An efficient route to synthesize macroscopic amounts of graphene is highly desired and bulk characterization of such samples, in terms of the number of layers, is equally important. We present a Raman spectroscopy-based method to determine the typical upper limit of the number of graphene layers in chemically exfoliated graphene. We utilize a controlled vapour-phase potassium intercalation technique and identify a lightly doped stage, where the Raman modes of undoped and doped few-layer graphene flakes coexist. The spectra can be unambiguously distinguished from alkali doped graphite, and modeling with the typical upper limit of the layers yields an upper limit of flake thickness of five layers with a significant single-layer graphene content. Complementary statistical AFM measurements on individual few-layer graphene flakes find a consistent distribution of the layer numbers. Apart from individual graphene flakes, bulk graphene is exploited in numerous systems8. Notably, bulk single-layer and few-layer graphene (SLG and FLG) are proposed as applicable in efficient Li-ion batteries, components of photovoltaic cells, and are viable candidates for spintronics applications10. Thus, scalable methods are required both for high-yield production and characterization techniques. One of the major challenges, i.e., to establish mass-production techniques leading to high-quality SLG and FLG was overcome in recent years12. Amongst the numerous synthesis (mostly top-down) means towards mass graphene production, wet chemical exfoliation methods prevail in terms of material quality and synthesis facility16. Nevertheless, measurement methods suitable for large sample quantities are still lacking.Graphene, the latest discovered carbon allotrope19. It was demonstrated to be an ideal characterization tool not only in the laboratory but also at the mass-production level. In particular, investigations of single-layer graphene are facilitated by the lack of a band-gap in the band structure leading to resonant processes at all wavelengths20. Details of the Raman spectra can reveal the layer number of single graphene flakes or identify the edge of the flakes24, it also enables detailed studies of electrostatic gating, strain, or chemical modifications of graphene36.Raman spectroscopy evolved to be an essential probe for studies of carbon structures especially of nanocarbon38. Bulk analysis of the number of the layers of unmodified bulk graphene is not feasible by optical means. Raman spectra of SLG and FLG single flakes differ significantly in their 2D mode (the overtone of the defect-induced D mode), which provides a way to identify the number of graphene layers, N39. However, twisted multilayers might have 2D modes resembling SLG20. Furthermore, the methods reviewed in Ref. 40, i.e. the so-called C peak and the interlayer breathing modes43, the ZO’ modes. However, this cannot be applied for powder samples built up of a distribution of different thicknesses.Historically, Raman, phase contrast spectroscopy and AFM studies focused on the characterization of separated graphene flakesHere, we present a compelling analysis method based on Raman spectroscopy for analyzing the typical upper limit of the number of graphene layers in a powder sample. We utilize controlled vapour-phase potassium intercalation to distinguish SLG and FLG content through following a stepwise intercalation process. This technique enables us to track the evolution of doping in lightly doped stages, where the Raman modes of the undoped and the doped FLG flakes coexist. The applicability of this intercalation method is demonstrated on a few-layer restacked graphene, poly-dispersed powder prepared by wet chemical exfoliation. The method reveals that the studied samples are composed of graphene flakes with dominantly less than five layers. AFM statistical measurements on individual SLG and FLG flakes of the same material confirm the observed distribution.μm surface, revealing a distribution of flakes on the surface. Representative AFM images of the graphene flakes are presented in Fig. We performed a detailed AFM statistical analysis study on a large number of as-prepared ultrasound treated individual FLG flakes in a single batch chosen randomly. Figure Figure 45 that this approach may be misleading due to the improperly chosen measurement parameters, and complementary studies are required. In particular, partial restacking of the chemically exfoliated graphene flakes on the substrate may lead to bigger aggregates of multiple layers where the graphene sheets are misaligned. AFM, however, is unable to resolve this change of the flake morphology, and cannot identify the thickness of individual flakes.Although AFM-based thickness measurement of individual flakes is a standard method for graphene characterization, it was suggestedμm lateral and vertical resolution, which is large enough for a representative surface average without the biasing effects of nano-imaging.This diversity of the flakes highlights the need for bulk characterization methods, such as Raman spectroscopy. Micro-Raman spectroscopy in our case has about a 5 … 10 46. The corresponding Raman spectra (recorded at 514 nm) are depicted in Fig. 33. The 2D band of the starting material is best fitted with a single Lorentzian line, unlike the composed structure in graphite. Whereas the width of the Lorentzian hints that the material may be a mixture of flakes with different number of layers, the Raman response of the starting material is insufficient to determine the exact distribution of the thicknesses. Here, it is worth to note that FLG flakes are stable on their own (i.e. self/standing), no substrate was used during the Raman measurements.Starting from undoped FLG, we performed controlled temperature-gradient driven potassium doping experiments. Saturation doping was achieved in approximately 10 steps. We intentionally refer to “steps” in our experiments rather than “stages”, as the latter is reserved for the well-known intercalation stages of bulk graphiteinhomogeneous Step 3 and the homogeneous Step 5 reveals that only the intensity ratio of the G modes changes, the position of the 2D modes is unchanged. Step 5 is of particular importance, as this was found to be a stable phase. The highest doping level , we used each time twice longer intercalation times than at the previous Step. In the meantime, we kept a constant and large temperature gradient to homogenize the material at a small and fixed doping level. Comparison of the 8), where similar Raman bands appear upon intercalation. A detailed analysis is given in the Supplementary Materials, and it indicates that the position, the width (ΓFano), and the coupling strength of the electronic continuum, measured by the asymmetry parameter, q, differ. The Cz mode arises from from a vibration, where carbon atoms displace perpendicularly to the graphene sheets. Even though this mode exists at the M point of the Brillouin-zone, it is folded back to the Γ point due to a 2 × 2 doubling of the unit cell during intercalation50. The presence of the Cz mode is a clear indication of a 2 × 2 ordered potassium lattice present on the graphene sheets, or between the graphene layers50. As a consequence, this mode is naturally present in Stage I KC8 and is a clear indication of successful and high doping yield.It is intriguing to compare this spectrum with the Raman spectrum of Stage I potassium intercalated graphite  GIC. Nevertheless, the spectroscopic details are markedly different.The most surprising observation in Fig. 72. Although the doublet structure of the G mode appear similar, two details are different: (i) the G mode with the larger Raman shift lies with about a 6(1) cm−1 difference in the two types of materials, (ii) the 2D mode is markedly different in the two kinds of materials: the FLG contains a 2D mode component with a smaller Raman shift, which is absent in graphite. Albeit these difference may appear to be subtle, these enable us to qualitatively differentiate between the two types of materials. The figure also shows that Step 5 intercalated FLG can be resolved into a mixture of a Stage 3 GIC and the undoped material. This fitting procedure  is capable of explaining both the position of the G mode and the composite structure of the 2D mode. This indicates an interesting scenario for the Raman spectrum of the alkali intercalated FLG: it consists of a mixture of (1) entirely undoped pieces, whose Raman spectrum remains identical to that of the starting material, and (2) relatively highly doped phases .In Fig. 51. In this respect, charges in GIC and our stepwise doped FLG are both inhomogeneously distributed, however, the inhomogeneity is completely different. Intercalation in graphite proceeds from homogeneous and crystalline graphite and the charging inhomogeneity is due to the intercalation itself: it occurs due to the thermodynamic preference for fully doped alkali layers which are inevitably separated by uncharged graphene layers. However in our FLG material, the inhomogeneity is present a priori in the sample (in terms of the different layer numbers in the grains) and the inhomogeneous doping merely reflects this inhomogeneity as we show below.We emphasize that the origin of the doublet structure in GIC is related to the presence of charged and uncharged graphene layers; graphene layers in GIC, which are adjacent to an alkali layer are charged, whereas two which are further apart, remain neutral or uncharged1 around 1580 cm−1 and G2 around 1600 cm−1) are shown in Fig. c band) and uncharged layers (the Guc band), respectively51. Upon doping, the Gc band moves to lower Raman shift beyond experimental error  remains low and varies with stage numbers of the GIC. We note that the Guc band also shifts slightly between the different stages due to strain effect.To gain deeper insight into the composition of the multiple restacked FLG, we analyze the G- and 2D Raman-bands. Intercalation step dependence of the split G-bands  independent of the doping steps, ii) its Raman shift position lies between the position of charged G-band in KC24 and in KC36. This is a strong indication that the G2 band corresponds to graphene layers that appear to be doped as in Stage 2 or 3 graphite. It also means that in our FLG samples, no higher stages (or lower doping levels) can be achieved. Given the heterogeneous nature of the number of layers in the FLG sample, this reveals that our sample is free from flakes with more than 3–5 restacked graphene layers. This observation is in full agreement with our AFM statistical analysis.In FLG, the Gsee Fig.  unveils 1-line barely changes as a function of doping as long as a Lorentzian line fits best the Raman line. This exposes that the induced strain is not affected by the doping level, hence, the G1 line corresponds to a significant amount of undoped flakes. Presence of these undoped flakes along with appearance of Stage 3 doping in five-layer-thick flakes highlights the important contribution of undoped flakes with smaller thickness .Figure ωFano is the measured position of the G-line peak, ωA and ωNA are the calculated adiabatic and non-adiabatic phonon frequencies52. We approximate the latter two quantities with the ones calculated for KC8: ωA = 1223 cm−1 and ωNA = 1534 cm−1, as no exact calculation exists for FLG. This approximation was found to be valid in similar hexagonal carbon systems such as potassium doped multiwalled carbon nanotubes53.To further emphasize the differences between the FLG and the graphite powder, we extract a measure of the charge transfer, the electron-phonon coupling parameter (EPC). The electron-phonon scattering linewidth can be estimated from the positions of the Fano lineshape using the expressionFano is the linewidth of the Fano lineshape. In accordance with previous findings in GICs52, the γEPC of SGN18 and FLG follow the linewidth of the Fano lineshape linearly  doping yields a formation energy gain (ΔF) that decreases for lower stages up to Stage 2, and increases for Stage 1. This behavior is confirmed by calculations for all thicknesses with a layer number n ≥ 3. This staging phenomenon means that all flakes of the sample reach the same stage before a new stage is started to be formed, independently of the layer number. The same effect was found experimentally in bilayer graphene individual flakes, i.e., that the doping occurs first on one of the layers reaching a full Stage 2 doping  before accumulating in-between all layers59.This scenario is supported by first-principles studies2 peak has to be compared with that found in GICs. Secondly, the 2D peak has to be fitted to a combination of Lorentzian curves. The position of the lowest fitted Lorentzian curve must be correlated to the highest stage possible in the material.A practical protocol for the use of the present method is as follows. The chemically exfoliated graphene being studied needs to be intercalated with potassium in a stepwise manner along the protocol, which we present in the Methods Section. Clear evidence for a material, which consists of monolayer graphene only, is when the stepwise doping steps result in a fully intercalated material, without the presence of the intermediate stages (e.g. the G mode shifts continuously). When the intermediate steps are present, the typical upper limit of the layer number can be determined in three steps. First, the position of the GNote that for samples with flakes with a large fraction of 10 or more layers, our method is only capable to prove that the sample contains multilayer flakes.Here, we limited ourselves to determine the typical upper limit of the distribution of the number of layers in chemically exfoliated graphene powder samples. However, further development of our analysis of the Raman spectra could yield more information about the precise statistical distribution.61. Recently, multilayer graphene flakes with ABC stacking were successfully isolated exceeding 17 graphene layers62. This stacking could be identified using Raman spectroscopy as described in recent works62. Specifically, first-principles calculations demonstrated that the 2D band is broader, and it has several components for ABC stacking. In addition, a shoulder emerges at 2576 cm−1 for 1.96 eV laser energy. These Raman fingerprints would be all the more relevant to study due to the interesting arbitrary restacking previously observed in few-layer graphene flakes64.It would be of interest to investigate the stacking of the graphene layers using an extension of our Raman spectroscopy-based technique. Graphene multi-layers typically exhibit two structures of stacking: the more stable AB  stacking and the less stable rhombohedral, ABC stacking. The rhombohedral stacking could lead to promising correlated states displaying superconductivity and magnetismin-situ K doping of FLG samples. Our method uses the combination of the G-band position, its intensity, and the position of the 2D mode components to determine the typically thickest flakes in the sample and to confirm the presence of single-layer ones. The technique works well on FLG powder samples prepared using wet chemical exfoliation technique and was tested for three different mechanical processing routes. Statistical AFM shows that such samples consist of flakes with a non-uniform distribution of the number of graphene layers, and 90% of the flakes consist of mostly less than 5 layers. The Raman spectra of intermediately doped FLG samples can be best described as a sum of two components, corresponding to doped and undoped graphene flakes. The former was argued to arise from five-layer-thick flakes and the latter from flakes made of fewer numbers of graphene layers. Remarkable agreement of our AFM statistical analysis and our Raman data validates our method. Our method provides an efficient way to characterize graphene samples with an arbitrary distribution of the number of graphene layers, where other alternatives  fail.In conclusion, we presented a Raman spectroscopy-based technique to identify the maximal flake thickness in few-layer bulk graphene samples. The presented method is based on studying 66 and as shown in Fig. 67. Further details of the sample preparation are described in Ref. Few-layer graphene samples were prepared from saturation potassium doped SGN18 spherical graphite powder (Future Carbon) using DMSO solvent for the wet chemical exfoliation as described elsewhereμm surface of a Si/SiO2 wafer. Unlike the samples used for Raman studies, only partial restacking may occur due to drop-casting and presence of the substrate in this configuration used for AFM studies. These measurements were carried out using a Bruker Dimension Icon microscope in tapping mode. Bruker Scanasyst-Air silicon tips on nitride levers with a spring constant of 0.4 N/m were used to obtain images resolved by 512 × 512 pixels.The as-prepared bulk samples were characterized by atomic force microscopy (AFM) topographic measurements and images to determine the flake size and the typical number of graphene layers in the bulk material. For AFM statistical analysis, the samples were drop-casted on a 100 × 100 68. We first examined few, incompletely exfoliated nanosheets, which showed clear terraces. As the terrace step heights are multiples of the apparent thickness, we were able to identify the individual layer thickness to be ~1.2 nm.To overcome the difficulties of the determination of the individual layer thickness due to residual solvent, capillary forces, and adhesion, we took advantage of the frequently used step-height analysis procedurein situ under a vacuum of 4 × 10−8 mbar pressure. To perform the in situ intercalation process, we used a similar setup to the one presented and described in Fig. 1 of Ref. 46. Due to lack of substrate, significant restacking of individual graphene flakes occurs, leading to misaligned layers in the sample. The powder sample consists of sponge-like structures, as thermodynamic equilibrium is achieved to create a three-dimensional solid material. As a result of the described sample morphology, Raman measurements provide detailed information not only of single graphene flakes but a large ensemble of flakes that are restacked. As restacking creates only mechanical contact, electronic properties of individual flakes are preserved. Thus, the staging phenomenon necessary for our Raman-based technique occurs at the level of individual flakes.Potassium intercalation was carried out Raman measurements were performed at 458, 514.5, and at 568 nm laser excitation wavelengths. To facilitate our discussion, our manuscript focuses on the observations at 514.5 nm wavelength, only the statistical analysis of the G-band positions in Fig. 69. Potassium with a purity of 99.95% (Sigma-Aldrich MKBL0124V) was evaporated to the sample in several steps in a controlled in situ process. At the first step, the potassium was evaporated for about 2 minutes. In the following steps, we used each time twice longer intercalation times than at the previous Step with constant temperature gradient until homogeneous doping was concluded through Raman measurements of the intensity ratio of the G modes .A power of 0.5 mW was used to avoid laser-induced deintercalation71. The spectrometer was operated with a 1800 grooves/mm grating. A typical 0.5 mW laser power was used with a built-in microscope (Olympus LMPlan 50 × /0.50 inf./0/NN26.5).Following each intercalation step, Raman spectra were recorded on a modified broadband LabRAM spectrometer (Horiba Jobin-Yvon Inc.). The built-in interference filter was replaced by a broadband beam splitter plate with 30% reflection and 70% transmission. The principles of the broadband operation are described elsewhereSupplementary Material


A novel, ultrahigh-sensitivity wide-range resonant micro-accelerometer using two differential double-clamped monolayer graphene beams is designed and investigated by steady-state simulation via COMSOL Multiphysics software in this paper. Along with stiffness-enhanced optimized folded support beams, two symmetrical 3-GPa prestressed graphene nano-beams serve as resonant sensitive elements with a size of 10 μm × 1 μm (length × width) to increase the acceleration sensitivity while extending the measurement range. The simulation results show that the accelerometer with cascade-connected graphene and proof-mass assembly exhibits the ultrahigh sensitivity of 21,224 Hz/g and quality factor of 9773 in the range of 0–1000 g. This is remarkably superior to previously reported studies characterized by attaching proof mass to the graphene components directly. The proposed accelerometer shows great potential as an alternative to quartz and silicon-based resonant sensors in high-impact and highly sensitive inertial measurement applications. Due to their direct frequency signal output, stable performance, excellent repeatability and high sensitivity, resonant accelerometers as an important development of inertial area have been widely applied in aviation, aerospace and automation fields. The resonance frequency-dependent resonant sensitive element is a key part of resonant accelerometers. At present, resonant sensitive elements are mainly composed of precision transverse elastic alloy, monocrystalline silicon, fused quartz and piezoelectric quartz crystal. Resonant accelerometers using these materials generally demonstrate low sensitivity and large sizes ,2,3. TheQ), and then compared the results with the theoretical model of a vibrating beam under tension which indicated the feasibility of resonant graphene sensors [Graphene is a two-dimensional material with a single-layer thickness of 0.335 nm, which is the thinnest material known up to now . Since m sensors . Chen et sensors . Oshidar sensors . Recentl sensors ,16. Addi sensors .15 g level, which is improper for practical applications. Byun et al. also attached micro mass (4.0152 × 10−22–4.0152 × 10−20 g) to graphene nanoribbon to sense acceleration and examined the linear relationship between acceleration and resonance frequency on a log–log scale [2 [Due to excellent resonant characteristics, graphene resonators have been explored to measure the acceleration. Kang et al. investigated highly sensitive resonant accelerometers based on graphene nanoribbon via classical molecular dynamics simulation. They pointed out that the acceleration was closely related with structural parameters and built-in strain of graphene nanoribbon ; howeverog scale . Unfortuog scale . Note thscale [2 . NeverthQ of 9773 in the defined range of 0–1000 g, which is significantly higher than the previously reported research where Q values were confirmed as 1000 [Therefore, in this paper, we propose an architecture of differential resonant accelerometer by sensing graphene beams’ stretching applied by a proof mass. The COMSOL Multiphysics steady-state simulation of the proposal accelerometer without proof mass attached on graphene exhibits the acceleration sensitivity of 21,224 Hz/g and  as 1000 , 7000 [1 as 1000  and 9000 as 1000 , respecty-axis) of accelerometer, and two grooves for placing graphene resonant beams. As shown in y-axis direction. In other words, the fabricated graphene resonant beams operate under double-clamped boundary condition. In this case, the substrate and sensitive mass plate could be fabricated by thermal growth oxidation [xidation ,28, and xidation .In particular, the designed resonant sensing structure is symmetrical along the center of accelerometer, thereby making the whole sensitive structure highly symmetrical. Combined with high axial stiffness and low cross-sensitivity of folded support beams, the sensing structure can fundamentally eliminate the sensitive mass block’s movement caused by the acceleration in other directions and the rotation around other sensitive axes. More importantly, it can guarantee not only high acceleration induced displacement sensitivity and smaller off-axis crosstalk, but also high resonance frequency. Furthermore, the differential resonant structure can enhance the signal intensity and improve nonlinear characteristics of graphene accelerometer, which contributes to increasing the sensitivity and measurement accuracy, as well as better robustness to the interference resulting from conjugate suppression and compensation effect.y-axis) of the sensitive mass plate in vacuum environment. The measured acceleration imposed on sensitive mass plate is converted to a concentrated force by the folded support beams in the sensitive mass plate as shown in Referring to the model in mined as (2)K=2EhThe displacement caused by the folded support beams will result in a movement in one end of the DCGB and then lead to axial stress changes in the graphene resonant beams. The resonant frequency (ribed as (4)fR=n2y-axis) of accelerometer. The end of the DCGB connected to the sensitive mass block is free and the other end is clamped. Assuming the density of graphene (v) are 2208 kg/m3, 1.1 TPa and 0.41, respectively. In this case, we apply a concentrated force load to the sensitive mass block in axial direction (y-axis) to simulate the concentrated force induced by acceleration.The model is imported to COMSOL Multiphysics by utilizing “Livelink for Inventor” module, wherein the initial values of structural parameters of designed accelerometer are listed in Graphene beam is a key element of resonant accelerometer. In view of the aspect ratio of graphene, the module “Membrane Interface” is selected in COMOSL, which is exclusively oriented to prestressed-ultrathin membrane simulation. Then, the module “Swept Mesh” is established by regulating the number, size and distribution of mesh in the thickness direction of membrane, thereby ensuring the computing integrity and accuracy of simulation. It is worth noting that, currently the COMSOL simulations have been widely adopted to investigate mechanical, optical and thermal behaviors of graphene-based beam, coating and ribbon with single-layer, few-layer or multi-layer thickness ,32,33. FAccording to Equations (1) and (2), the dimension of folded support beam has an influence on the equivalent stiffness  for DCGB with zero bending stiffness is solved by using a density value ρ (ρ = 2208.96 kg/m3) larger than theoretical value grapheneρ (grapheneρ = 2208 kg/m3) in this paper. Therefore, there is a small deviation existing between the simulation and theoretical results. That is, the former is slightly higher than the latter, as shown in 10 N/m2 [7 N/m2 to 1010 N/m2. It should be mentioned that, as depicted in inset in 9 N/m2 (3 GPa). For ensuring the DCGB under compression to work in resonant status with high natural frequency (1.954 × 107 Hz) for the proposed accelerometer, the prestress of the designed graphene beam is confirmed as 3 × 109 N/m2 (3 GPa). The prestress of DCGB could be applied by improving the surface roughness of insulating layers and electrodes [Given the resonant frequency shift closely depends on the prestress of graphene beam, a proper prestress should be determined. We investigated the frequency variations with the prestress ranging from 10n et al. , the twoectrodes ,18 or emy-axis). The sensitive mass block in the center is driven into motion and causes the deformation of two support beams and two DCGBs, resulting in a distinct graphene beams’ resonance frequency change of 0.75 × 107 Hz compared to natural frequency (1.95 × 107 Hz).By utilizing the parameters as defined above, we further study the input-output performance of resonant accelerometer in response to measured acceleration. As shown in aforementioned Based on the established simulation model, we can achieve the resonance frequency variation of two DCGB's as a function of acceleration. From Q means stronger ability to resist external interference, better working stability and easier detection of resonance frequency shift. Q factor can be calculated by 7 Hz, which is consistent with the simulation result mentioned above. The corresponding calculated Q value is 9773, which is by far larger than the previous results obtained at room temperature, such as a Q factor of 1000 for the epitaxial graphene resonator in Reference [Q factor of 3000 for the drum graphene resonator in Reference [Q= 7000) in identical dimension order (μm) [Q factor accordingly [The quality factor is one of the important parameters for resonant sensors, which is strongly related to the energy loss speed in each oscillation cycle. For resonant accelerometer, higher eference  and Q faeference . It is ader (μm) . In thisder (μm) . In addider (μm) ,20,21, sordingly ,37.To sum up, the simulation results in this study demonstrate the excellent performance of the suggested differential resonant graphene accelerometer, such as ultrahigh sensitivity, high resonance frequency and wide acceleration range, owing to the excellent mechanical properties of graphene with high prestress, non-proof mass sensitive structure and optimized folded support beam design. It is worth noting that, although the simulation analyses in this paper are performed under non-damping and room temperature conditions, the results obtained herein could provide great potential in applications of high-g accelerometers wherein the thermal and damping effects would be considered.9 N/m2 (3 GPa) based on simulation analysis and theoretical model. The resonant response simulation of the designed accelerometer demonstrates an ultrahigh sensitivity of 21,224 Hz/g, which is significantly better than the results reported in previous studies. Furthermore, a superior Q factor of 9773 is achieved, which well represents the accelerometer’s excellent stability and anti-interference ability. Therefore, these effective efforts validate the advantages of resonant graphene accelerometer over conventional quartz and silicon devices in high-g measurement and provide potential applications in the field of NEMS inertia measurement.An ultrahigh-sensitivity, differential resonant graphene accelerometer which can offer high-g acceleration measurement is designed and investigated via COMSOL Multiphysics in this paper. The simulation results agree well with the theoretical resonance frequency in consideration of prestress in graphene beams, thus verifying the applications of the graphene accelerometer. In order to enable both of the two graphene beams in resonant condition, the prestress of graphene beams is set as 3 × 10


The negative colossal magnetoresistance (CMR) of manganite-cobaltite and positive one of graphene gives the possibility to increase the sensitivity to magnetic field of the hybrid sensor. The performed magnetoresistance (MR) measurements of individual few layer (n = 1–5) graphene structures revealed the highest MR values for three-layer graphene (3LG), whereas additional Co-doping increased the MR values of nanostructured manganite films. The connection of 3LG graphene and Co-doped magnanite film in a voltage divider configuration significantly increased the sensitivity of the hybrid sensor at low and intermediate magnetic fields (1–2 T): 70 mV/VT of hybrid sensor in comparison with 56 mV/VT for 3LG and 12 mV/VT for Co-doped magnanite film, respectively, and broadened the magnetic field operation range (0.1–20) T of the produced sensor prototype.The demand to increase the sensitivity to magnetic field in a broad magnetic field ranges has led to the research of novel materials for sensor applications. Therefore, the hybrid system consisting of two different magnetoresistive materials – nanostructured Co-doped manganite La MR) sensors are widely used in daily applications1, for example, in information storage, where the data are received from a magnetic hard disk with MR read sensor that is very sensitive to low magnetic fields3. A group of materials, so called mixed-valence manganites, Ln1−xAxMnO3 7 are the materials that have been intensively investigated for magnetic sensor applications due to the Colossal Magnetoresistance (CMR) phenomenon10. Indeed, manganite-based MR sensors, consisting of nanostructured thin film deposited on a polycrystalline substrate were already successfully developed11 measuring the magnitude of magnetic field independently of its direction 12. The demand for such MR sensors is increasing rapidly due to the development of advanced scientific and industrial devices and techniques such as non-destructive pulsed-field magnets, electromagnetic launchers, forming and welding, contactless high current measurement technique, etc. However, several existing problems concerning MR sensors have to be solved, like: (i) MR anisotropy (MRA) of nanostructured manganite films at fields <0.3 T (limited by demagnetization or ‘shape’ effect)13, (ii) sensitivity to magnetic field at temperatures higher than room temperature (limited by Curie temperature)14, and (iii) sensitivity at fields close to megagauss which is limited by MR saturation of manganites11. These problems are related to a very rich and complex physics of manganites (due to the electron-lattice and electron-electron interactions in these films) which then leads to the close relation of their structural, electronic and magnetotransport properties, therefore, the investigation of new manganite materials and innovative concepts for such sensors is of highest importance.Magnetoresistive  sensor based on combination of a single layer graphene (SLG) and La0.82Sr0.18MnO3 manganite film was recently proposed and investigated up to 2.3 T20. In the mentioned paper, the hybrid single layer graphene/manganite sensor showed the increased sensitivity in comparison to the individual graphene and manganite sensors. Therefore, further investigations concerning optimization of the magnetoresistive properties and the efforts to increase the sensitivity of these hybrid GM devices are of great interest.At this point, two-dimensional graphene has become a material of scientific and technological interest, due to remarkable electronic, mechanical, and thermal properties making graphene a highly attractive material for a large number of different applications1−xSrx(Mn1−yCoy)zO3 films with partial Co substitution for Mn, as well as the results of a novel hybrid manganite-cobaltite/three-layer graphene sensor prototype are reported.In this study, the comprehensive investigations in magnetic fields up to 21 T of individual graphene structures consisting of few graphene layers and nanostructured manganite La2O3 substrate by wet chemical etching procedure (see Methods). A subsequent transfer of single layers of graphene and placement on top of another graphene layer resulted in few-layer graphene samples. The resonance Raman spectroscopy provides rich structural information on graphene-based materials24. It was demonstrated that the intensity ratio of the 2D band to G (I2D/IG) can be used to probe the number of graphene layers27. Figure 2O3 substrate. The characteristic Raman feature for a single layer graphene is enhancement of I2D/IG ratio . For our studied samples the I2D/IG ratios were found to be 2.48, 1.44, and 0.47 for SLG, 3LG, 5LG, respectively. This is consistent with previously reported data28. In addition, slight G peak frequency downshift was observed with increasing number of graphene layers indicating that I2D/IG ratio trend is associated with changes in number of graphene layers26. The ratio ID/IG sensitively probes the defects in graphene23. For our studied samples the intensity of D band was found to be very low (ID/IG = 0.062 for single layer) indicating high structural quality of the graphene films.Firstly, the structural quality of the prepared single and few-layers graphene was examined. The commercially available CVD grown single layer graphene (SLG) was transferred from Cu foil on AlHowever, the XPS analysis revealed the possible organic contamination with the increase of graphene layers. The residuals of PMMA polymer, used as support layer for the transfer of graphene, could explain the organic contamination which was confirmed by surface sensitive XPS measurements and Auger C KVV spectra. The main component 1 in the C 1 s photoemission spectra see Fig.  at bindi29. Despite the D parameter change to graphitic value of 21.2 eV for 5LG sample, the dominant presence of graphene is determined. XPS and XAES results revealed the existence of polymer traces with the increase of graphene layers, probably due to incompletely removed PMMA, captured at the interface between graphene layers.The characteristic D parameter determined from XAES measurements was in the range of 13.0−13.9 eV see Fig.  for almovs. magnetic flux density for different number of graphene layers (1–5) is presented in Fig. MR behavior of graphene samples was investigated using Corbino disk configuration . This geometry eliminates any Hall voltage contribution and allows achieving the maximal values of longitudinal magnetoresistance (caused by Gauss effect) in graphene. The manetoresistance was determined from the following equation at a certain magnetic field (magnetic flux density B): MR(%) = [(R(B)-R(0))/R(0)] × 100%, where R(0) – resistance in zero magnetic field and R(B) – resistance in magnetic field. The observed MR values subsequently increased for the SLG, 2LG and 3LG and started to decrease for 4LG, 5LG samples, in the investigated magnetic field range  and depends strongly on the number of layers, on the stacking sequence 31 or stacking in rotational disorder between the placed graphene sheets, resulting in turbostratic graphite. The significant influence of graphene stacking sequence on the properties of the system was noticed in the field-effect transistor configuration which was used with changeable gate voltage31. Liao et al.32 observed the reduction of MR of 3LG in comparison with 2LG, with the current applied perpendicularly to graphene. The authors obtained inconsistent decrease of the MR with the increase of number of graphene layers (up to 10) and explained it by decrease of charge carrier mobility evaluated from quadratic magnetoresistance dependence on magnetic flux density MR~B2. In our study, in comparison to ref.32 the current flows in-plane. However, the change of carrier mobility could be also the most probable reason for the MR decrease in multi-layer graphene. We propose, that the grain boundaries/impurities (due to graphene growth and transfer processes) in graphene are acting like scattering centers for charge carriers and reduce their mobility. The vertical transport of charge carriers at graphene/graphene interface allows to “overcome” different inhomogeneities and leads to the reduced electrical resistance and increased magnetoresistance in the few-layer graphene structure. The carrier mobility (µ) values in graphene layers were extracted from the fitting the quadratic part of MR dependences of the experimental curves (MR~(μB)2) presented in Fig. 2/Vs for SLG; 1530 cm2/Vs for 2LG; 2420 cm2/Vs for 3LG; 1790 cm2/Vs for 4LG and 1780 cm2/Vs for 5LG. It is obvious, that three-layer graphene exhibits highest mobility values. High μ values could be explained by small number of defects which is confirmed by intensity of the D-band in Raman spectra  it is possible to induce a large non-saturating linear magnetoresistance, desirable property for practical applications32. Linear MR has been observed in multilayer graphene35, in exfoliated36 and in epitaxial bilayer graphene19 up to 62 T. Our study agrees well with the literature data since we also observe linear magnetoresistance dependence on magnetic flux density. The interesting point is that the region of quadratic (B2) dependence of magnetoresistance differs depending on the number of the stacked graphene layers. In our experiments, in the sequence of SLG, 2LG and 3LG, the quadratic part of MR decreases  and the linear part starts at lower and lower magnetic field. The reduction of quadratic MR could be explained by the interlayer tunneling of charge carriers in few layer graphene in comparison with SLG37. For the multi-layer graphene (4LG and 5LG) the quadratic parts increase again in comparison with 3LG up to 1 T, but still has linear dependence of the MR with higher values in comparison with SLG  magnetic field, the single- and three-layer graphene were further investigated in pulsed magnetic fields up to 21 T. The experiments revealed that room temperature MR of three-layer graphene increased by 2 times (up to 758%) at 21 T . Moreover, the 3LG showed the MR increase in all investigated range starting already from very low magnetic field of 0.1 T.According to the observed 1 T Fig.  in compaMR of 60% at 14 T magnetic field and at room temperature was reported32. In our study, the MR of 300%, at the same 14 T magnetic field, was achieved for transferred SLG. Nevertheless, one has to be aware by comparing the results of different groups since the MR values depend on the growth/production method of graphene, number of the graphene layers, supporting substrate and configuration of the measurement system15. In the case of 2LG graphene Kisslinger and others19 observed MR of about 800% at 21 T and 300 K, which is in a good agreement with our results. Additionally, our results demonstrate that magnetic field sensor consisting of 3LG (instead of SLG) could lead to higher sensitivity to the magnetic field in the range of [0.1−20 T].In comparison to the CVD grown few-layer graphene, the 39 and have tendency of saturation at high magnetic fields11. Therefore, the idea to combine two materials (graphene and manganite) having different magnetoresistance behavior was realized. In order to increase the MR of lanthanum manganite (La1−xSrxMnO3), the La1−xSrx(Mn1−yCoy)zO3 (LSMCO) films were grown by Pulsed-Injection Metal-Organic Chemical Vapor Deposition (PI MOCVD) technique on polycrystalline Al2O3 substrate. The films grew either by layer-then-island or by island growth mechanisms resulting in the polycrystalline structure with the crystallite size varying from 70 up to 120 nm (nanostructured film) and surface roughness of 18−19 nm  exhibit higher resistivity and magnetoresistance values at room temperature43 thus, the substitution of Co for Mn in manganite films could increase the sensitivity of the sensors to magnetic field. Figure MR(%) = [(R(B) − R(0))/R(0)] × 100%. Acording to our results, the certain content of Co (y = 0.14) increased the MR from 55% to 63% at 21 T magnetic field in comparison with undoped manganite film.Previously, it was demonstrated, that the lanthanum manganites exhibit high magnetoresistance values at relatively low magnetic fields nm Fig. -inset. TMR was observed in the whole range of investigated magnetic field, showing the increase of sensitivity to magnetic field for Co-doped manganites not only at high (up to 21), but also at low magnetic fields (lower thant 2 T). The dopping at so-called B-site (Co substitution for Mn), destroys the long range ferromagnetic ordering and Mn3+-O-Mn4+ network, resulting in the change of electric and magnetic properties. Thus, the Co substitution for Mn results in the decrease of the transition temperature from the paramagnetic to ferromagnetic phase and increase of the resistivity, resulting in the increase of room temperature magnetoresistance43. However, to compare MR values of manganite-cobaltite films and MR of graphene layers is complicated and not straightforward. As follows, in manganite-cobaltite case, the zero field resistance (R(0)) was the largest at zero magnetic field due to the negative MR, while it was smallest one for graphene, due to positive MR effect. Therefore, the MR values of manganite-cobaltite films were always less than 100% . In our case, the current was much higher and therefore no dependence of the MR on the flowing current magnitude was observed for various graphene layers having different resistances.In order to increase the sensitivity of graphene at low magnetic fields, the hybrid sensor was designed  and hybrid LSMCO/3LG sensor at magnetic fields up to 0.5 T is presented in Fig. ΔVr is normalized to the source voltage VS: ΔVr/VS = /VS, where Vr,B=0 – response voltage without magnetic field and Vr,B – at magnetic field, respectively. The individual elements (LSMCO and 3LG) were connected in series with the ballast resistors in the same voltage divider configuration, by replacing other element (LSMCO or 3LG) with ballast resistor, respectively, depending on the investigated element (3LG or LSMCO). In order to compare the response of individual elements with the hybrid device, the ballast resistors with the same values as zero field resistance of the other element in the hybrid structure, were used for the measurements. The measurements revealed, that in magnetic fields up to 0.2 T the response signal of individual 3LG was lower than signal of individual LSMCO sample: e. g. at magnetic field of B = 0.07 T it differs more than two times. This could be attributed to different MR behavior of both materials at low magnetic fields: quadratic dependence of graphene and almost linear of LSMCO. Therefore, the connection of 3LG and LSMCO in voltage divider geometry lead to increased sensitivity of the device, in particular, at the low magnetic field range. This effect could be explained by different signs of magnetoresistance - positive MR of graphene and negative MR of nanostructured manganite-cobaltite film /RLSMCO(0). It is obvious, that hybrid LSMCO/3LG sensor exhibits significantly higher response signal in all range of magnetic field in comparison with individual elements. However, depending on the ratio r, one can tune the highest response signal of the hybrid sensor and its sensitivity to the lower or higher magnetic field ranges. The normalized sensitivity of the hybrid sensor was evaluated as S = (δΔVres/δB)/VS. The sensitivity values of individual 3LG, LSMCO elements and the hybrid LSMCO/3LG sensor . These results demonstrate improvement of sensitivity of Co-doped manganite/3-layer graphene sensor in comparison with previously published results on manganite/single layer graphene sensor20 exhibiting normalized sensitivities of 16 mV/VT at 0.5 T and 29 mV/VT at 2 T. However, at 20 T the sensitivity of the LSMCO/3LG sensor drops to 7 mV/VT for r = 0.5, respectively. It has to be noted that in the case of r = 0.1 the sensitivity of hybrid sensor is lower at low fields (20 mV/VT at 0.5 T), however, it is high enough in a wide range of magnetic fields (40 mV/VT at 2.5–7 T and 17 mV/VT at 20 T) and no saturation of the response signal was observed up to 20 T  mV/VT46, however mentioned sensors were operating only in mT range of permanent magnetic field.Additionally, the response signals of hybrid sensor and individual samples in pulsed magnetic fields up to 20 T is presented in Fig. see Fig. . It is d12. Additionally, LSMCO/3LG sensor could be also used at lower operation temperatures since three-layer graphene exhibits even higher MR at cryogenic temperatures .In conclusion, the obtained results show the potential application of Co-doped manganite/3LG sensor for magnetic field measurement application at room temperature in very broad magnetic field range [0.1−20 T] with a certain sensitivity values depending on the magnitude of magnetic field. It has to be noted, that for the device fabrication the sensitivity dependence on the magnetic field can be solved by modern electronics with in advance stored calibration data, converting the measured signals into magnetic field values2O3) with formed Ag contacts in Corbino disk configuration for electric and magnetotransport measurements. PMMA polymer is used as supporting and protecting layer from folding and physical damages of graphene layer. The same transfer procedure was applied in order to get few-layer graphene samples, only the PMMA layer was removed each time with acetone and iso-propanol before placing the next graphene layer on the top of the previous one.The commercially available, high quality monolayer graphene grown by CVD method on Cu foil was used for the experiments. The single graphene layer grown on Cu foil was transferred to target substrate by applying wet chemical etching procedure. During this procedure Cu substrate was etched by ammonium persulfate solution from the bottom, leaving a floating graphene/PMMA flake on the surface of etching solution. Afterword, the SLG + PMMA flake was rinsed several times in deionized water and transferred to the target substrate  equipped with thermoelectrically cooled (−70 °C) CCD camera and microscope. Spectra were excited with continuous-wave diode-pumped solid state laser providing 532 nm radiation. The 50x/0.75 NA objective lens and 1800 lines/mm grating were used to record the spectra. The accumulation time was 50 s. To avoid damage of the sample, the laser power at the sample was restricted to 0.6 mW. The laser beam was focused to ~2 μm diameter spot on the sample. Raman scattering wavenumber axis was calibrated by the polystyrene Standard Raman spectrum. Parameters of the bands were determined by fitting the experimental spectra with Gaussian-Lorentzian shape components using GRAMS/A1 8.0 (Thermo Scientific) software.47 of carbon. Spectroscopic data were processed by the Avantage v.5 software (Thermo Fisher Scientific Ltd.).X-ray photoemission (XPS) and X-ray induced Auger (XAES) spectra were collected by using an Escalab 250Xi  spectrometer with a monochromatic Al Kα excitation source and six-channeltron detection system. The photoemission spectra were collected in standard mode of electromagnetic lenses from the area of about 1 mm in diameter at 40 eV pass energy of the analyzer. For Auger electron spectroscopy (AES) was employed an Escalab MkII spectrometer  with electron gun LEG200 operated at low energy (1–2 keV) and low beam current (few nA). All C KVV spectra (XAES and AES) were acquired at 100 eV pass energy and later smoothed by moving average algorithm for 11 times with a width of 1.8 eV. Afterwards, these spectra were differentiated by using a width of 7 data points for the determination of D parameter1−xSrx(Mn1−yCoy)zO3 (LSMCO) films were grown on polycrystalline Al2O3 substrates by pulsed-injection MOCVD technique with supplying the mixture of precursor solution in certain micro-doses48. The 360 nm thick LSMCO films were deposited at 750 °C temperature and 10 Torr pressure with a partial 3.5 Torr oxygen pressure. The Co doping level in the films (y = 0.14) was chosen in order to achieve higher magnetoresistance values of the films at room temperatures in comparison with manganite films without Co doping. Other deposition parameters  were controlled and optimized in order to get LSMCO films with a certain column-like nanostructure responsible for high magnetoresistance values in a wide temperature and magnetic field range.The nanostructured LaThe morphology of the films was investigated using Atomic Force Microscopy (Veeco Multi-Mode Nanoscope II) (AFM) and Scanning Electron Microscopy (Hitachi SU70) (SEM). The nanostructure of the films was studied using Transmission Electron Microscopy (Tecnai G2 F20 X-TWIN) (TEM).2 atmosphere. The distance between electrodes was constant and equal to 50 μm, whereas the width of the sample was changed to match the zero field resistance ratio between graphene and LSMCO samples.For the magnetoresistance measurements of manganite-cobaltite films the traditional photolithography and mesa etching was applied to form rectangular shaped LSMCO films. The Ag electrodes with a Cr sublayer were thermally deposited over the LSMCO films and post annealed at 450 °C for 1 h in OMR measurements of manganite-cobaltite and graphene samples were performed at 300 K in permanent magnetic field up to 2.35 T by using electromagnet. The mobile 43 kJ pulsed magnetic field generator was used for investigation of the samples in pulsed magnetic field up to 21 T. The generator was capable to generate pulsed magnetic fields of half-sine shape with amplitudes up to 40 T and pulse durations in the range of 0.3–2 ms depending on the chosen inductor (fabricated in the laboratory). The high pulsed magnetic field generator consists of high voltage power supply, capacitor bank, high power switch and a special multi-shot magnetic field coil (inductor). The high-field facilities were designed and are used for the investigations of properties of different materials in high magnetic fields at the Center for Physical Sciences and Technology in Vilnius, Lithuania49.The resistivity and MR measurements showed the linear dependence of magnetoresistance for all graphene structures reaching the highest MR values of 758% at 21 T for the three-layer graphene (two times higher as single one). The revealed highest response of 3LG to magnetic field lead to the development of hybrid magnetic sensor prototype consisting of Co-doped manganite/three-layer graphene system for the sensing applications at room temperature. The prototype sensor demonstrated that combination of materials with positive (graphene) and negative (Co-doped manganite) magnetoresistance effects leads to a significant increase of the sensitivity  and broadens the magnetic field operation range from 0.1 up to 20 T. Moreover, the possibility to tune the sensitivities depending on the magnetic field range during the fabrication stage, by changing the ratio of zero field resistances of 3LG and LSMCO, was demonstrated.In the present study, the quality and magnetoresistance properties of graphene layers (n = 1–5) were measured. The Raman measurements confirmed the good quality of graphene samples, whereas the


FOXC2 and PITX2 variants in CG. Variants of the proximal promoter and transcribed sequence of these two genes were analyzed by Sanger sequencing in a cohort of 133 CG families. To investigate possible oligogenic inheritance involving FOXC2 or PITX2 and CYP1B1, we also analyzed FOXC2 and PITX2 variants in a group of 25 CG cases who were known to carry CYP1B1 glaucoma-associated genotypes. The functional effect of three identified variants was assessed by transactivation luciferase reporter assays, protein stability and subcellular localization analyses. We found eight probands (6.0%) who carried four rare FOXC2 variants in the heterozygous state. In addition, we found an elevated frequency (8%) of heterozygous and rare PITX2 variants in the group of CG cases who were known to carry CYP1B1 glaucoma-associated genotypes, and one of these PITX2 variants arose de novo. To the best of our knowledge, two of the identified variants ; and PITX2: c.535C>A, p.(P179T)) have not been previously identified. Examination of the genotype-phenotype correlation in this group suggests that the presence of the infrequent PITX2 variants increase the severity of the phenotype. Transactivation reporter analyses showed partial functional alteration of three identified amino acid substitutions (FOXC2: p.(C498R) and p.(H395N); PITX2: p.(P179T)). In summary, the increased frequency in PCG patients of rare FOXC2 and PITX2 variants with mild functional alterations, suggests they play a role as putative modifier factors in this disease further supporting that CG is not a simple monogenic disease and provides novel insights into the complex pathological mechanisms that underlie CG.Congenital glaucoma (CG) is a severe and inherited childhood optical neuropathy that leads to irreversible visual loss and blindness in children. CG pathogenesis remains largely unexplained in most patients. Herein we have extended our previous studies to evaluate the role of  CYTOCHROME P450, SUBFAMILY I, POLYPEPTIDE 1  mutations [CYP1B1 [LATENT TRANSFORMING GROWTH FACTOR-β-BINDING PROTEIN 2  [TEK  underlie PCG with variable expressivity [G-PATCH DOMAIN-CONTAINING PROTEIN 3  variants are present in a some congenital glaucoma (CG) cases [CYP1B1 mutations are the predominant known genetic cause of this type of glaucoma in different worldwide populations [LTBP2 gene alterations have been identified only in a few families [MYOCILIN  [FORKHEAD BOX C1  alterations [CYP1B1 genotypes [FOXC1 variants with moderate functional dysregulation as either causative of modifier factors in CG [Primary congenital glaucoma  is a severe and irreversible neonatal or infantile optic neuropathy of uncertain pathogenesis. The immature iridocorneal angle appearance observed clinically indicates that it results from arrested maturation of tissues derived from cranial neural crest cells. This alteration leads to increased aqueous humor (AH) outflow resistance, elevated intraocular pressure (IOP) and optic nerve degeneration. The diagnosis is based in finding isolated trabeculodysgenesis with no other ocular anomalies , 5. PCG  602091)  as the cG) cases . Null CYulations –13, and families , 15. MYOerations , 19 havers in CG –20.FOXC2, MIM# 602402) and PAIRED-LIKE HOMEODOMAIN TRANSCRIPTION FACTOR 2 , are genes encoding transcription factors that control the expression of other genes and form part of a complex regulatory network involved in ocular development and disease [FOXC2 is structurally and functionally closely related to FOXC1. Their chromosomal localization and genomic organization suggest a common origin from an ancestral gene through chromosomal duplications [FOXC2 function cause lymphedema-distichiasis syndrome, an autosomal dominant disease that affects the formation of the lymphatic vasculature system and is also characterized by a double row of eyelashes (distichiasis) [FORKHEAD BOX C2  .PAIRED-LIKE HOMEODOMAIN TRANSCRIPTION FACTOR 2  gene encodes a homeodomain-containing transcription factor. It presents a pattern of expression similar to that of FOXC1 and FOXC2, and it is also present in the periocular mesenchyme. Disruption of PITX2 is one of the major causes of Axenfeld-Rieger syndrome (ARS) [PITX2, like FOXC1, is dose sensitive and mutations that alter the level of functional PITX2 protein (either increased or decreased) by different mechanisms are pathogenic [PITX2 gene is composed of six exons, two alternative promoters located upstream of exons 1 and 4, and two alternative transcription start sites located in exons 2 and 4 [The me (ARS) . ARS is me (ARS) . Systemithogenic . The PIT 2 and 4 . The imm 2 and 4 .FOXC2 and PITX2 variants in CG patients, providing further evidence for the role of rare variants of these two genes involved in ocular anterior segment development as putative modifier factors.In this study we show an increased frequency of rare and functionally altered One hundred and thirty-three unrelated families affected by PCG participated in this study. The study and informed written consent procedures were approved by the Ethics Committee for Human Research of the Hospital Clínico San Carlos, Madrid (Spain), approval number 13/388-E) and followed the tenets of the Declaration of Helsinki. The clinical examination and diagnosis of patients were performed as previously described . GlaucomQIAamp DNA Blood Mini Kit . The DNA sequence variation analyses were carried out by automatic Sanger sequencing. The promoter region (nucleotides -1 to -1500), the coding region and both the 5'- and 3’-untranslated regions (UTRs) of FOXC2 were amplified by PCR, using the primers and conditions described in PITX2 were amplified as previously described [FOXC2 variants were screened for LTBP2 mutations using a customized gene panel [The genomic DNA was extracted from peripheral blood, using the escribed . The varne panel . LibraryFOXC2 variants p.(C498R) (rs61753346) and p.(H395N) were directly amplified and cloned from the genomic DNA of carriers using the following primers, which incorporated the EcoRI or BamHI restriction sites at the 5’ end (indicated in bold): FOXC2-Up-EcoRI, GAATTCCGCGCTCTCTCGCTCTCAGG-3'5'-GG and FOXC2-Dw-BamHI, GGATCCCCGTATTTCGTGCAGTCGTAGG-3'5'-GG. The PCR products were cloned into a modified version of the EcoRI/BamHI restriction sites of the pcDNA3.1(-) vector  [FOXC2 coding sequence from DNA of a control subject and the cloned DNA was completely sequenced to rule out the presence of rare FOXC2 variants. The PITX2 coding variants p.(P179T) was obtained by site-directed mutagenesis using the QuickChange Site-Directed Mutagenesis Kit , with the following primers (mutant nucleotides are indicated in bold): F, ACAGGCTATTCCTACAAC-3’5’-CCTACGACGACATGTAC and R, TGTACATGTCGTCGTAGG-3’5’-GTTGTAGGAATAGCCTG. A cDNA encoding the wild-type PITX2C (NP_000316) isoform, cloned into the pcMV6 vector , was used as a template for the site-directed mutagenesis reaction. The presence of the mutation in the site-directed mutagenesis product was confirmed by Sanger sequencing. The mutagenized cDNA was subcloned into the SgfI and MluI restriction sites of pcMV6 to avoid undesirable mutations in the vector.The R) rs617546 and p.CXCR4 luciferase reporter vector (50 ng), and the pMirTarget vector (50 ng), which expresses the red fluorescent protein (RFP) as a transfection control. The cells were cultured in Dulbecco’s modified Eagle’s medium (DMEM) supplemented with 10% fetal bovine serum (FBS) and antibiotics , at 37°C in a fully humidified 5% CO2 atmosphere. Forty-eight hours after transfection the cells were harvested and assayed for firefly luciferase activity using the Luciferase Assay System (Promega) according to the manufacturer’s instructions. FOXC2 and PITX2 proteins were detected by Western blot as expression control and transactivation activity was normalized for these proteins. RFP and endogenous lactate dehydrogenase (LDH) were also assessed by Western blot as transfection and loading controls, respectively.Transactivation assays were performed using the Luciferase Assay System (Promega) following the instructions of the manufacturer. HEK-293T cells were transfected with the different recombinant versions of FOXC2 and PITX2C, cloned in the expression vector pcMV6 (500 ng each), along with the recombinant pGL3-basic-The protein stability was studied via western blot of transfected cells treated with cycloheximide (300 μg/ml) at different times as previously reported . The FOXFor Western blot analyses the cell lysates were prepared and fractionated using sodium dodecyl sulfate-polyacrylamide gel electrophoresis (SDS-PAGE), using the Mini-PROTEAN III Gel Electrophoresis System . A commercial mouse monoclonal anti-myc  was used as the primary antibody, diluted at 1:1000. Horseradish peroxidase-conjugated antibodies against mouse IgG  were diluted to 1:1000–1:4000. Chemiluminiscence detection and the densitometry for protein band quantification was performed as previously described . As an ahttp://genetics.bwh.harvard.edu/pph2/) programs. The MicroSNiPer software was used to identify putative miRNA target sequences (http://cbdb.nimh.nih.gov/microsniper) . Both nucleotide and amino acid sequence alignments were carried out with Clustalw . The novel FOXC2 and PITX2 variants were named using directions from Mutalyzer , according to RefSeq NM_005251.2 and NM_000325.5, respectively. The novel variants have been submitted to dbSNP (http://www.ncbi.nlm.nih.gov/SNP/) and to ClinVar (https://www.ncbi.nlm.nih.gov/clinvar/). Structure prediction of mutant and wild-type proteins were evaluated by QUARK (https://zhanglab.ccmb.med.umich.edu/QUARK2/) [The deleterious effect of mutations was predicted online with the Sorting Intolerant From Tolerant (SIFT)  and PolyPhen-2 (QUARK2/)  and furtt-test or the one-way analysis of variance (ANOVA). A Bonferroni correction was applied to adjust tests for multiple comparisons. The data were statistically processed by the SigmaStat 2.0 software .The statistical comparisons between groups were performed using either the FOXC1 variants [FOXC2 and PITX2 variants in CG using the same group of patients. The main clinical and demographic features of these patients have previously been reported [MYOC and CYP1B1. Eight of the patients (6.0%) carried rare FOXC2 variants in the heterozygous state , which is annotated in Ensemble and GeneCards databases as a promoter covering the entire FOXC2 sequence and most of the antisense lncRNA FOXC2-AS1 gene (FOXC2-AS1 intron 1 (n.145+174G>A) and the regulatory feature (promoter) (FOXC2 and a modifier outcome on both FOXC2-AS1 and the overlapping promoter. The nucleotide change c.*38T>G was located in the 3’UTR of the gene (FOXC2 nucleotide change was predicted to cause a novel amino acid change (p.(H395N)) located in the second transactivation domain of the protein  , and it the gene  and was the gene . This vathe gene  in the tthe gene , andacco protein  with inf protein . The str protein . In addipromoter . The patpromoter .FOXC2 and PITX2 defects on patients with CYP1B1 glaucoma-associated genetic backgrounds. To that end, a group of 25 CG cases who were known to carry CYP1B1 pathogenic genotypes, were selected based on DNA availability from a previously reported cohort of patients [PITX2 variants were identified in two (8%) of these patients (PCG-133 and PCG-139). As far as we know, one of these variants , CYP1B1 variants none of them presented this PITX2 nucleotide change (de novo. The patient who carried this variant (PCG-133) was diagnosed with bilateral PCG at the age of three months and presented advanced optic nerve excavations and was subjected to six surgical operations per eye to control IOP , that resulted in an inferred missense mutation (p.(A188T)), ),  also locA188T)), . This nuA188T)), . The varA188T)), . The carFOXC2 and PITX2 wild-type amino acids or nucleotides were highly conserved , could be cloned for functional evaluation. It showed approximately 130% increased transactivation compared with the wild type  stability was not significantly different from that of FOXC2 wild-type protein  and c.*38T>G (non coding 3’ UTR) also mapped at a promoter, which overlapped with FOXC2 and FOXC2-AS1 genes. FOXC2-AS1 encodes a lncRNA transcribed from the negative strand of FOXC2. Both molecules have been located in the cytoplasm and are concordantly upregulated in doxorubicin-resistant osteosarcoma cell lines and tissues [FOXC2-AS1 stabilizes FOXC2 mRNA by forming an RNA-RNA duplex, which subsequently upregulated steady-state FOXC2 expression at both the transcription and post-transcription levels [FOXC2 activity through FOXC2-AS1 dysregulation. Further studies are required to confirm these ideas. Finally, the high evolutionary conservation of wild type FOXC2 and PITX2 nucleotide or amino acid positions involved in the variants identified in this study, further indicate that their substitution by infrequent residues might modify their function.The two missense FOXC2 variants (p.(H395N) and (p.(C498R)) and one of the PITX2 amino acid substitutions (p.(P179T)) were inferred to cause a moderate functional effect at least by one bioinformatic analysis and, experimentally, they were found to be associated with moderately disrupted transactivation. The functional impact of the second PITX2 amino acid substitution, p.(A188T), could not be functionally evaluated due to DNA cloning difficulties. In fact, the two FOXC2 amino acid changes were found to be hypomorphic whereas the PITX2 amino acid substitution (p.(P179T)) behaved experimentally as a hypermorphic variant. Additional structural and functional analysis indicated that p.(H395N) alters polypeptide chain conformation and decreases protein stability, which can explain the associated reduced transactivation. However, our data showed that the transactivation effect of p.(C498R) is not due either to altered protein stability, protein structure or subcellular localization, suggesting that the substitution of a basic (arginine) for a polar (cysteine) amino acid may disrupt protein interactions in the second inhibitory domain of the protein, leading to reduced protein activity. Similarly, our results indicate that increased transactivation associated with the p.P179T PITX2 mutation is not related with altered protein stability, protein conformation or subcellular localization. Proline possesses a hydrophobic side chain, whereas threonine side chain has both hydrophilic and hydrophobic functions. Therefore, this amino acid replacement may affect protein interactions taking place in the transcriptional inhibitory domain where it is located, leading to increased PITX2 activity. In this line, an increased side chain polarity associated with amino acid substitution p.(A188T) could also interfere protein interactions involving the first PITX2 transcriptional inhibitory domain, leading to a functional alteration. Additional studies are required to evaluate these hypotheses. Interestingly, according to Ensembl Regulatory Build,  tissues . It has n levels . TherefoFOXC1 and PITX2A. In fact, it has been reported that PITX2A protein can act as a negative regulator of FOXC1 through formation of PITX2A-FOXC1 heterocomplexes [PITX2/CYP1B1, PITX2/FOXC1 [FOXC1/CYP1B1 [CYP1B1/GJA1 [CYP1B1 and MYOC has been reported in early-onset congenital glaucoma [PITX2 and CYP1B1 variants and showed very early glaucoma onset and severe disease, characterised by a large number of surgical operations required to control IOP. These data suggest a modifier role of rare PITX2 variants on age at onset and severity on CYP1B1 associated congenital glaucoma. Our data also revealed the presence of one PITX2 de novo mutation in one PCG patient, indicating that this type of genetic variants might also be involved in congenital glaucoma, adding additional genetic complexity to the genetic basis of this disease. We should keep in mind that reduced penetrance, low frequency of these variants and limited number of available patients due to low prevalence of the disease, will also complicate studies to determine the role of this type of genetic variation in CG. Combination of genome, transcriptome and even proteome analysis of individual patients, along with functional studies in animal models will be required to gain new insights into the genetic complexity of congenital glaucoma.Modifier genes are increasingly recognized as an important source of phenotypic variation that may explain the relationship of phenotype to genotype. Individually, genetic modifiers have small effect size on the phenotypic outcome, leading to reduced penetrance, i.e., not all individuals harbouring a pathogenic variant will exhibit clinical signs of the associated disorder . Normal omplexes . Therefoomplexes , 46, andomplexes , 47, 48.X2/FOXC1 , FOXC1/C1/CYP1B1  and CYP11B1/GJA1 ) provideglaucoma  and PCG glaucoma . LikewisIn summary, this study further supports the concept that mild functional alterations associated with variants in genes involved in ocular development may act as modifier factors in non-Mendelian CG cases, providing novel insights to understand the genetic complexity of this disease.S1 FigArrows over the peaks indicate the location of mutations. The mutant nucleotides are indicated in red and those in heterozygosis are shown below the corresponding wild type position. The sequences of the different variants were obtained from the corresponding probands in each family and the wild type sequences from control subjects.(TIF)Click here for additional data file.S1 Table(DOCX)Click here for additional data file.


Meanwhile, the sensing performance and sensing mechanism of the sensors are discussed. The improved experimental schemes are raised and the critical research directions of graphene/metal-oxide sensors in the future are proposed.Owing to the excellent sensitivity to gases, metal-oxide semiconductors (MOS) are widely used as materials for gas sensing. Usually, MOS gas sensors have some common shortages, such as relatively poor selectivity and high operating temperature. Graphene has drawn much attention as a gas sensing material in recent years because it can even work at room temperature, which reduces power consumption. However, the low sensitivity and long recovery time of the graphene-based sensors limit its further development. The combination of metal-oxide semiconductors and graphene may significantly improve the sensing performance, especially the selectivity and response/recovery rate at room temperature. In this review, we have summarized the latest progress of graphene/metal-oxide gas sensors for the detection of NO The ele2. The sensor showed high sensitivity to NO2 at low concentrations. In another work, Hu et al. [3 with an ultra-fast response time of 1.4 s at room temperature. The sensors based on rGO exhibited a rapid and high response to target gas at room temperature. However, these sensors show a common shortage. Since the binding force between graphene and gas molecules is van der Waals force or even covalent bonds [The mass production of single-layered graphene is difficult. Another problem is that pristine graphene does not have a bandgap, which means it is not suitable for semiconductor gas sensors . Graphenu et al.  fabricatnt bonds , the recnt bonds –20.2), titanium dioxide (TiO2), zinc oxide (ZnO), copper oxide (CuO), tungsten oxide (WO3), indium oxide (In2O3), ferric oxide (Fe2O3) and cobalt oxide (Co3O4) are important materials for gas sensors [2−. With the increase of temperature, O2− is transformed into O−. When the temperature is above 397 °C, the oxygen anions are converted into O2−. The reaction equations are as follows:Metal-oxide semiconductors (MOS), including tin oxide . In general, people should not be exposed to an environment with more than 35 ppm NH3 for more than 15 min or an environment with more than 10 ppm CO for more than 10 min. Also, people should not be exposed to workplaces with more than 0.08 ppm formaldehyde for more than 30 min according to World Health Organization [There are many kinds of toxic gases from industrial processes and car emissions around us, such as NOnization . Therefonization . Compare2–rGO sensor [2 and rGO formed p–n heterojunctions. The enhancement mechanism of the p–n heterojunction is shown in Ec and Ev are the energies of conduction band and valence band of the two components, respectively, while the Fermi level energy (Ef) is between these two bands. It can be seen from 2, meaning that electrons transfer from rGO to SnO2 in the heterojunctions. The Schottky barrier is only 0.2 eV due to the changed Fermi level of the composite structure after achieving a dynamic balance of the electron flow, indicating that the electrons are able to pass through the energy barrier. In summary, the SnO2–rGO sensor allow for the transition of electrons even at room temperature because of its low Schottky barrier. When exposed to air at room temperature, oxygen molecules obtain electrons from n-type SnO2–rGO hybrids to form O2−. The electron depletion layers generated on the interfaces of SnO2 grains and rGO sheets owing to the loss of electrons and the resistance is the initial resistance of the sensor. The initial resistance of SnO2-rGO sensor is much lower than that of a SnO2 sensor since because of the high conductivity of rGO, indicating that variations of resistance can be detected at room temperature. Moreover, the electron depletion layers in the SnO2–rGO sensor where the electrons are constantly moving between SnO2 and rGO are wider than those in a SnO2 sensor. When exposed to reducing gases , the gas molecules act as electron donors to SnO2–rGO so that the resistance of composite sensor decreases dramatically, leading to high sensitivity and rapid response.Reduced graphene oxide (rGO), which plays the role of a p-type semiconductor, can form heterojunctions when forming composites with most metal-oxide semiconductors. In the example of a SnOO sensor , SnO2 anIsotypical p–p homojunctions are formed when rGO forms composites with most p-type metal-oxide semiconductors. The enhancement mechanism of p–p homojunctions is similar to that of p–n heterojunctions. However, p–n junctions cause the recombination of charge carriers with a decreased amount of charge carriers, while p–p junctions lead to spatially separated charge carriers with an unchanged amount of charge carriers. The following summarizes the parameters of several room-temperature gas sensors based on rGO/metal oxides, which exhibit enhanced sensing performance mainly due to the formation of heterojunctions.3 with ultra-fast response/recovery times of 78 s/188 s, which was much better than that of a pure rGO sensor (low response and endless recovery time). The composite sensor with the optimal amount of GO (1.5 mL) was highly sensitive to low concentrations of NH3 and responded dramatically, which could be ascribed to the p–n heterojunctions formed between ZnO and rGO. However, the influence of humidity on the sensor response was not negligible due to the residual oxygen (high hydrophilicity) on rGO and the active sites (oxygen vacancies) on ZnO. Chen and co-workers [2–rGO sensor, which was synthesized via growing SnO2 nanorods on a GO surface, showed a response of 1.3 to 200 ppm NH3 with instant response/recovery times of only 8 s/13 s at room temperature. The rapid response and particularly the ultra-fast recovery were extremely inspiring because the recovery time of sensors based on rGO is usually long at room temperature. As a comparison, the pure SnO2 sensor is only able to work at 160 °C. The outstanding sensing performance was accounted for by the p–n heterojunctions according to the explanation of the authors. NH3 adsorbed on rGO has a smaller adsorption energy than other gases. Strong hydrogen bonds were formed between hydrogen atoms (NH3) and the residual oxygen atoms on rGO, facilitating the interaction of NH3 with rGO. Thus the selectivity to NH3 was good. The authors found that water molecules affected the sensor at low humidity levels. However, the response of the composite sensor increased with increasing humidity at high humidity levels. At low humidity levels, ammonia and water molecules competed for adsorption, which reduces the sensing performance of NH3. As humidity levels increased, ammonia adsorbed on the surface of sensor by dissolving into water, leading to the higher response. This is quite different from the usual opinion that sensitivity will decrease at high humidity levels.Tai et al.  deposite-workers  demonstrEg1 (1.2 eV) is not equal to Eg2 (0.4 eV) [3 was not as well as that of a pure ZnO sensor, the ZnO–rGO sensor exhibited rapid response/recovery times with a detection limit of 50 ppb while the pure ZnO sensor was unrecoverable at room temperature. The authors stated that the p–p junctions constituted ZnO and rGO reduced the response/recovery times dramatically. An interesting phenomenon found by the authors was that the pure ZnO sensor showed characteristics of a p-type semiconductor during the test. Two reasons were given for this. One was that a little carbon and nitrogen might be doped into ZnO due to the addition of graphite and nitrogen gas protection during the carbothermal reduction (1150 °C). Another was the formation of Schottky barriers between ZnO and the metal electrodes, which caused ZnO to exhibit p-type semiconductor properties. The gas-sensing performance parameters of the abovementioned sensors for reducing gases based on metal oxides and rGO enhanced by the formation of semiconductor interfaces are listed in Zhang et al.  synthesi(0.4 eV) –51. The (0.4 eV)  mixed Zn2–rGO sensor as an example [2 and rGO formed p–n heterojunctions during the recombination process. The Fermi level of SnO2 is higher than that of rGO since the work function of SnO2 (4.55 eV) is lower than that of rGO (4.75 eV). As a result, electrons transfer from SnO2 to the conduction band of rGO, not only leading to the bending of the energy bands but also forming potential barriers at the interfaces between SnO2 and rGO. When exposed to air at room temperature, oxygen molecules form O2− by obtaining electrons from p-type SnO2–rGO hybrids. Thus, hole accumulation layers and potential barriers are generated at the interfaces of SnO2–rGO hybrids. There are wider hole accumulation layers and higher potential barriers in SnO2–rGO than in pristine SnO2 owing to the existence of the immanent potential barrier and the imprisoned electrons in rGO. When exposed to an oxidizing gas (NO2), which has a high electron affinity, the electrons are continuously captured from SnO2–rGO hybrids, leading to wider hole accumulation layers and increased potential barriers. Consequently, the resistance of composite sensor decreased sharply, leading to high sensitivity and rapid response time. The reaction equations are as follows:Similarly, taking a SnO example , as disp example , SnO2 an2 and the response/recovery times were only 25 s/15 s with good stability at room temperature. As a comparison, the response of the ZnO sensor was 6.2 and the response time was over 150 s under the same conditions. The pure rGO sensor only showed a response of 1.2 and a recovery time up to 58 s. The effect of humidity on this sensor was also tested. It was found that the sensor was relatively stable to different humidity levels. The authors ascribed the enhanced sensing properties to both the p–n heterojunctions and the increased carrier concentrations. Specifically, the addition of rGO increased carrier concentrations and provided conductive pathways, which favored the transfer of electrons. The CeO2–rGO sensor reported by Jiang et al. [2 sensing performance of CeO2–rGO sensor. Zhang et al. [3)3·9H2O solution, then synthesized α-Fe2O3–rGO hybrids via hydrothermal method. During the process of testing, the authors found that the doping amount of graphene significantly affected the sensing properties of the α-Fe2O3–rGO sensor. The α-Fe2O3–rGO sensor with the optimal doping amount of 12.2% showed the highest sensitivity and rapid response time to NO2 at room temperature. In addition, the sensor exhibited excellent selectivity to NO2 because other interference gases required high operating temperatures to react with the surface of this sensor. As a comparison, the pure α-Fe2O3 sensor does not work at room temperature. The formation of p–n heterojunctions was an important factor for the exceptional sensing performance. There were two different depletion layers and potential barriers in the composites due to the heterojunctions, one at the α-Fe2O3 grain boundaries and another at the interfaces of α-Fe2O3 and rGO. It appeared clear that O2 adsorption on the α-Fe2O3 grain boundaries modified the width of the α-Fe2O3 depletion layer, which conversely altered the width of depletion layer at the α-Fe2O3 and rGO interfaces leading to higher sensitivity. In addition, the periodic exposure of the α-Fe2O3–rGO sensor to 0.1 ppm NO2 indicated that the process of response was repeatable. However, a drift of the baseline was easily noted, which usually appears in room-temperature sensors, and further optimization is needed to control the drift. In another work, p–p junctions were accounted for the outstanding sensitivity to NO2 of Co3O4–rGO sensors at room temperature [In an early work, Liu et al.  stated tg et al.  added GOperature .2 nanoparticles were grown on S-rGO sheets to prepare SnO2–S-rGO hybrids. The SnO2–S-rGO sensor exhibited exceptional sensitivity to NO2 with a detection limit of 450 ppb at room temperature. The sensor also showed good repeatability and was not affected by water molecules. In contrast, a sensor based on S-rGO showed high sensitivity to NO2, but its response/recovery times were long (more than a few minutes). A gas sensor based on SnO2–rGO [2 in their previous study, but it could only work at 50–55 °C. The authors ascribed the better sensing performance of the SnO2–S–rGO sensor to the following factors: The addition of sulfonic acid enhances the dispersibility of rGO. At same time, the conductivity of S-rGO is better than that of rGO. Most importantly, the p–n heterojunctions between S-rGO and SnO2 lead to the good sensing properties. Because noble metals offer exceptional catalytic activity, the same group [2–rGO ternary hybrids by reducing AgNO3 on the dispersion of SnO2–rGO, and the sensing properties were tested at room temperature. The authors demonstrated that the response and recovery of this sensor were much faster than that of SnO2–rGO sensor, which could only work at 50–55 °C. The doping of Ag nanoparticles not only improved the electron transfer rate of the sensor, but also increased the number of active sites on the surface of the sensor. Moreover, the introduction of Ag nanoparticles reduced the Schottky barrier of the ternary composites so that those electrons with lower energy were able to cross the energy barrier at room temperature. Thus the sensor can work near room temperature. The most important reason for the excellent NO2 sensing performance of the ternary composite sensor was still the p–n heterojunctions. The gas-sensing performance parameters of the abovementioned NO2 sensors based on metal oxides and rGO enhanced by the formation of semiconductor interfaces are listed in Liu et al.  preparedSnO2–rGO  also exh2 nanoparticles prevent graphene from agglomerating, which in turn leads to a high specific surface area. Graphene enhances the conductivity of the composite materials, enabling the composite sensors to achieve a high response at low operating temperatures. Moreover, the introduction of graphene provides more adsorption sites at the surface of the composite so that the response can be significantly improved.By changing morphology and structure of the sensor materials, one can obtain large specific surface area, more conductive pathways and more active sites, which significantly improve sensing performance of the sensor. As shown in 2 at room temperature, while the pure rGO sensor only exhibited 7.7% response under the same conditions. Apart from the conductive 3D network channels provided by rGO, the large specific surface area of the composite sensor also contributed to the high response. Liu et al. [2, GO, sodium acetate and sodium citrate in an autoclave. Different from other drying processes, the hybrids were obtained through freeze-drying to keep the 3D structure. ZnO nanoparticles were well wrapped in graphene sheets, while the graphene sheets were well dispersed in the hybrids. The ZnO–rGO sensor exhibited 8% response to 50 ppm NO2 with rather rapid response/recovery times (132 s/164 s) while the pure rGO sensor showed a response of 6.4% with longer response/recovery times (149 s/243 s) at room temperature. The sensor showed exceptional selectivity to NO2. Other gases, such as C3H6O, CH3(CH2)3OH, CH3OH and H2, were not able to react with the oxygen ions that adsorbed on the surface of this composite sensor at low operating temperatures. NO2 is a strongly electron-withdrawing molecule that enhanced the electron-withdrawing ability of oxygen functional groups through its electron-deficient N atom interacting with active sites on rGO. In addition, the authors did an interesting comparative experiment with the ZnO–rGO sensor, a pure rGO sensor and a sensor fabricated by physically mixing ZnO and graphene dispersions. The last sensor exhibited the worst gas-sensing performance to NO2 due to the agglomeration of graphene sheets and ZnO particles. Similarly, this group [2 nanoparticles, synthesized via the method mentioned above, exhibited a higher response to NO2 and faster response/recovery times than a 2D SnO2–graphene sensor, fabricated by the same method without freeze-drying process.ZnO is widely used as a typical wide-bandgap (3.37 eV) metal-oxide gas sensor material. However, the problem with ZnO gas sensors is their poor selectivity . Li et au et al.  developeis group  demonstr2), as a wide-bandgap semiconductor, has been widely used as photocatalyst, and in solar cells and gas sensors [3 sensor using rGO decorated with TiO2 nanocrystals. There were two different morphologies in these sensing materials: rGO either laid on the surface of TiO2 nanoparticles, partly wrapping them, or made “bridges” at the interface between two nanoparticles. Due to the “bridges” existing between TiO2 nanoparticles, the initial resistance of TiO2–rGO sensor was greatly reduced, indicating that the sensor was able to work at room temperature. The partly “wrapping” microstructure enhanced the number of adsorption sites. Moreover, the TiO2–rGO sensor showed a better selectivity to NH3 than the pure rGO sensor owing to the acidic surface of TiO2 preferentially adsorbing primarily NH3. Ye et al. [2–rGO sensor, fabricated via depositing GO and TiO2 nanoparticles on IDEs and then heating them for reduction, showed a 1.5-times higher response than a rGO sensor at room temperature. In contrast, a pure TiO2 sensor did not respond to NH3 at room temperature, which proved that the introduction of graphene reduced the operating temperature of the TiO2 sensor. The same group [3 room-temperature sensor based on TiO2–rGO hybrids, synthesized by a hydrothermal method, exhibited a two-times higher response and a much shorter response time than the TiO2–rGO sensor fabricated by direct deposition as mentioned above. Apart from the porous and undulating graphene sheets due to the introduction of TiO2 nanoparticles, the mellow and regular TiO2 nanoparticles also contributed to the improvement of the gas-sensing performance. However, the stability was poor because the composite sensor was sensitive to water molecules. Further optimization is needed to control the influence of humidity on the sensors.Titanium dioxide  hybrids fabricated via hydrothermal synthesis using GO and SnCl2 as precursors exhibited a 3D nanostructure with high specific surface area (94.9 m2/g). During the hydrothermal process, GO served as a template to promote the preferential growth of SnO2 nanoparticles and prevented SnO2 nanoparticles from agglomeration. The response of the composite sensor to NH3 at 10 ppm was 5.09% and the response/recovery time was less than 1 min at room temperature, whereas the sensor based on SnO2 did not respond to NH3 and the sensor based on GN showed a response of only 2.7 % at room temperature. The authors claimed that the introduction of GN not only increased the specific surface area, but also improved the conductivity of the sensor at room temperature. Bo et al. [2 nanoparticles on the VG networks. The SnO2–VG room-temperature sensor was capable of detecting as low as 20 ppb of formaldehyde and showed a response of 4.6% to 5 ppm formaldehyde, which was three-times higher than that of the graphene sensor. Compared with normal graphene sheets, the specific surface area of vertical graphene sheets is extremely increased. The SnO2 nanoparticles on the VG sheets with the 3D structure provided numerous adsorption sites for target gas molecules.SnOof gases –77, so tn et al.  demonstro et al.  grew ver3)3 to develop In2O3–rGO hybrids through a facile one-step microwave-assisted hydrothermal method. The response of the In2O3–rGO sensor to 5 ppm NO2 was 37.81% with excellent stability and selectivity at room temperature. It should be noted that graphene improved the conductivity of the sensing materials, while the addition of In2O3 nanocubes prevented rGO sheets from re-accumulation, as shown in 2O nanorods–rGO hybrids obtained after annealing showed a porous structure with a high surface area to volume ratio. The porous structure promoted the diffusion of gases, improving the reaction of gases with oxygen species on the surface of the hybrid material. The Cu2O–rGO composites exhibited exceptional catalytic activity and acted as high-efficiency catalysts for the reduction of oxygen molecules, leading to an excellent response. This room-temperature sensor exhibited a linear response to the concentration of NH3 with rapid response/recovery times. The gas-sensing performance parameters of the abovementioned graphene/metal-oxide gas sensors enhanced by improved morphology and structure are listed in Apart from the abovementioned common MOS, some other MOS also have exceptional sensing properties after mixing with graphene. Yang et al.  added a In the past years, some scholars have found that excellent sensing properties can also be achieved when MOS are directly mixed with pristine graphene. The formation of heterojunctions is improper to explain this phenomenon since pristine graphene is a conducting material. Also, the morphology and structure of the composite sensor have not been altered through the introduction of graphene. Another strengthening mechanism, namely a synergetic coupling effect between graphene and metal oxides, is proposed. In detail, the enhanced sensing performance is accounted for by chemical bonds between graphene and metal oxides. Many XPS studies have claimed that there indeed exist chemical bonds between metal oxides and graphene.3, a transition-metal oxide semiconductor is widely used as a gas sensor because of its small bandgap (2.585 eV) and its physical and chemical stability [2 sensor based on WO3 nanospheres wrapped in graphene sheets, prepared by a simple sol–gel technique, showed a linear response to low concentrations of NO2, while the pure WO3 and graphene sensors did not respond to NO2 at room temperature. The authors claimed that the reason for room-temperature sensing of the composite sensor was the effective charge transfer between graphene and WO3 nanospheres by chemical bonds. The research group confirmed that there existed C–O–W chemical bonds between WO3 and graphene by Raman and XPS measurements. The proposed sensing mechanism is shown in 2 molecules, the gas molecules adsorbed on WO3 nanospheres cause the energy band to bend upward via obtaining electrons from WO3 and move the Fermi level of WO3 from the conduction band to the valence band. The reactions mentioned above shift the work function of WO3 [3 and graphene. Because of the continuous loss of electrons of WO3, electrons are transferred from the graphene sheets to WO3 through chemical bonds to maintain the adsorption. The chemical bonds are considered as electrons bridges during the response, which improved the sensing performance of the graphene–WO3 sensor. In addition, the excellent conductivity of graphene may also enhance the response because graphene offers conductive pathways that enhanced the efficiency of charge-carrier transfer in the composites. Zhang et al. [3O4–graphene gas sensor through a traditional hydrothermal method. The XPS results certified that Co–O–C bonds were formed at the interfaces of Co3O4 and graphene. The Co cations in Co3O4 were strongly attracted by the oxygen anions in the oxygen functional groups of the graphene. Consequently, the centers of Co3+ and Co2+ served as additional active adsorption sites for NO2. At the same time, the electrons transferred from graphene to Co3O4 through the Co–O–C bonds lead to an additional increase of the width of hole accumulation layers, leading to a high sensitivity at room temperature. In another work, Feng et al. [3O4 nanocrystals by using electrospinning. The Co3O4–rGO room-temperature sensor showed excellent sensitivity to low concentrations of NH3, an ultrafast response time of only 4 s with exceptional selectivity to NH3. The authors claimed that the introduction of rGO exhibiting a strong attraction to NH3 played a crucial role in improving the sensitivity and selectivity. However, the most important factor was that the chemical bonds (C–O–Co) formed between rGO and Co3O4, which facilitated the electrons transfer at the interfaces of rGO and Co3O4.WOtability –89. In atability  reportedn of WO3  to be adg et al.  fabricatg et al.  develope2 and graphene. The response of the SnO–graphene sensor to 50 ppm NH3 reached 21% with fast response/recovery times of only 15 s/30 s at room temperature. The most encouraging result was that the sensor could maintain its excellent response for more than one year. The authors demonstrated that the formation of C–O–Sn chemical bonds, confirmed by XPS results, was responsible for the superior performance of the SnO–graphene sensor. Zhang et al. [2 than a pure NiO sensor. The authors also ascribed the enhanced sensor sensitivity to the Ni–O–C bonds formed at the interfaces of NiO nanoparticles and graphene sheets. The NO2 molecules capture electrons from the surface of NiO nanoparticles when the sensor is exposed to NO2. More and more electrons are transferred from graphene to NiO by Ni–O–C bonds because of the lower Fermi level of NiO, increasing the sensitivity and response rate of the NiO–graphene sensor. However, the sensitivity of the NiO–graphene sensor was not as good as that of the pure NiO sensor at low concentrations of NO2 (0.25–1 ppm). In order to enhance the sensitivity of this sensor, the authors [2 to the NiO–graphene hybrids (p-type) to form p–n heterojunctions and found that the sensitivity of the SnO2–NiO–graphene sensor was about 10-times that of the NiO–graphene sensor. The authors claimed that the low adsorption energy of NO2 on the surface of SnO2 and the formation of p–n heterojunctions contributed to the exceptional sensitivity. The gas-sensing performance parameters of the abovementioned metal oxide–graphene gas sensors enhanced by a synergetic coupling effect are listed in Kumar and co-workers  fabricatg et al.  fabricat authors  added n-2–rGO fabricated by Zhang et al. [In this review article, the progress of room-temperature graphene/metal-oxide gas sensors has been summarized. The introduction of graphene or its derivatives greatly improves the sensitivity and reduces the operating temperature of MOS gas sensors. The advances in the field of graphene/metal-oxide gas sensors over the past few years have greatly broadened the research directions of gas sensors. However, the achievements reported in this article are still in the stage of basic research and require a lot of effort to put these researches into practical use. The main bottleneck is that the gas sensors are susceptible to environmental humidity 4783,87,96.g et al.  showed aA small number of these composite sensors seem to exhibit the phenomenon of resistance-baseline drift, which is due to the formation of chemical bonds during the adsorption of gas molecules at the material surface, resulting in slow or even impossible recovery. Some researchers have stated that the recovery time can be reduced by heating or irradiating with UV light. Unfortunately, the measures will damage the structure of rGO films and increase the cost. Another point we have to mention is that the sensors based on metal oxides and graphene exhibit complicated sensing mechanisms owing to the additional reactions between gas molecules and various sensing materials.Future works should focus on the following aspects: The most important thing is to simulate the influence of humidity on the sensor performance in a real environment and study the mechanism of humidity influence. A hydrophobic coating can be applied to the surface of the sensor to reduce the effects of humidity. Regarding the synthesis, it is important to quantify and control the effect of hydrothermal or solvothermal parameters on the micromorphology of hybrids. An important way to enhance the selectivity of composite sensors is mixing or doping with other phases . Similarly, the recovery time can be reduced by surface functionalization of graphene/metal-oxide hybrids with specific functional molecules such as boron or nitrogen. In addition, more effort should be dedicated to the detection of more environmental gases, especially VOCs. Last but not least, the influence of baseline drift in resistance can be eliminated by defining a new method, which calculates the response values by using the resistance value after each response cycle.