicates considerable spin filtering of spatial nature. These findings bring out the significance of CVD h-BN as a novel spin tunnel barrier, well beyond its use as a 2D insulating barrier that can provide interface resistance for addressing the conductivity mismatch problem. It also demonstrates the possibility of achieving different polarization with different thicknesses of h-BN offering different spin dependent interface conditions and spin filtering. The possibility of high spin polarization in thicker multilayer h-BN together with better quality graphene would be a key ingredient for efficiency in future graphene spintronic devices. In addition, the novel spin filtering effects of ferromagnetic tunnel contacts with thicker h-BN barriers can be explored for engineering spintronic devices with other 2D crystals56572. Graphene flakes are exfoliated on a cleaned SiO2/Si substrate with predefined Ti/Au markers by repeated peeling of highly oriented pyrolytic graphite (HOPG) using the conventional scotch tape technique. The graphene flakes were selected using a combination of optical and atomic-force microscopy. The CVD h-BN layer used as tunnel barrier in our experiment, was grown by chemical vapor deposition (CVD) on a copper foil and has high crystalline quality with ~90% coverage (procured from Graphene supermarket). The CVD h-BN is placed on graphene by wet chemical etching and transfer processes, respectively. The h-BN surface is first covered with PMMA, and then isolated from the Cu substrate by etching in a FeCl3 solution. The h-BN/PMMA layer is washed with 10% HCl and deionized water, and subsequently transferred onto the chip containing graphene flakes. The chip is then baked at 150 °C and washed with acetone to remove the PMMA. The resulting chip containing a h-BN|graphene heterostructure is annealed at 400° C in Ar/H2 atmosphere. The annealing improves the adhesion between the h-BN and graphene and ensures the removal of possible residues of resist remaining from the transfer process of h-BN. The h-BN layer is patterned in regions covering graphene flakes by electron beam lithography and subsequent Ar-etching. In the 2nd step of e-beam lithography, electrodes were patterned and cobalt of 65 nm deposited by electron beam evaporation at a rate of 2 Å/sec at a pressure of <5 × 10−7 mbar with sample being placed in normal position without any specific tilt of the sample stage. The ferromagnetic Co deposition was followed by deposition of 20 nm thick Au capping layer without breaking the vacuum to protect them from possible oxidation. The electrical and spin transport measurements were performed in a liquid helium cryostat with a superconducting magnet.The Graphene devices with Co|h-BN tunnel contacts were prepared on a highly doped Si substrate with 285 nm SiOHow to cite this article: Kamalakar, M. V. et al. Inversion of Spin Signal and Spin Filtering in Ferromagnet|Hexagonal Boron Nitride-Graphene van der Waals Heterostructures. Sci. Rep.6, 21168; doi: 10.1038/srep21168 (2016).


HSP70) gene polymorphisms in the susceptibility to IPF is unknown. The aim of this study was to explore the association between HSP70 polymorphisms and IPF susceptibility in the Mexican population.Idiopathic pulmonary fibrosis (IPF) is a fatal lung disease of unknown etiology. Genetic variation within different major histocompatibility complex (MHC) loci contributes to the susceptibility to IPF. The effect of 70 kDa heat shock proteins (HSPA1L (rs2075800), +2437 of HSP HSPA1L A>G (rs2227956), +190 of HSPA1A G>C (rs1043618) and +1267 of HSPA1B G>A (rs1061581).Four HSP70 single nucleotide polymorphisms (SNPs) were evaluated using real time PCR assays in 168 IPF patients and 205 controls: +2763 C>T of HSPA1B (rs1061581) GA and AA genotypes were associated with a lower risk of IPF compared with GG . Similarly, HSPA1L (rs2227956) AG genotype  and the dominant model AG + GG genotypes were also associated with a lower risk of IPF . In contrast, the HSPA1L (rs2075800) TT genotype was associated with susceptibility to IPF .The analysis of the recessive model revealed a significant decrease in the frequency of the genotype HSPA1B AA (rs1061581) in IPF patients  when compared to controls. Using a multivariate logistic regression analysis in a codominant model the HSPA1 (rs1043618) polymorphisms are associated with a decreased risk of IPF.Our findings indicate that HSPA1B (rs1061581), HSPA1L (rs2227956) and The online version of this article (doi:10.1186/s12890-015-0127-7) contains supplementary material, which is available to authorized users. MUC5B  in both studied groups. No significant differences in the distribution of haplotypes between IPF patients and healthy controls were detected. To calculate linkage disequilibrium in the control group the Haploview 4.2 program was used [HSPA1B (rs1061581) with HSP1AL (rs2075800) and HSPA1B (rs1061581) with HSPA1A (rs1043618).Finally, we also performed a haplotype analysis Table  and we fwas used . In Fig.HSP70 polymorphisms is unknown [HSPA1B 1267A/G polymorphism [HSP70 genes in IPF patients and healthy controls from Mexican ancestry. The most striking findings confirmed by a multivariate analysis were that the GA and AA genotypes of the HSPA1B (rs1061581) polymorphism were associated with a lower risk of IPF. Also, AG genotype and the AG + GG genotypes (in a dominant model) of the SNP HSPA1L (rs2227956) were also associated with a lower risk to develop IPF. On the other hand, the HSPA1L (rs2075800) TT genotype was significantly associated with susceptibility to IPF.Hsp70s, together with their Hsp40 co-chaperones, are the most prominent chaperone families that participate in chaperone-assisted proteosomal degradation of misfolded proteins . Previou unknown . Interesmorphism . Likewismorphism . In thisHSPA1L gene and leads to an amino acid change at position 493 from a non-polar hydrophobic Met to a polar neutral Thr. Amino acid 493 is present in the 18 kDa peptide-binding domain on the beta sheet that forms the floor of the peptide binding groove [HSPA1L polymorphic G allele translates to a Met residue, an hydrophilic amino acid that may affect the interaction of the HSP70 with hydrophobic proteins and consequently impairs its ability to assemble and transport proteins within cells [HSPA1B AG/GG genotype, which seem to protect from IPF, have an increased risk of cancer [HSPA1A, HSPA1B-179 C>T and HSPA1B 1267 A>G have been linked with significant alterations in the mRNA and protein expression [HSP70-HOM +2437 A>G (Met493Thr), affecting the substrate specificity and chaperone activity of HSP-70 HOM, has been described [HSP70 gene-disease associations in IPF are unclear. In this context, strong evidence has revealed the accumulation of unfolded and misfolded proteins with severe endoplasmic reticulum (ER) stress in alveolar epithelial cells lining areas of fibrosis [HSPA1B 1267 A>G polymorphism in the differential expression of HSPA1B mRNA [The polymorphism rs2227956 is located in the coding region of the g groove . The HSPin cells , 39. Intf cancer . Previoupression , 26. Alsfibrosis . It is pA1B mRNA . In addip = 0.005). In comparison with former data of associated functional polymorphisms [HSPA1B [In our study, the combination of the two genotypes associated with protection was only found in healthy controls and not in IPF patients (orphisms , 27, 44, [HSPA1B .HSP70 polymorphisms association with the protection to IPF with a statistical power greater that 80 %, even with the stringently defined number of IPF cases studied. Another limitation was the imbalanced distribution of males and females among IPF and control group. In this regard, a possible bias of the genetic association analysis may be due to the gender disparities. However, in previous genetic-association studies of MHC genes, no marked differences in the MHC genes frequencies have been detected between males and females from Mexican ancestry [To determine whether the associations described in the three models of inheritance are independent of LD between the SNPs analyzed, a multivariate analysis was performed. As a result we confirmed that these associations are independent of LD. The differences in the association with IPF between the three models of inheritance may indicate that the segregation of one or two genotypes of protection or risk act independently, either as protective or risk IPF factors, respectively. The present study has some limitations and replications in other populations are needed to verify these findings, since polymorphisms suggesting fibrosis protection are relatively rare. Notwithstanding, we were able to confirm ancestry . FurtherHSP70 SNPs genotypes between Mexicans and other ethnic groups (Additional file HSPA1L gene is the most common genotype, whereas CC genotype is the least frequent genotype in Italian Caucasians [HSPA1L with rs2227956 are similar in all population groups. Likewise, for the HSPA1A rs1043618 gene, the GG genotype is more frequent in Chinese, Taiwanese, Polish [Finally, we compared the frequencies of ucasians  and Croaucasians  subjectsucasians  and Taiwucasians  as well , Polish , 51 and HSPA1L and HSPA1B genes may influence the susceptibility of developing IPF in our population. Additional investigations with other populations are needed to confirm the significance of our findings and the functionality properties of these polymorphisms.In summary, the present study suggests that genetic variation in the