Curcuma longa L. Modification of the graphene’s surface with these compounds also increased the thermal and chemical stability of PVC/graphene nanocomposites.A large amount of graphene-related research is its use as a filler for polymer composites, including thin nanocomposite films. However, its use is limited by the need for large-scale methods to obtain high–quality filler, as well as its poor dispersion in the polymer matrix. This work presents polymer thin-film composites based on poly(vinyl chloride) (PVC) and graphene, whose surfaces were modified by curcuminoids. TGA, UV–vis, Raman spectroscopy, XPS, TEM, and SEM methods have confirmed the effectiveness of the graphene modification due to π–π interactions. The dispersion of graphene in the PVC solution was investigated by the turbidimetric method. SEM, AFM, and Raman spectroscopy methods evaluated the thin-film composite’s structure. The research showed significant improvements in terms of graphene’s dispersion (in solutions and PVC composites) following the application of curcuminoids. The best results were obtained for materials modified with compounds obtained from the extraction of the rhizome of  Research concerning the application of graphene materials is carried out in many areas, such as medicine, chemistry, materials engineering, energetics, and electronics. A large area of graphene–related research regards its use as a filler for polymer nanocomposites ,2,3,4,5.g) [An interesting and still relatively unexplored group of polymer nanocomposites with graphene are poly(vinyl chloride)–based materials (PVC/GN). Analysis of the literature has shown that in these materials with relatively well-dispersed graphene, regardless of the method of preparation , result in an increase in tensile strength and modulus of elasticity ,14,15,16g) ,24, loweg) ,26. HoweThe agglomerative behavior of graphene is due to strong π–π interactions and van der Waals forces ,28,29,30The π–π interactions refer to non–covalent interactions  between the π bonds of aromatic rings. However, this is a misleading description, because the direct sandwich overlap is electrostatically repulsive. Instead, we are dealing with staggered stacking  or π–teeing (perpendicular T–shaped), which makes the described interactions electrostatically attractive. The displacement occurs due to the superposition of carbon atoms with partially negative charges over hydrogen atoms with partially positive charges. Similarly, T–type perpendicular overlap is electrostatically attractive because here, too, there is an overlap of different charges ,33,46.A promising method of graphene modification using π–π stacking is the modification of its surface with curcuminoids. The interaction between the aromatic rings of curcumin and the GN surface was used to obtain graphene by reducing graphene oxide, through ultrasonic exfoliation of graphite ,34,47. HCurcuma longa L. (CE) obtained by extraction of the powdered rhizome of the plant, which was presented in our earlier work [Therefore, the present work aimed to modify surface graphene with curcuminoids and to investigate the potential of this modification to improve graphene dispersion in poly(vinyl chloride)/graphene nanocomposites obtained by the solvent-casting method. In this work, we used a commercially available mixture of three types of curcumin (CU)  and an extract of ier work . In addiCurcuma longa L. extract, and GN–CU and GN–CE colloidal suspensions. The analysis showed that the maximum absorption for graphene was observed at 272 nm, corresponding to π–π* transitions of C=C bonds [Curcuma extract, the described transitions were observed at 233 nm and 422 nm, respectively. The shift of the absorption maximum (233 nm) toward the ultraviolet spectrum is due to the presence of the other CE components. A maximum wavelength of 233 nm is characteristic of ar–turmerone [Before the morphology studies were conducted, the UV-vis studies were performed, in order to determine if the GN was successfully modified with curcuminoids. =C bonds ,54. Two =C bonds . In the urmerone ,57. Two urmerone . In addiurmerone ,55,59.2 hybridization was expected. Unfortunately, it was impossible to match sp2 in any way, so an additional sp3 hybridization peak was used. This fact may be due to the partial defecting of the graphene used and its partial oxidation  [Next, X–ray photoelectron spectroscopy (XPS) was used to characterize graphene surfaces’ functionalization accurately. xidized) ,47,60.2 and C sp3 peaks at 0.8 eV, especially for the starting sample, agrees very well with literature data on peak distances of thin carbon films, which are, respectively, 0.7 eV for graphene oxide and reduced graphene oxide [2 peak around 284.5 eV is also characteristic of this type of material [The distance between the C spne oxide , 0.8 eV ne oxide ,63, and ne oxide ,65,66. Fmaterial  and genematerial ,61,67.2 [The rest of the smaller carbon peaks , Table 12 . Conside2 ,70,71. T2 ,72.Curcuma longa L was stable up to about 175 °C, and its degradation occurred between 175 °C and about 450 °C  can be seen in the SEM images B,C 83,8,883,84. −1), which indicates the residual compression stress in the material. The RBS parameter (Raman Band Separtion) is relatively high and equals 287 (GN–CE) to 293 (GN–CU), whereas non modified filler has a value of ~291. This confirms the different mechanisms of carbon structure modification by both curcuminoids.The following step was the evaluation of the structure of graphene materials using Raman spectroscopy, an effective method for studying carbon materials . Recorde−1 for GN to 1600 cm−1 for graphene modified with curcumin and Curcuma longa L. extract, further confirming the better foliation of the material after surface functionalization with curcuminoids. Bands typical for disordered carbons are not visible .In paper , it was As we pointed out in our earlier work , in nanoThe analysis shows that the destabilization and stratification of graphene dispersion occur due to re-agglomeration and sedimentation of the filler ,89. The The Turbiscan Stability Index (TSI) (a specific Turbiscan parameter) was determined to compare the destabilization kinetics of the prepared systems, in particular, the differences between the CU and CE effect. TSI is an efficient method of comparing several systems, where the higher its value is, the less stable the system is ,92. OtheThe structure of nanocomposite polymer thin films obtained by solvent casting was studied using Raman spectroscopy and atomic force microscopy. In addition, the morphologies of the obtained materials were also observed on a macro- and micrometer scale, using images from digital photography and SEM. As mentioned in the introduction, the problem with graphene aggregation occurs in materials containing about 0.5 wt.% filler or more, so the structure evaluation was carried out for materials with 1 wt.% GN.−1 to C–Cl in the trans configuration of the polymer [−1 to C–Cl stretching [−1 to C–C stretching [−1 to C–H symmetrical stretching in CH2 group [−1 to C–H asymmetrical stretching in CH2 group [−1. It was attributed to the C–O bonds of curcuminoids [−1 , being the lowest in GN–CE and the highest in GN–CU. This can be explained by the layer-stacking enhancement in the matrix’s presence. Thus, the best filler dispersion was observed in GN–CU. Contrarily, the D-bands have been shifted towards higher Rs . The D (first–order peak) moves towards lower frequencies in the case of tensile stress and towards higher ones in the presence of compressive stress, as observed in nanocomposites. Additionally, the ID to IG ratio decreases, especially in the case of the GN–CU, indicating better structural ordering. However, the sheer increase in the intensity of the PVC-derived band (2916 cm−1) in the materials after its functionalization may indicate better dispersion of these materials and deposition inside the polymer matrix [Regarding the Raman spectra of the nanocomposites , D– and  polymer , bands aretching ,94, a baretching , a band H2 group  and a baH2 group ,93. In auminoids . The stuAFM also investigated the surface structure of the obtained nanocomposites in contact with a surface . AFM allows for the direct evaluation of the dispersion of graphene in the PVC matrix and the determination of roughness ,96. FiguThe image of the surface topography of the nanocomposites shows some corrugation of the samples due to the evaporation of solvent during the formation of the polymer film . AFM PVCThe observations made earlier were confirmed by analyzing the homogeneity of the polymer films on a macroscopic scale. The presented digital images of nanocomposites containing 1 wt.% GN, GN–CU, and GN–CE The same observations were confirmed by SEM imaging on the surface of the PVC/1%GN cryogenic breakthrough /graphene nanocomposites. Materials containing 0.01, 0.1, 0.5, and 1 wt% GN were produced.Observations of the structure demonstrated the best improvement in the dispersion of graphene in materials modified by Thermal properties were determined by the thermogravimetric method and the Congo Red test, with which the thermal stability time was measured. The results obtained are summarized in Thermal degradation of PVC and PVC/GN nanocomposites takes place in two stages related to the decomposition of the polymer. The first step occurs at 200 °C to 375 °C, and is related to the dehydrochlorination of PVC and the formation of conjugated polyene structures. The second step in the range from 375 °C to 600 °C corresponds to thermal cracking of the carbonaceous, conjugated polyene sequences and the formation of residual chars ,23.Due to the presence of residual THF in PVC, the nanocomposite films’ thermal stability was evaluated based on the differential thermogravimetry (DTG) ,19. DTG The max. DTG II analysis  showed nA commonly used method for assessing the thermal stability of PVC materials is the Congo Red test. The results presented in At the same time, it should be emphasized that discrepancies in the results of thermal stability measured by two methods are not unusual, due to the conditions under which the tests were conducted. Thermal stability by the Congo Red method is determined at a constant temperature that is close to the upper temperature range of PVC processing. In the TGA method, the sample is heated under a nitrogen atmosphere at a specified heating rate, over a wide temperature range, usually up to about 900 °C.An improvement of the mechanical properties of PVC/GN nanocomposites depends on the homogeneous dispersion of graphene in the polymer matrix and the structure of the filler. The lack of defects on the GN surface results in its high integrity and good mechanical properties, which directly determines the improvement of the mechanical properties of nanocomposites ,99,100. 14 Ω, and significant change occurs when the GN concentration in the composites is 0.5% and 1%. The lowest resistivity, which equals 1.4 × 107 Ω, is characteristic of the PVC/1%GN sample, which makes it possible to conclude that it is an antistatic material [14 Ω) was shown by the sample PVC/1%GN–CE.material ,101. Nan5 Ωm. Such a large change in resistivity (for PVC was 1.7 × 1014 Ωm) indicates that the dispersion of the GN was sufficient to create conduction paths. The PVC/GN–CE composites, as in the case of surface resistivity, did not show large changes in volume resistivity as the filler content in the matrix increased.The volume resistivity of PVC/GN composites also decreases with the increase in graphene concentration in the matrix. The materials containing 1% of GN were characterized by the lowest resistivity, and equaled 9.8 × 10References ,101 presThe numerous applications of poly(vinyl chloride) are due to its resistance to solvents ,102,103.dS is swelling degree, %,ES is equilibrium swelling, upper asymptote, %,Mt is time in which the swelling occurs with a maximum rate, s,t is time of exposure to the swelling agent, s,p is comparison parameter, 1 s−1.It was found that all the obtained materials undergo limited swelling, while the dependence of swelling degree on the exposure time to the swelling agent has a shape of a sigmoid function. Therefore, Equation (1) was used to approximate the swelling curves ,22.(1)SR2 are summarized in The parameters of the equation and the coefficient of determination M) also increases, which further confirms the correlation between acetone resistance of the materials and filler surface modifications. In materials with unmodified GN, a decrease in tM values was observed with an increase in its content in the PVC matrix.PVC/GN nanocomposites were found to have a lower equilibrium swelling ratio compared to unmodified PVC; the lowest SE value observed for PVC/1%GN, i.e., 43.5%, shows an improvement of about 5% in terms of the chemical resistance on the polymer matrix. Materials obtained with the addition of extract-modified graphene were characterized by a lower value of SE compared to corresponding materials with unmodified graphene. This effect was attributed to the better dispersion of the filler in the polymer matrix, as evidenced by significantly lower SE values for materials with 0.5 and 1% GN–CE. They were 32.1 and 27.3%, respectively, which represents a reduction in the equilibrium degree of swelling relative to PVC of 14.1%, in the case of PVC/0.5%GN–CE, and of 20.9%, in PVC/1%GN–CE. In addition, it should be noted that the agglomeration of the filler in the polymer matrix can lead to lower resistance to solvents in the form of acetone, which was observed in PVC/0.5%GN, where the SE was 1.3% higher relative to PVC. As the GN–CE content of the nanocomposites increases, the time in which the swelling occurs with a maximum rate  solution using unmodified suspensive poly(vinyl chloride) Neralit 601 , with a K number of 59–61, a bulk density of 0.56–0.63 g cm−3, a specific density of 1.39 g cm−3, and 97% purity. As a solvent for PVC and an environment used for graphene modification, tetrahydrofuran (THF)  was used. Curcumin from Curcuma longa L. turmeric powder  and Curcuma longa L. rhizome extract (made in-house [Graphene–based nanopowder with a flake thickness of 1.6 nm (maximum of 3 atomic monolayers), a flake length of 10 µm, and a specific surface area of 400 ÷ 800 min-house ) was use3 solutions of CU and CE in THF with a concentration of 2.25 mg cm−3 were prepared by ultrasound  using a SONOPULS rod-shaped probe homogenizer from Bandelin. Graphene was then introduced into the solution, which was prepared so that the dispersion concentration was 1.5 mg cm−3. Next, GN was sonicated under the same conditions as curcuminoids for 2 h. It was then stirred with a magnetic stirrer for 15 h at 200 rpm. After this time, the material was centrifuged 4 times at 12,000 rpm for 4 min and washed after each step with methanol, in order to remove the undeposited modifier. After centrifugation, the graphene was dried at 45 °C for 65 h to evaporate residual solvents.In order to functionalize the graphene’s surface with curcuminoids in the first step, 100 cmIn the first step, poly(vinyl chloride) was dissolved in THF at 25 °C for about 48 h, yielding a solution of 3 wt%. Then, the dry graphene was added to the solution and dispersed using ultrasound  at 50 °C for 2 weeks, in order to remove the THF residue from the polymer films. All samples were coded to take into account the graphene content and the presence of CE, e.g., a sample containing GN at a concentration of 0.01 wt.% was coded as PVC/1%GN, whereas a sample containing the same graphene modified by −3 and dispersions of GN, GN–CU, and GN–CE with a concentration of 0.03 mg cm−3 were prepared. The study was conducted in the 210–800 nm range. The experimental results were normalized in the range 0–1 using the OriginLab software.Absorption studies were conducted using a UV-vis/NIR spectrometer (Perkin Elmer Lambda 1050+) in a quartz cuvette. For this purpose, solutions of CU, CE in methanol with a concentration of 0.008 mg cmX–ray photoelectron spectroscopy: a PHI 5000 VersaProbe  spectrometer was used to conduct XPS measurements under the following conditions: monochromatic Al Kα radiation (hν = 1486.6 eV), an X-ray source operating at 25 W, 15 kV, 100 µm spot, pass energy of 23.5 eV, energy step of 0.1 eV. Obtained XPS spectra were analyzed with CasaXPS software, using the set of the sensitivity factors native for the hardware. Shirley background and Gaussian–Lorentzian peak shape were used for deconvolution of all spectra.−1 in an open ceramic crucible, under a nitrogen atmosphere, at 30 to 700 °C for fillers and at 30 to 900 °C for nanocomposites. In addition, static thermal stability tests were determined for the nanocomposite materials using the Congo Red test. The test was carried out at 200 °C, measuring the thermal stability time, i.e., the time during which the sample shows no signs of degradation in the form of hydrogen chloride release and color change of the indicator paper.The thermal stability of the fillers and nanocomposites was assessed by the thermogravimetric method using a TG 209 F3 Tarsus apparatus (Netzsch). The heating rate was about 10 °C minRaman spectra were collected with a Raman DXR microscope  using two laser lines, 532 nm and 780 nm, where the shorter wavelength was used for polymer composites, and fillers were studied using the longer wavelength line. The specimens were measured with 50 repetitions of 50 s each (20 in successive control measurements). The aperture was 50 µm, the lenses were 10 and 50 mm, and the laser beam power was 0.1 mW for 780 nm and 1 mW for 532 nm. Each sample was measured at least dozens of times (from 80 to >300) at different locations, and the final signal was obtained as an average and statistically representative one. The morphology of the samples was studied using scanning electron microscopy (SEM), and a Zeiss Crossbeam 350 microscope; in the case of polymer nanocomposites, cryogenic breakthroughs with a sputtered gold layer were observed. As a second technique, a Zeiss Libra 120 Plus transmission electron microscope (TEM)  operating at 120 kV was used, where the aqueous suspension was placed on a copper grid coated with a formvar polymer layer. Finally, the sample was left in the air (under a fume hood) to dry.Stability testing of the obtained graphene dispersions in PVC solution was carried out by multiple light scattering using the Turbiscan Lab apparatus , where suspensions of GN, GN–CU and GN–CE in 3% PVC solution in THF were placed in glass cylindrical test tubes, with a working height of 54 mm. The samples were then placed in a Turbiscan Lab instrument and scanned with light at approximately 880 nm. Scanning was carried out every 1 h for 24 h at room temperature (23 °C). The Turbiscan Stability Index (TSI) was determined from the results.−1.Filler dispersion in the matrix was also determined using an atomic force microscope Nanosurf  in contact mode with the PPP–XYCONTR probe , where the radius of the aluminum-coated blade was 7 nm, while the line angle was 20 degrees. 2D and 3D images were taken within a 50 × 50 µm area. To determine the surface topography of the area, 500 lines were measured , and the number of measurement points per line was 5000. The scanning speed of the sample was 5 µm s−1.A study of tensile properties was carried out using an in-house-built tensile tester, with the parameters presented in our earlier work . The tesThe electrical properties of nanocomposites were determined by measuring system consisting of a 6517A electrometer and a 8009 measuring chamber . The surface and volume resistivity were measured. The tests were carried out on film samples with a diameter of 70 mm in air, at a temperature of 23 °C and humidity of 50%, and at a voltage of 10 V.dS) was determined (see Equation (2)), depending on the immersion time in the swelling agent. The changes in the samples’ diameters were set on the basis of photos using the NIS Elements 4.0 software. The frequency picture-taking depended on the exposure time to the swelling agent. The measurement temperature was 20 °C and the initial diameter of the samples h0 = 10 mm.h is sample diameter after time t (mm),h0 is initial sample diameter (mm).The resistance of the obtained materials to swelling in acetone was tested in accordance with the method proposed in our earlier study . The chaIn order to analyse the obtained results, Origin 8.6 Pro software with implemented statistical analysis modules was used. ANOVA with Tukey’s post–hoc test was used to compare the significance of the difference for the mean values of the obtained results. The normal distribution was confirmed by the Shapiro–Wilk test, and the homogeneity of variance by the Levene’s test. All analyses were performed assuming a significance level below 0.05.Curcuma longa L.Graphene is one of the most promising nanofillers of nanocomposites, improving their performance. However, its uniform distribution in the polymer matrix is challenging. Several studies refer the need to reduce the agglomeration of GN and its derivatives during nanocomposite fabrication, in order to achieve improved properties ,106,107.Observations of the modified materials’ morphology showed better CE dispersion on the graphene’s surface. Spectroscopic studies have shown that modification occurs via a π–π stacking mechanism. TGA analysis revealed a higher amount of deposited curcumin on GN flakes i.e., 11.4 wt.%, while for the extract it was 6.9 wt.%. In addition, it is worth noting that the modified materials have a high thermal stability of about 340 °C. The analysis of the stability of GN, GN–CU, and GN–CE dispersions, as well as observations of the polymer films structure obtained after solvent evaporation from the dispersions, showed a significant improvement in the homogeneity of the nanocomposites after modification by curcuminoids. At the same time, extract-modified graphene showed better dispersion improvement in PVC. That indicates that not only the amount of modifier deposited on the graphene surface, but also its size and dispersion important in improving PVC/GN homogeneity. The analysis of the properties of nanocomposites with GN-CE has shown a better thermal and chemical stability of these materials. On the other hand, based on the results of the Congo Red test, it was stated that the proposed nanocomposites will require thermal stabilization at the processing stage by traditional methods, which is a well-known phenomenon in the field of the PVC technology.Summarizing, the proposed naturally derived compounds in the form of CE can be an alternative to previously proposed modifications of graphene for use in PVC and other polymer composites. In addition, the high thermal stability of the proposed modification, as well as the relatively low cost and possibility of obtaining the modifier on a large scale, provide opportunities for the technological use of CE in the future production of graphene polymer nanocomposites, through traditional processing methods at high temperatures.


Experiments that can resolve gas flows of a few atoms per hour through micrometre-sized membranes found that monocrystalline graphene is completely impermeable to helium, the smallest atom5. Such membranes were also shown to be impermeable to all ions, including the smallest one, lithium7. By contrast, graphene was reported to be highly permeable to protons, nuclei of hydrogen atoms9. There is no consensus, however, either on the mechanism behind the unexpectedly high proton permeability14 or even on whether it requires defects in graphene’s crystal lattice17. Here, using high-resolution scanning electrochemical cell microscopy, we show that, although proton permeation through mechanically exfoliated monolayers of graphene and hexagonal boron nitride cannot be attributed to any structural defects, nanoscale non-flatness of two-dimensional membranes greatly facilitates proton transport. The spatial distribution of proton currents visualized by scanning electrochemical cell microscopy reveals marked inhomogeneities that are strongly correlated with nanoscale wrinkles and other features where strain is accumulated. Our results highlight nanoscale morphology as an important parameter enabling proton transport through two-dimensional crystals, mostly considered and modelled as flat, and indicate that strain and curvature can be used as additional degrees of freedom to control the proton permeability of two-dimensional materials.Defect-free graphene is impermeable to all atoms A study using high-resolution scanning electrochemical cell microscopy attributes proton permeation through defect-free graphene and hexagonal boron nitride to transport across areas of the structure that are under strain. Additional experiments with hydrogen’s heavier isotope deuterium revealed that the initial energy of incoming protons is not given by thermal excitations (about 25 meV) but is instead about 0.2 eV owing to zero-point oscillations of protons bound to oxygen atoms in the proton-conductive media9. This correction lifts the total energy barriers, E, posed by the crystals to about 1.0 eV and about 0.5 eV for graphene and hBN, respectively. Despite these insights, the mechanism for proton permeation through the 2D crystals remains controversial. The general consensus from density functional theory calculations is that the energy barriers should be notably larger14. The studies  have yielded a rather wide range of E but always exceeding the value of about 1 eV found experimentally. The spread of values arises from the various assumptions made in the models, such as whether the process is slower than the lattice relaxation timescale14, protons tunnel quantum mechanically12 or protons locally hydrogenate the carbon lattice  before transfer19. This uncertainty has motivated an alternative explanation widely speculated in the literature, namely that proton permeation takes place through structural defects in the crystal lattice. This hypothesis is based on experiments using graphene grown by chemical vapour deposition (CVD)17, which has grain boundaries, pinholes and other imperfections that appear during growth and transfer22. Experiments using CVD graphene typically report very high proton permeation rates and, sometimes, even the loss of graphene’s impermeability to other ions16. However, the explanation that assumes atomic-scale defects as the only proton conductive sites is inapplicable to mechanically exfoliated graphene. Indeed, transmission and tunnelling electron microscopy have failed to observe any vacancies or other atomic-scale imperfections for scans over relatively large areas of such crystals. Even more decisively, gas permeation experiments that can easily detect a single-angstrom-scale defect permeable to gases within micrometre-sized membranes5 detected none in exfoliated graphene and hBN monolayers6. Further experimental evidence is necessary to understand proton transport through defect-free 2D crystals and resolve the existing controversy.Measurements of proton transport through two-dimensional (2D) crystals demonstrated that these crystals pose an energy barrier for incoming protons of about 0.8 eV and about 0.3 eV for graphene and hexagonal boron nitride (hBN), respectivelyx) substrates (2 (see the section in the Eapp and Ebias set so as to fix Ecollector = −0.5 V at the Pt electrode (H+ collector) with respect to Ag/AgCl (see the section in the + pump) and the Nafion–Pt collector, and a current (Icollector) is detected only when the probe is at locations where H+ transmission occurs  and high current (fA) resolution using scanning electrochemical cell microscopy (SECCM). The devices for this study consisted of graphene and hBN monolayer crystals, which were suspended over micrometre-sized holes (2 μm in diameter) etched into silicon nitride  currents of about 10 fA are observed because the SiNx substrate blocks proton transport , each exhibiting a proton current17 of about 0.3–1 nA under similar Ebias. Our present SECCM measurements using higher-quality CVD graphene have not found such isolated highly conductive defects (possibly pinholes) and instead show enhanced permeation over large areas, mostly at what seems to be grain boundaries, a common feature of CVD-grown graphene20  images of the 2D crystals. Figure x aperture is covered with monolayer hBN and the other half with tetralayer  in the SECCM map is concentrated around wrinkles and the rim. However, the maps for hBN monolayers also reveal many active areas that do not correspond to any obvious morphological features. In the statistical distribution  or static (caused by strain or adsorbates), as previously revealed by transmission electron microscopy26. Although our AFM maps do show a difference in apparent roughness for monolayer and tetralayer regions in Fig. L of several nanometres26) to be quantified using our AFM. Nevertheless, direct evidence from transmission electron microscopy has previously found26 that static ripples have typical h/L ≈ 0.1, comparable to the aspect ratios observed for the wrinkles. On this basis, we propose that unavoidable nanorippling of 2D crystals enhances their proton permeability in much the same way as larger wrinkles. As nanoripples have smaller L, their smaller areas result in smaller proton currents within individual pixels. To corroborate this microscopic picture provided by SECCM, we integrated the observed currents over the entire area of the 2D membranes. This allows for an estimate of the proton conductivities of graphene and hBN monolayers as about 10 mS cm−2 and about 300 mS cm−2, respectively. Taking into account our experimental uncertainty (factor of about three that comes mostly from assessing the area contacted by the probe during the scans), the found conductivities are in good agreement with the previous measurements of average proton conductivities of the 2D crystals8.We attribute the smaller proton currents away from apparent morphological features to transport through nanoscale ripples that are ubiquitous in 2D crystals8. For example, monolayer hBN presents a sparser electron density than graphene and, accordingly, is more permeable to protons8. Strain and curvature modify the electron-density distribution within 2D materials  seem to be consistent with the experiments reporting barriers of about 1.0 eV (ref. 8). Finally, strain and curvature can be exploited to enhance proton conductivity of 2D crystals, which is of interest for various applications involving proton transport30.In conclusion, our experiments show that strain-inducing morphological features in otherwise defect-free 2D crystals are associated with enhanced proton conductivity around them. A notable example of this is graphene wrinkles that do not require any crystal-lattice defects but result in high proton currents, not dissimilar to the case of grain boundaries in CVD graphene . Our finx) using photolithography, wet etching and reactive ion etching, following the protocol previously reported8. Our devices had several apertures next to each other and were 2 μm in diameter each  to obtain ≈10-μm-thick films. The devices were annealed in a water-saturated environment at 130 °C to crosslink the polymer. The collector electrode was prepared by laminating a Pt foil  onto a cylindrical carbon block with a hot compression mounting machine (SimpliMet). The exposed Pt surface was then subjected to mechanical and electrochemical polishing. Gaskets were used to cover the carbon block and expose only the Pt surface for contact with Nafion–2D crystal devices34. For measurements, the Nafion film was hydrated with deionized water and allowed to equilibrate before placing it in contact with the Pt collector.Micrometre-sized apertures were etched into silicon-nitride-coated silicon substrates (500 nm SiN2-laser puller (Sutter Instruments P-2000). The nanopipette was then filled with 100 mM HCl electrolyte, and a silicone oil layer was added on top of the electrolyte solution in the tip to minimize evaporation during prolonged scanning procedures35. Two AgCl-coated Ag wires, fabricated by electrochemically oxidizing Ag wires (0.125 mm in diameter) in saturated KCl solution36, were used as quasi-reference counter electrodes (QRCEs). Each of the nanopipette channels was fitted with a QRCE positioned about 3–4 cm away from the tip end36.Nanopipettes for SECCM were fabricated from quartz theta capillaries with filaments . The capillary  was pulled to form a fine sharp point with a tip opening diameter of about 200 nm, using a CO37. The SECCM probe was mounted on a z piezoelectric positioner  whereas the studied graphene or hBN device was mounted on an x–y piezoelectric positioner . The SECCM probe was moved to the initial scan position using an X–Y-micropositioner  controlled with Picomotor Actuators . An optical camera provided a visual guide for the probe positioning. The microscopy stage and all positioners were enclosed in a Faraday cage with heat sinks and vacuum panels to minimize noise and thermal drift. The Faraday cage was placed on an optical tabletop with tuned damping  balanced on a high-performance laminar flow isolator .SECCM was carried out using a home-built workstationhttps://www.warwick.ac.uk/electrochemistry/wec-spm). Two home-built electrometers were used for current measurements, together with home-built eighth-order brick-wall filters with the time constant of the current amplifier set to 10 ms. Influence of acquisition parameters on the noise level is detailed elsewhere38. The current data were acquired every 4 μs, and 256 samples were averaged to give a data acquisition rate of about 1 ms.Data acquisition and instrumental control were carried out using an FPGA card (PCIe-7852R) running the Warwick Electrochemical Scanning Probe Microscopy software  traces of the reduction of protons at the bottom Pt collector electrode  (ref. 42). This potential is chosen as Ecollector = −0.5 V versus Ag/AgCl QRCE . This was the maximum voltage used in ref. 8 in which the linear-response proton transport was studied in similar devices. Ecollector drives the electrochemical proton reduction at the Pt electrode, which results in current Icollector . SECCM topographical maps of mechanically exfoliated graphene devices (not shown) were similar to those subsequently obtained by AFM, as expected. Such maps, recorded synchronously with proton transport activity, were especially valuable as they revealed morphology around proton-conducting sites in CVD graphene ) and Idc = Idc = Idc first sharply increasing and then returning to the initial value. These steps were clearly visible throughout scanning of entire samples.Extended Data Fig. Icollector . Extended Data Fig. Idc exhibits the same features both in areas of high proton conductivity (blue curve) and in areas where no proton transport takes place (red). This shows that meniscus wetting of the sample is independent of proton transport through 2D crystals. Note, however, that the magnitude of Idc does change for active and inactive areas because the Ag/AgCl electrodes are also the counter electrodes for proton conductivity measurements42. This change served as independent confirmation of those sites where there was notable proton permeation through 2D crystals.The described behaviour was observed independently of 44.Note that the above also rules out changes in the droplet cell size as the source of the observed spatial inhomogeneity of the proton currents. The consistency of the SECCM cell size across the surface is also in accord with the following considerations. First, wrinkles protrude at most a few nanometres from flat areas of graphene, which leads to only small variations in the involved surface area as compared to the area probed in each pixel (about 200 nm in diameter). Therefore, this cannot explain the orders of magnitude difference in the observed SECCM activity. Second, the roughness associated with the wrinkles is much less than a typical surface roughness of a wide range of samples previously studied by SECCM for which consistent meniscus cell size was observed or deducedk = 0.4 N m−1, resonant frequency of 70 kHz and a tip radius of 2 nm. The resulting AFM maps were used to estimate strain across different areas of the 2D membranes. From AFM traces through the membrane centre, we estimate that the membranes were globally strained by typically 0.5%. However, the strain ε was distributed not uniformly but accumulated around the aperture rim23, leading to ε several times higher than away from it23. This yields ε of a few percent around the rim. Such strain is also expected to accumulate around wrinkles in the 2D membranes, whose complex morphology cannot be attained using strain-free  deformations26. From the height (h) and base (L) of the wrinkles measured in AFM, we also estimate strain of a few percent, consistent with the above expectations.High-resolution topography and adhesion AFM imaging was carried out under ambient conditions with a Bruker Dimension Icon AFM using the PeakForce mode. The instrument was equipped with SCANASYST-AIR silicon tips (Bruker). The tips had a nominal spring constant, For scanning electron microscopy (SEM) characterization, we used a Zeiss Gemini500 scanning electron microscope, using an In-lens secondary electron detector, accelerating voltages of 0.5–2 keV and a working distance of about 2 mm.8. The images of Extended Data Fig. x substrate beyond individual apertures. These wrinkles exhibit notable proton-conducting activity in the SECCM maps that occurs not only above the Nafion region but extends onto the SiNx substrate. We attribute this observation to water that fills the space between the substrate and wrinkles and thus provides a proton-conducting medium inside the wrinkles.Extended Data Fig. 46 as well as gas permeation measurements5. None of those studies could detect any structural defects in the membranes. Nevertheless, it was important to ensure that the fabrication procedures used in the present report did not lead to accidental tears, cracks or pinholes that would break the continuity of the graphene lattice and leak protons through.Suspended membranes made from exfoliated 2D crystals have previously been characterized extensively using AFM, SEM, Raman spectroscopy, transmission electron microscopy and scanning tunnelling microscopy47, and graphene is no exception. To understand the formation of wrinkles in our devices we note that graphene sheets are initially suspended over holes, rather than supported. The membranes are therefore stretched laterally because of adhesion to the holes’ sidewalls and free to relax in the out-of-plane direction. In most cases, this results in wrinkle-free membranes2. The situation changes after depositing Nafion. Adhesion to sidewalls disappears in the presence of water (as observed in ref. 48) so that graphene is no longer stretched over the holes. The membrane therefore becomes looser, which unavoidably results in the formation of wrinkles. In addition, the now loose graphene sheets conform to the porous Nafion polymer surface, which further contributes to the wrinkling and rippling. Importantly, the wrinkles and roughness do not lead to cracks, tears or pinholes that would allow unimpeded proton permeation through them. This conclusion is supported by many experimental observations. For the sake of brevity, we describe below only three of them.The formation of wrinkles in supported thin sheets is a universal phenomenon that arises from non-uniform adhesion between the sheet and the substrate. For example, this phenomenon has been extensively studied for 2D polymers49). Our graphene monocrystals do not exhibit any discernible D peak, which allows us to put an upper bound on the atomic-scale defect density of about 109 cm−2. This translates into no more than 10 single-atom vacancies for our entire membranes of 2 μm in diameter . By contrast, the reported wrinkles are hundreds of nanometres long, and if there were any breakdown of crystallinity along them, an intense D peak would also be apparent. Occasionally, we found devices with accidental cracks formed during fabrication, and those exhibited a strong D peak. They were discarded. All of the devices reported in the manuscript had no discernible D peak  activate the so-called D peak in its Raman spectrum. The intensity of this peak increases with defect density . These experiments have found similar proton conductivity as in devices measured using Nafion. Unfortunately, we cannot remove Nafion after measurements, but we could remove electrolytes. In the latter case, the membranes did not show any D peak or any damage under AFM or SEM, which demonstrates that the membranes were not damaged during proton conductivity measurements. As the conductivity using electrolytes is the same as in the case of Nafion, we can safely conclude that Nafion does not damage graphene membranes either.The second piece of evidence that rules out lattice defects in our membranes comes from measurements using liquid electrolytes  exhibited notable helium leakage. This was measured using a He-leak detector that allowed us to resolve flows as low as 1050. To this end, the Cu foil that was covered on both sides with graphene was first exposed from one side to oxygen plasma, which removed graphene from that side. The CVD graphene remaining on the other side was then hot-pressed against the Nafion film, and the Cu foil was etched away in an ammonium persulfate solution. The resulting graphene-on-Nafion stack was left in deionized water for days to remove etchant residues. For SECCM measurements, centimetre-sized graphene-on-Nafion samples were fixed17 to the Pt electrode (as described above) and characterized using the same procedures as for micrometre-sized 2D crystals.For these measurements, centimetre-scale pieces of CVD graphene (grown on Cu) were transferred onto a Nafion N212 film as reported previously17. Statistics of the currents collected over large areas . We have also found that grain boundaries in CVD graphene were often accompanied by local corrugations .Extended Data Fig. 51. Even in the absence of gross defects , which sometimes are prevalent in CVD graphene films17, nanoscale pinholes can result in isolated hotspots with proton currents17 up to 1 nA. For higher-quality CVD graphene without such defects, proton conductivity is likely to be dominated by grain boundaries. Even in the last case, considerable variability in proton permeability is expected because of different grain sizes, depending on growth conditions, so that graphene films with smaller grains and thus higher density of grain boundaries would exhibit higher proton conductivity, in agreement with the previous report7.The experiments described in this section provide important insights into the large variability in proton permeability reported in the literature for CVD graphene films52 implemented in the Vienna ab-initio Simulation Package53 to model pseudopotentials of protons, and C and H atoms. The exchange-correlation potential was taken into account by considering the generalized gradient approximation within the Perdew–Burke–Ernzerhof form54. The weak van der Waals forces between graphene and proton were also included by using the DFT-D2 method of ref. 55. For geometry optimizations, a kinetic energy cutoff of 500 eV was used for the plane-wave basis. The convergence criterion of the total force on each atom was reduced to 10−5 eV Å−1 and the convergence criterion for the energy was set at 10−6 eV. For calculating the proton barrier, we used the proton pseudopotential from the hydrogen atom and then removed an electron from the whole system.We used the projector augmented-wave methodh , which was sufficient to prevent proton–proton interactions between neighbouring supercells. The carbon cells were isolated with a vacuum gap larger than 10 Å, which ensured the absence of edge-to-edge interactions. The ripples were modelled by fixing the out-of-plane positions of carbon atoms so that the crystal forms a Gaussian profile of height E for proton permeation is calculated by subtracting the minimum energy from the maximum one. For the case of flat unstrained graphene, the energy barrier found using these approximations is about 1.37 eV, in good agreement with the earlier theory14. As various approaches used to calculate E yield a rather large spread in the predicted values18 and the exact value of the energy barrier for flat monolayer graphene remains debatable14, here we avoid this uncertainty by focusing on relative changes in E that are arising from strain and curvature. Finally, note that E is expected to vary across membranes becoming lower around wrinkles and ripples and higher in flatter and unstrained areas. As this strain is mostly random, it is reasonable to expect that in the first approximation the distribution of E is normal . As proton currents depend exponentially on E, their distribution should then be log-normal, which is consistent with our SECCM observations.We calculated the total energy of the proton–graphene system as a function of the position of the proton in the direction perpendicular to the centre of the hexagonal ring in the graphene crystal lattice (Extended Data Fig. Any methods, additional references, Nature Portfolio reporting summaries, source data, extended data, supplementary information, acknowledgements, peer review information; details of author contributions and competing interests; and statements of data and code availability are available at 10.1038/s41586-023-06247-6.


This Topic on the “Application of Graphene-Based Materials”, which consists of a total of twenty-six articles, including two review articles, written by research groups of experts in the field, considers the most recent research and trends on the synthesis and characterization of graphene-based materials, including nanohybrids, intended for a vast array of high-demanding technological applications, namely batteries/fuel cells, aerogels, laser technology, sensors, electronic/magnetic devices, catalysts, etc.One of the reviews included in the present Topic considers the most recent research of graphene-based materials, with the authors even calling it “graphene technology”, to be used in lasers . Namely,The unique characteristics of graphene-based materials make them ideal candidates to be considered for energy storage purposes and in the high-trending market of batteries/fuel cells. Guo et al.  considerDue to its pressure-sensitive properties, graphene (and its derivatives) has found interest in the development of sensors, and, due to its highly tunable electrical conductive properties, more specifically, piezoresistive sensors. Florencio and co-workers  have preOne of the most interesting uses of graphene-based materials in the field of materials science is in the development of ultralight aerogels for different applications, from enhanced wave absorption to liquid adsorption. In this sense, Mei et al.  considerOwing to their highly interesting transport properties, namely high and tunable electrical conductivity and magnetic characteristics, graphene-based materials have generated a great deal of interest in applications where electron charge transport may be of crucial importance, as in electronics. For instance, the group of Roshchupkin  investigIn addition to their highly desirable characteristics for charge transport, graphene-based materials also display theoretically high mechanical characteristics, such as high strength, while keeping a significant flexibility. As a consequence, especially in the field of materials science, a great number of research groups have focused their interest in using graphene or its derivatives as nanoreinforcements, alone or in combination with other (nano)reinforcements, in different types of matrices, especially polymer-based. An example of this use for graphene can be seen in the work of Lim and co-workers , which c2, and reduced graphene oxide. The nanocomposite showed an extremely high level of photocatalytic activity through hydrogen production by water splitting, far exceeding that of undoped TiO2 and reduced graphene oxide, hence showing promising application as hydrogen evolution catalyst.A high interest has been dedicated to the possible use of graphene-based materials as part of a (nano)catalyst system. An example of this use is the work of Zhao and colleagues , which hAlongside all the abovementioned possibilities of graphene-based materials, there are others that, although less typical, are also being considered by researchers, presenting another example of the versatility of graphene. Teshigawara et al.  proposedOwing to its particular flat-like surface, together with its high specific area, graphene materials have been considered as interlayers in the development of several devices. An example of this can be seen in the work of Badokas and co-workers , in whicLess seen, but also considered, are the possible applications of graphene related to food. In this Topic, the study of Peng et al.  dealt wi


Chiral graphene hybrid materials have attracted significantattentionin recent years due to their various applications in the areas ofchiral catalysis, chiral separation and recognition, enantioselectivesensing, etc. On the other hand, chiral materials are also known toexhibit chirality-dependent spin transmission, commonly dubbed “chiralityinduced spin selectivity” or CISS. However, CISS propertiesof chiral graphene materials are largely unexplored. As such, it isnot clear whether graphene is even a promising material for the CISSeffect given its weak spin–orbit interaction. Here, we reportthe CISS effect in chiral graphene sheets, in which a graphene derivative(reduced graphene oxide or rGO) is noncovalently functionalized withchiral Fmoc-FF  supramolecular fibers. The grapheneflakes acquire a “conformational chirality” postfunctionalization,which, combined with other factors, is presumably responsible forthe CISS signal. The CISS signal correlates with the supramolecularchirality of the medium, which depends on the thickness of grapheneused. Quite interestingly, the noncovalent supramolecular chiral functionalizationof conductive materials offers a simple and straightforward methodologyto induce chirality and CISS properties in a multitude of easily accessibleadvanced conductive materials. Thus,chiral media can act as spin polarizers or analyzers, which are criticalcomponents of spintronic devices. The physical origin of CISS is stillunder debate, although it is generally assumed that the spin–orbitinteraction of the medium plays a critical role.1Chirality-induced spin selectivity (CISS)refers to a set of phenomenain which chirality of a material imparts significant spin selectivityto various electronic processes.6 However, these molecules are electrical insulators, which pose challengesfor their integration into electronic devices and circuits. In addition,the formation of reliable electrical contacts with molecules givesrise to additional difficulties. On the other hand, various inorganicmaterials that are considered promising platforms for future electronicsand spintronics are generally nonchiral, and hence these materialsdo not exhibit the CISS effect in their natural forms. Some alternativeapproaches have been investigated to bridge this gap. For example,it has been reported that CISS can be induced in carbon nanotubesby attaching chiral molecules on the tube walls.14 Several studies on chiral inorganic crystals have also been reported.18 However, rational integration of chiral organic materials with emergingadvanced materials with promising electronic and spintronic propertiesis expected to significantly expand the application of CISS in thearea of solid-state spintronics. Therefore, the development of strategiesthat allow interesting materials to be endowed with chirality in asimple and effective way is a highly desired goal in this field.Many organic molecules such as DNA, amino acids,peptides, helicenes,etc. are inherently chiral, and hence they are commonly employed forCISS studies.20 From the viewpoint of spintronics,room-temperature spin transport and micron-scale spin relaxation lengthin pristine graphene layers have been reported.22 Such features originate due to the weak intrinsic spin–orbitinteraction of graphene and the weak hyperfine interaction of electronspin with carbon nuclei, which suppress the spin relaxation mechanisms.On the other hand, it has been reported that various chiral organicmolecules can impart chirality to various graphene derivatives suchas graphene oxide (GO) and reduced (or deoxygenated) graphene oxide(rGO) via covalent as well as noncovalent interactions including hydrophobiceffect, π–π stacking, electrostatic effect, etc.30 In most cases, the chirality of the resultant hybrid materials originatesfrom the chirality of the attached moiety. These have been classifiedas “configurational chirality” , “conformational chirality” , or “hierarchical chirality” .26 Reference .Graphene is covered by the hydrophobic peptide fibers in which thearomatic rings of Fmoc-FF molecules bind efficiently with the planarsurface of graphene via the π-π interaction. Controlexperiments are performed with Fmoc-GG (Fmoc-diglycine) molecules,which are achiral. As described below, the chiral molecules self-assemblein helical supramolecular structures along with the attached grapheneand induce the CISS effect in two-terminal magnetotransport measurements.The role of graphene flake thickness has been investigated by usingsingle-layer graphene (SLG) and multilayer graphene (MLG) flakes (bothin rGO form).Figure S2 (Supporting Information).37 This band determines the supramolecular arrangementof these peptides and therefore is the one related to the CISS effect.13CD spectra of Fmoc-FF supramolecular fibers showedmirror imagesfor both enantiomers, indicating that the chirality of the supramolecularaggregates is dictated by the intrinsic chirality of the peptide.These aggregates show two characteristic Cotton bands, one around220 nm which corresponds to n−π* transition of the aminoacids and another band around 260–270 nm which correspondsto the π–π* transition of the fluorenyl groupsand was previously justified as superhelical arrangements of thesepeptides.38 As such, whensolid particles are present in the media, they are engulfed by thepeptide fibers in formation, including them in the supramolecularpeptide structure and giving rise to hybrid materials.41 This processfavors a strong interaction between the peptide fibers and the solidmaterial that ends up with its surface completely functionalized withthe peptides. Moreover, if the solid material is flexible, it cansuffer morphological alterations in its structure, the peptide fibersacting as a template.43 In this process, the supramoleculararrangement of the peptides can also be altered by the interactionof the peptides with the solid particles, resulting in some casesin an inversion of the supramolecular chirality, as observed by changesin the sign of the Cotton bands in CD spectra. We have previouslyobserved this effect when the self-assembly of Fmoc-FF in the presenceof SWCNT  is triggered by Na2CO3 or GdL12 or in the presenceof CNTs of different diameters.14 In thiscase, we observed the same effect upon passing from the peptide solutions, the well-known D (∼1350 cm–1), G (∼1580 cm–1), and D′(∼1610 cm–1) peaks are present.44 While the strongest peak in unfunctionalized(raw) SLG and MLG is the G peak, the D′ peak appears as the dominant peak in the functionalizedgraphene samples. The appearance of the strong D′peak after functionalization indicates the presence of defects andstructural disorder. It should be noted that the D′ peak is relatively weak and can be challenging to distinguishfrom background noise in the raw samples; however, the presence ofa strong D′ peak in conjunction with the D and G peaks after functionalization indicateslattice distortion. As commented upon above, this can arise from theself-assembly of the peptide molecules, which inflicts strain on theattached SLG and MLG layers, as indicated by the twisted and rolledlayers of graphene in –1, which is associatedwith the functionalization-induced structural changes and the vibrationalmodes of the functional groups,45 and alsoat ∼1294 cm–1, which is related to the oxygen-containingfunctional groups and carboxyl (−COOH) groups.46 In addition, we observe weak Raman active bands such as2D (∼2710 cm–1) and D+G (∼2950 cm–1) at higher wavenumbers. The presence of defects anddisorders broadens and weakens the 2D peak significantly, which isa characteristic of rGO.47Raman spectra of raw andfunctionalized graphene flakes, measuredin ambient air with an excitation wavelength of 532 nm, are shownin 48 A FESEM image of the final device is shown in A two-terminal planar CISS device is fabricatedby placing a sliceof the gel described above between Au–Ni contacts. The deviceis vacuum annealed to improve the electrical contacts as well as toimprove the connectivity between the individual chiral flakes.I–V) characteristics of the Fmoc-FF+SLG samples with both chiralities,measured at zero magnetic field. The I–V plotsas well as the temperature-dependence of resistance of the Fmoc-GG+SLG/MLGsamples are shown in the Supporting Information . The I–V characteristics have been found to be nonlinear relative to biasand semiconducting in terms of temperature dependence. Similar behaviorhas been reported in the literature for monolayer51 as well asmultilayer rGO.52 In rGO, regions of highlyconductive graphene islands with delocalized states are separatedfrom each other via disordered regions or “barriers”with localized states, and conduction is limited by the temperatureactivated hopping mechanism through the localized states.52 In this case, the attachmentof chiral molecules is also responsible for these barriers. As biasvoltage is increased, electrons acquire sufficient energy and percolatethrough the flakes via multiple branches (field-assisted tunneling),which gives rise to the nonlinear bias-dependence. It is to be notedthat in the present case, the graphene flakes are functionalized withchiral molecules, which is expected to induce CISS during transport.I–V responsein rGO or randomly distributed graphene islands is typically explainedin terms of the variable-range hopping (VRH) model,52 in which resistance R scales withtemperature T as follows:The temperature-dependent R ∝ exp[(T0/T)p], where T0 is a parameter knownas “characteristic temperature”and p is a fractional exponent equal to 1/(d + 1), where d is the dimensionality ofthe system. T1/3. A linear dependence is found, which isconsistent with the 2D structures of the rGO flakes.Figure S3). In the case of MLG, only the surface layer is functionalized,whereas the inner layers are not, and the transport occurs via bothinner layers and the surface layers. The unfunctionalized inner layersoffer fewer barriers and hence weaker temperature dependence of resistivityfor the MLG samples.The temperaturedependence of resistance is weaker in MLG samplescompared with the SLG samples (R(+12 kG) = R(−12 kG), as shownin the insets. A negative MR, defined as R(0) – R(±12 kG)/R(±12 kG), of ∼15%is observed at 11 K for the Fmoc-GG+SLG samples, which gradually weakenswith increasing temperature. As described above, the system is nota weakly disordered metallic system, and hence the standard weak-localizationtheory cannot be used to explain this negative MR. However, even inthe VRH regime, such negative MR has been reported to appear, dueto magnetic field driven carrier delocalization.53R(+12 kG) = R(−12 kG), as shown in the insets. Thicker MLG layersare analogous to graphite, which is known to exhibit positive in-planeMR in the presence of an out-of-plane magnetic field.54 This is presumably because transport occurs via multiplelayers, which offers significantly more carrier pathways than SLG.In each path, the carriers experience deflections due to magneticfield induced Lorentz force, which tends to deviate the carriers,resulting in a positive MR.In general, MLG samples show lowerresistance values compared withtheir SLG counterparts. We believe that the availability of multipleinner current paths via the inner layers of MLG causes lower resistance.The inner layers are not functionalized as discussed above, and hence,they offer fewer barriers, which lowers the overall resistance ofthe MLG samples.R(+12kG) <>R (−12 kG) for L and D chirality, respectively.ForMLG samples, as shown in R(+12 kG)<>R (−12 kG) for L and D chirality,respectively.2 The electronstraveling through chiral (L/D) functionalized graphene flakes acquirea chirality-dependent spin polarization (up/down), which is eithertransmitted or blocked depending on the magnetization of the Ni spindetector, resulting in a chirality dependent asymmetry in the MR.Opposite chiralities induce opposite spin polarizations, and hence,the MR asymmetry is reversed for L and D chiralities. It is to benoted that in our experiments we have kept the magnetic field perpendicularto the sample plane and hence perpendicular to the planar currentpaths. This ensures that the results are minimally influenced by anyelectromagnetochiral effect.55 Observationof a CISS signal in such a transverse configuration also confirmsthe existence of transverse spin components perpendicular to the currentpath. This has been reported before in functionalized carbon nanotubes13 and is true for functionalized graphene as well.For ideal helical systems, the net transverse spin component is expectedto be zero.56 However, in reality, thedirection of spin depends on various factors such as electron energyor the details of the chiral medium and its coupling with the electrodes.56 These can result in nonzero transverse spinpolarizations.Such chirality-dependent MR asymmetry can be explainedby invokingthe standard CISS phenomenology.14 Even though the detailsabout the mechanism of chiral induction are not fully known, it isworth noting that simple CD measurements of supramolecular aggregatesgave predictable information about the asymmetry in the MR responses.Consequently, achiral fibers with null CD resulted in no spin selection.Although a correlation between the CISS effect and the sign and magnitudeof the CD of the conductive material has been suggested in earlierreports,59 the present result shows that the CD signal of thechiral inductor could also infer the CISS response of the compositematerial.The opposite MR asymmetries of the SLG and MLGsamples correlatewith their respective chiralities, as discussed earlier in RAsym(B) = [R(B) – R(−B)]/2, shownin RSym(B) = [R(B) + R(−B)]/2, as plotted in the insetsof Sym = |RSym(0 kG)– RSym(±12 kG)|. Similarly,the strength of the CISS-dependent asymmetric component is computedas δAsym = RAsym(+ 12kG) – RAsym(−12 kG), andthe CISS signal is represented as δAsym/δSym × 100%.Since the role of the chiral molecules is to introducean asymmetryin the MR response, we isolate this contribution by computing theodd component of the MR function, defined as Asym/δSym responsesfor SLG and MLG samples, respectively, showing their temperature dependence.The CISS signal has been found to decrease with increasing temperature.While this is in contrast with the molecular systems, such temperaturedependence is often observed in solid-state CISSdevices.60 We also observe that the δAsym/δSym values of the MLG samples are generallylower than those of the SLG samples under the same conditions. Thisis expected, because, as discussed above, transport in MLG occursvia multiple inner layers, and these layers are less affected by thechiral molecules, which are attached on the surface layers. We measuredfour different SLG samples and nine different MLG samples. For theSLG samples, since the signal was stronger, we typically measuredeach device two or three times at a given temperature and bias. Forthe MLG samples, since the signal was weaker, we typically measuredeach device five or six times. The signals are reproducible, and theerror bars shown in Figure S4 (Supporting Information), we have added data obtained from a different set of samples, showingthe reproducibility and variability of these results.1 It is generally thought that spin–orbitinteraction, along with spatial inversion asymmetry  and time inversion asymmetry , are the necessary components for observation of this phenomenon.1 Spin–orbit interaction of pristine grapheneis weak, due to the low atomic number of carbon,61 and hence the observation of the CISS effect discussedabove is somewhat surprising. However, in the present case, thereare several factors that may mitigate this effect:A self-consistenttheoretical model of the CISS effect currentlydoes not exist, and it is a subject of significant research activity.61 Such curvature-induced spin–orbitinteraction is estimated to be an order-of-magnitude stronger thanthe intrinsic spin–orbit coupling.61 In our case, on the molecular level, the chiral entities form ahelical supramolecular structure. As discussed above, the chiral moleculesare strongly attached to the graphene layers via π–πinteractions. Hence, when these molecules self-assemble to form superhelicalstructures, so do the graphene layers attached to them , which makes the graphenelayers curved, and this can enhance the spin–orbit coupling.The TEM images in (a) First,it has been reported by theoretical calculations thatcurved graphene layers have an additional spin–orbit term thatarises due to the mixing of σ and π bands due to localcurvature.62 Such impurities can cause local latticedistortion via sp3 hybridization, which can significantlyenhance the local spin–orbit interaction.64 According to some studies, such enhanced spin–orbit interactionscan approach values comparable to those in zinc-blende semiconductors.63 Such spin–orbit enhancements are presumablyresponsible for the CISS signal in rGO, even though it is not expectedin pristine, impurity-free graphene. We note that such adatom-inducedenhancement of the spin–orbit interaction is responsible forsmall spin relaxation times in graphene.65(b) Second, and perhaps moreimportantly, the system under considerationis rGO, instead of pristine graphene. Due to the synthesis mechanismof rGO, there exist impurity atoms (such as hydrogen) on the graphenelattice.66 Carbon vacancies,which typically arise during GO synthesis and its thermal exfoliationreduction, are well-known magnetic defects.67 Such local magnetic interactions can alsocontribute to the observed CISS effect.Apart from the above factors that enhance the spin–orbitcoupling in rGO, we note that there are several sources of magneticdefects, as well. The hydrogen adatoms discussed above can give riseto local magnetic moments.I, computed as I(+12 kG) – I(−12 kG) asa function of the applied voltage bias. In a two-terminal measurementgeometry with only one magnetic contact, as shown in I is expected to be zero according to Onsager’s reciprocityprinciple, at least in the linear region.68 As seen from I approaches zero as voltage bias approacheszero, i.e., as the device enters the linear region of operation. Thecurrent differential ΔI due to the CISS effectincreases with bias and manifests in the nonlinear region of transport in the present measurement geometry. This resultis consistent with that observed recently in CNT based CISS devices.13 The CISS signal saturates around ∼0.5V, and hence, the MR curves discussed in 14Graphene is a promising 2D material foremerging electronic andspintronic devices; however, its compatibility with the CISS effectwas never explored before. In fact, weak spin–orbit interactionof graphene makes the occurrence of CISS unlikely in this material.Nevertheless, in this work, we demonstrated that the CISS effect manifestsin rGO flakes functionalized by chiral dipeptide molecules. The chiralmolecules organize in supramolecular structures, resulting in a “conformationalchirality” of the attached graphene flakes. This introducesa curvature in the graphene flakes, as shown by the microscopic imagesdiscussed in this paper. Such curvatures as well as impurities inrGO can enhance the spin–orbit interaction and result in theCISS signal. The simple and straightforward methodology describedin this work can be used to induce chirality and CISS properties ina multitude of easily accessible advanced 2D conductive materials,vastly expanding device design opportunities and associated applications.The observed CISS signal correlates with the supramolecular chiralityof the medium, emphasizing the importance of “global chirality”as opposed to “local chirality” in the CISS phenomenon.Figure S1 (Supporting Information). As purchased SLG samples , prepared by thermal exfoliationreduction and hydrogen reduction of graphite oxide (GO), consist ofone to five atomic layers of graphene with a typical flake size of∼0.5–5μm. The oxygen content is estimated to be∼7–7.5%. The MLG samples were procured from TCI Europe and are6–8 nm thick nanoplatelets with a width of ∼15 μm.No exfoliation step has been employed for the MLG samples; hence,it is expected that the chiral functionalization (described below)occurs only on the surface, whereas the inner layers remain imperviousto it. This allows an investigation of the effect of sample thicknesson the observed behavior.The molecular structures of Fmoc-FF and Fmoc-GGare shown in l-diphenylalanine (Fmoc-FF(L)) and N-fluorenylmethoxycarbonyl-d-diphenylalanine (Fmoc-FF (D)) were purchased from LifeTein,USA. N-Fluoroenylmethoxycarbonyl-diglycine (Fmoc-GG) was purchasedfrom Fluorochem, UK. All Fmoc-peptides were used without further purification.N-Fluorenylmethoxycarbonyl-To obtain the Fmoc-peptides + SLG/MLGcomposites, basic solutions of Fmoc-FF (L/D) and Fmoc-GG were firstprepared. Fmoc-FF (L/D) and Fmoc-GG peptides were weighed separatelyinto a vial, and deionized water was added to obtain a final concentrationof 10 mM. The suspension was sonicated (HSt Powersonic 405 ultrasonicbath) for 1 h. Then, a NaOH solution (0.5 M) was added dropwise untila clear solution was obtained (pH = 10.7). The pH was measured usinga HACH sensor PH 3 pH meter. The pH meter was calibrated using pH4, pH 7, and pH 10 buffer solutions.To prepare the grapheneflake peptide solution (SLG or MLG), 0.7 mg of each graphene flakewas separately weighed in a vial tube. The graphene flakes were suspendedin 1 mL of a basic aqueous solution of Fmoc-FF (L/D) or Fmoc-GG (preparedabove). The suspension was sonicated for 2.5 h in a cold ultrasonicbath and then centrifuged for 5 min at 10 000 rpm (Sigma 1–14centrifuge). Finally, the supernatant was carefully collected.12 Fmoc-GG + SLG/MLG hydrogels were obtained using Na2CO3 as a gelling agent by adding a final concentrationof 25 mM sodium carbonate and mixed by vortexing.11Final hydrogels were obtained using GdL (glucono-δ-lactone)as a gelling agent for the Fmoc-FF + SLG/MLG composite solutions byadding 2 mol equiv of GdL and mixing by vortexing.The CD spectra were recordedusing a Jasco J-815 spectropolarimeter with a xenon lamp of 150 W.The samples were measured into a 0.1 mm quartz cell (Hellma 0.1 mmquartz Suprasils), and the spectra were obtained from 200 to 350 nmwith a 1 nm step and 1 s integration time per step at 20 °C.Measurements were performed with a 1:1 dilution ratio of peptide basicsolutions or composite solutions in MiliQ water to keep sample absorbancebelow 2.–1, andeach spectrum resulted from the average of three acquisitions, with50 s of accumulation each. The samples were deposited on glass slidesdirectly from the bottle and crushed to form a thin film.Raman spectra of commercial graphene flakes were collectedon a Raman microscope NRS-5100  equipped with a Peltiercooled charge-coupled device (CCD) (1064 × 256 pixels) detector.The excitation line, provided by a diode laser emitting at 532 nm,was focused on the surface of the sample through a 100× objectivelens. The spectral resolution was 2.1 cmForTEM, the samples were sonicated for 10 min and then drop castonto a TEM Cu grid, coated with lacey carbon film. After 24 h dryingin a normal atmosphere, the grid was mounted on the JEM-ARM200cf S/TEM,to obtain the TEM images.


Impedance spectroscopy confirms that theorganic semiconductor is depleted, and the energy band diagram resultsin two electron blocking interfaces. The resulting rectifying natureof the Gr/C60 interface could be exploited in organic hot electrontransistors and vertical organic permeable-base transistors.Among the familyof 2D materials, graphene is the idealcandidateas top or interlayer electrode for hybrid van der Waals heterostructuresmade of organic thin films and 2D materials due to its high conductivityand mobility and its inherent ability of forming neat interfaces withoutdiffusing in the adjacent organic layer. Understanding the chargeinjection mechanism at graphene/organic semiconductor interfaces istherefore crucial to develop organic electronic devices. In particular,Gr/C60 interfaces are promising building blocks for future n-typevertical organic transistors exploiting graphene as tunneling baseelectrode in a two back-to-back Gr/C60 Schottky diode configuration.This work delves into the charge transport mechanism across Au/C60/Grvertical heterostructures fabricated on Si/SiO Infact, graphene is typically implemented as a gateable bottom electrodeinto barristors.12 However, more complex multilayer architectures could exploit theunique properties of graphene, i.e., optical transparency,13 flexibility,14 andthe ability to conform to three-dimensional and/or rough surfaces.15 For instance, graphene could be implementedas the base electrode in organic hot electron transistors, where ultrathinorganic layers are required, or replace the metallic base electrodein vertical organic permeable-base transistors.17 This would enable flexible vertical organic transistors operatingat high frequencies18 that are suitablefor applications such as wearable electronics with communication functions,or RF tags. Flexible vertical transistors could also be verticallyintegrated in OLEDs,16 enabling high-resolutionorganic displays with enhanced contrast. Within this framework, understandingthe charge transport mechanisms occurring at the interface betweengraphene and organic semiconductors (OSC) is crucial for the developmentof new graphene-based optoelectronic devices.20Graphene is the ideal candidate as topor interlayer electrodefor hybrid van der Waals heterostructures made of organic thin filmsand 2D materials due to its high conductivity and mobility and itsinherent ability of forming neat interfaces without diffusing in theadjacent organic layer, in contrast to commonly used metal electrodes.21 In this case, however, the graphene sheet is transferredon top of a thermally evaporated thin film of small molecules, i.e.,the n-type organic semiconductor C60, and subsequently patterned intothe top electrode. The fabrication processes combines techniques commonlyused in the semiconductor industry, and therefore it can be upscaledto wafer level. The charge injection mechanisms at the interfaces,i.e., Gr/C60 and Au/C60, are investigated by temperature-dependent I–V measurements and impedance spectroscopy.The current–voltage characteristic is modeled by the doubleSchottky barrier (DSB) model at room and high temperature (above 300K) and by Fowler–Nordheim (FN) tunneling at low temperature(below 100 K). The energy barrier height at the Gr/C60 and Au/C60interfaces is extracted from the DSB and from the FN tunneling models,while the static dielectric constant of C60 is determined by impedancespectroscopy. The study concludes with the energy band diagram ofthe heterostructure explaining the charge injection at the two interfaces.The fabrication techniques presentedin this work are similar towhat previously reported for vertical VdW heterostructures made ofthe p-type polymer P3HT and CVD grown graphene.Vertical Stacks, with diameters spanning from 5 to 50 μm and C60 film thicknessof 80 nm, and (ii) 34 Graphene Bridges to characterizethe in-plane resistance of graphene after transfer on C60. Vertical Stack, fabricated followingthe protocol reported in the Supporting Information, togetherwith the electrical scheme implemented for the measurements, whilethe schematic of the Graphene Bridge and the correspondingelectrical scheme are shown in the Supporting Information . Vertical Stack having a diameterof 10 μm. From the AFM height image one can extract the actualdiameter of the graphene electrode , while theheight profile shows the thickness of the bottom electrodes (35 nm)and of C60 (80 nm). Refer to the Supporting Information for the AFM height image of a representative Graphene Bridge device . 2 (335 nm), Ti/Au (5 nm/30 nm), and C60 (80 nm). Theresolution of the instrument does not allow to observe the grapheneelectrode on top of the stack. 2/Gr) shows the characteristic G (1582 cm–1) and 2D (2674 cm–1) peaks, as well as the D (1340cm–1) peak, with a weak amplitude, possibly resultingfrom fabrication-induced defects. The C60 film (SiO2/C60),the C60 film covered by graphene (SiO2/C60/Gr), and thefull stack (SiO2/Au/C60/Gr) Raman spectra are virtuallyidentical and display the characteristic Raman-active vibrations ofC60,22 i.e., Hg (1) at 264 cm–1, Hg (2) at 431 cm–1,Ag (1) at 492 cm–1, Hg (3)at 709 cm–1, Hg (4) at 774 cm–1, Hg (5) at 1099 cm–1, Hg (6) at 1244 cm–1, Ag (2) at 1463 cm–1, and Hg (8) at 1568 cm–1. The typical Raman scattering of Si,23 i.e., the first-order optical mode (520 cm–1)and second-order scattering band (940–980 cm–1), due to the Si substrate, is observed in all spectra. The resultsof the AFM, SEM, and Raman analysis show that the methods used tofabricate the vertical VdW devices are compatible with graphene andC60 and do not degrade these materials.This study was performed onone chip that includes two distinctgroups of devices: (i) 119 Au/C60/Gr ormation S1. Figurormation S1. Vertiormation S1. Vertiormation S1. VertiI–V characteristics measured in air and a vacuum and by studyingthe scaling relationship between the system resistance and capacitanceand the device area, where the device resistance and capacitance areobtained from impedance measurements under vacuum. Oxygen and moistureare known to affect the electrical properties of C60 and graphene.29 In particular, oxygen and moisture typically act as electron trapcenters for C60, reducing the electrical conductivity of the film,29 while they typically p-dope graphene on SiO2,26 as shown in the Supporting Information for the Graphene Bridges devices  and asdemonstrated on graphene field effect transistors . In order to desorb oxygen and water from C60 andfrom the graphene surface, and therefore ensure the consistency andreproducibility of measurements done in different systems, all measurementsare done in a vacuum after annealing . I–V traces offive devices per area measured under vacuum (∼10–6 mbar) at room temperature (293 K) after annealing for 12 h at 110°C. For the purpose of clarity, the figure does not display the I–V data in the range from −5to +5 V, where the current is below the sensitivity of the instrument.The inset of J–V tracesof all devices overlap and, thus, that the transport characteristicis not affected by the device area. The small variability between J–Vs is ascribed to the device inhomogeneitiesdue to the fabrication. After annealing, the J–V characteristics become very similar, while a larger variabilityis observed for the J–V tracestaken in air and shown in Figure S5. Interestingly,the stacks become less conductive in a vacuum after annealing , against what one would expect ifthe charge transport were bulk limited by C60 due to oxygen and moisturedesorption. This observation suggests that the transport is not bulk-limitedbut injection-limited and that the oxygen and moisture have an impacton the energetics of the interfaces. All J–Vs display the same nonlinear and asymmetric S-shaped characteristic:the current initially increases slowly with increasing voltage andthen rapidly increases at higher voltages; the same occurs for negativevoltages, but the magnitude of the current is higher. Similar trendsare typically described by various analytical models for the chargeinjection at the metal/semiconductor interface, i.e., thermionic emission(TE),30 the modified thermionic emission(MTE) developed for graphene/semiconductor interfaces,32 direct tunneling, and Fowler–Nordheim (FN) tunneling.30 However, in most cases the charge injectionat the metal/OSC follows a hybrid process,33 which includes tunneling and thermionic emission34 with deviations from the standard theory due to defects,surface inhomogeneity, and image-charge barrier lowering, resultingin a bias-dependent barrier height.35 This hybridprocess can be described by a nonideal Schottky diode model with idealityfactor n,30 where thereverse current is I = SA**T2 exp(−Φ)/kT),and where S is the contact area, T is the absolute temperature, A** is the reducedeffective Richardson constant, k is the Boltzmannconstant, and Φ = Φ0 ± (1 – 1/n)qV is the voltage-dependent energy barrierwith nominal barrier height Φ0 and ideality factor n and where q is the elementary charge.The model can be extended to a double barrier system, where differentpotential barrier heights are formed at two interfaces in series.This model is called the double Schottky barrier (DSB) model and isdescribed by35VerticalStacks of this work, IS1,2 = S1,2A**T2 exp are the reverse saturation currents, and Φ1,2 = Φ01,02 ± qV are the voltage-dependentenergy barriers at the Gr/C60 and Au/C60 interfaces, respectively.In this model, the current is limited by the reverse currents. When V > 0, the Gr/C60 diode (SB1) is forward biasedand the Au/C60 diode (SB2) is reversed biased. Therefore,the measured current is mostly given by the reverse current of SB2, i.e., IS2. On the contrary, when V < 0, the currentis limited by SB1, i.e., IS1. The DSB model is applied to the I–Vs in S1 = π(d1/2)2 for the graphene top electrode and S2 = π(d2/2)2 for the Au bottom electrode, where d1 and d2 are the diameter of thecontact area and d2 is 2 μm largerthan d1, while the reduced effective Richardsonconstant for both interfaces is set to A** = 100A m–2 K–1. It is worth observingthat A** is extracted from temperature-dependentmeasurements, and it varies in the range 101–103 A m–2 K–1 depending onthe device and voltage , as typically reportedfor metal/OSC interfaces.36 The variation of theenergy barriers extracted with different A** fromthe aforementioned range is less than 0.1 eV , and it does not affect the conclusionsof this work. Figure S9). 01,02 and ideality factors n1.2 of the Gr/C60 and Au/C60 interfaces, respectively. Thenormal distributions are given by Φ01 = 0.58 ±0.01 eV, Φ02 = 0.65 ± 0.02 eV, n1 = 1.033 ± 0.002, and n2 = 1.036 ± 0.003. The potential barrier height of the Au/C60interface is similar to previously reported value for C60 field-effecttransistors with Au electrodes.37 Impedancespectroscopy reveals that the Vertical Stacks exhibitthe typical nonideal capacitor characteristics. R and capacitance C of the Vertical Stacks obtained fromimpedance spectroscopy by fitting the modulus and phase with a nonidealcapacitor model, i.e., an R||C circuit. Refer tothe Supporting Information for detailson the fit . The resistancedepends on the applied bias, as expected from the I–V characteristics, while the capacitanceis bias-independent, therefore suggesting that the C60 is fully depleted38 and that the voltage drops linearly over thewhole structure. As a consequence, the electrostatic doping of graphene predicted by the MTE32 can be discarded, and the position of the EF of the graphene electrode is determined solely by thechemical doping of graphene. It is known that C60 adsorbed on graphenetypically results in p-doping.39 This isdemonstrated by a comparison between field-effect measurements ona set of Gr channels before and after deposition of C60 . The charge neutrality point of thegraphene field transistors shifts to positive gate biases (ca. 20V) when a thin film of C60 is deposited on top of graphene, resultingin an average p-doping of about 1.4 × 1012 cm–2 . S, where S is the area of the top electrode, whilethe capacitance scales as S  and ϵ0 is the vacuum permittivity.This value is in agreement with previous works.41 The results of the DSB analysis, supported by the impedance measurements,suggest that the vertical charge transport in the Au/C60/Gr is limitedby charge injection. Further insight into the energy barrier and chargeinjection mechanisms at the Gr/C60 and Au/C60 interfaces is obtainedby temperature-dependent measurements.The effects of the environmentand of the devicearea on the chargetransport are investigated by comparing the J–V characteristics as a function of temperature of a representative50 μm device, in linear and logarithmic scale, from 50 to 285K in steps of 5 K ; I–V characteristics for a representative 50μm device as a function of temperature, from 300 and 380 K insteps of 5 K; I–V characteristic ofa representative 50 μm in the high-temperature range between300 and 380 K. In this range, the charge injection is dominated bythermionic emission, and the DSB model can fit all I–Vs. The inset of 01, Φ02 and ideality factors n1, n2 as a function of temperature,extracted from the DSB fit for the Gr/C60 and Au/C60 interfaces, respectively.The ideality factor of both interfaces decreases with increasing temperature,indicating that the device approaches the ideal TE injection modelat higher temperatures, with values around n1 = 1.026 and n2 = 1.031. Simultaneously,the potential barrier heights decrease with temperature. This shouldnot come as a surprise, as it is known that barrier inhomogeneitycan lead to temperature-dependent energy barriers.43 Below 100 K, the J–Vs areapproximately temperature-independent and well described by the FNtunneling mechanism under the WKB (Wentzel–Kramers–Brillouin)approximation in the formula30m is the effectivemass of the electron in the semiconductor and ℏ is the reducedPlank constant. Here, the free electron mass is considered. The insetof J–V traceof a representative 50 μm device and the corresponding FN fit,while J/V2 vs 1/V plot at high bias  for 5, 10, 15, 20, 25, 30, and 50 μm representativedevices at 50 K. The J–Vsare very similar, and the energy barrier of the two interfaces canbe derived from the slope of the ln J/V2 vs 1/V plot at |V|> 8 V, resulting in Φ01 = 0.54 ± 0.07 eVandΦ02 = 0.60 ± 0.05 eV, which are similar to thoseobtained from the DSB model fit at high temperature. Refer to the Supporting Information  for the I–V traces in the range from −10to 10 V. It is worth observing that the asymmetry of the J–Vs characteristics changes from high tolow temperatures. In particular, the rectification ratio at ±9V, that is RR = |(J(+9 V)/J(−9 V)|, is ca. 1.14 at 50 K and 0.13 at 380 K.Although it is hard to explain quantitatively the phenomenon, oneshould observe that the injection mechanism changes from nonidealTE at high temperatures to FN tunneling at low temperatures, whichare described by different models with different prefactors. The latterare known to deviate significantly from theoretical values in nonidealand/or metal/OSC systems.36 A transition from FNto direct tunneling would only be possible with a thinner semiconductor.The C60 thickness for transition between FN tunneling and direct tunnelingat room temperatures can be approximated by t = ΦB/E, where E is the electricfield.30 Assuming an energy barrier of0.5 eV, which corresponds roughly to the measured value for the Gr/C60interface, and an electric field of E = 10 V/80 nm= 1.25 MV/cm, this results in a C60 thickness of ∼4 nm. Few-nanometerthin vertical stacks can be theoretically achieved using the architectureand fabrication process described in this work, where graphene isused as the top electrode. In fact, this fabrication process allowsto avoid the typical metal atoms intercalation and diffusion observedfor metal electrodes evaporated on organic films.1 The study of ultrathin stacks could provide insights intothe quantum charge transport across OSC/Gr VdW vertical heterostructures,besides potentially enabling FN tunneling at room temperature andtherefore making possible the fabrication of organic flexible memories.The latter, which exploit FN tunneling to perform the write and eraseoperations, could be developed by stacking multiple C60/Gr interfaces.Temperature-dependentelectrical measurements arenecessary todistinguish the different charge injection mechanisms contributingto the electrical current in the devices and possibly explain thenonideal thermionic emission observed at room temperature. Low- andhigh-temperature-dependent electrical measurements are done in twodifferent systems as described in the EF,Gr = ELUMO – Φ01 = −4.7 eV, whichis possibly shifted with respect to the Dirac point due to p-dopinginduced by C60 and/or to graphene defects, and (ii) the Au work functionis ca. W = ELUMO –Φ02 = −4.8 eV, as typically observed for Ausurface with organic adsorbates, due to the pillow effect and/or interfacestates, which results in an interface dipole and possibly Fermi levelpinning at the Au/OSC interface.45 It is worthobserving that because the transport is injection limited, the thermalannealing has an observable effect on the energetics of the barrier:oxygen and moisture desorption affect the surface states, and thuspossibly the Fermi level pinning, resulting in more or less conductingdevices, as evinced by the I–V characteristics before and after annealing .This work describes the field and thermalinduced charge transportmechanisms in Au/C60/Gr VdW vertical heterostructures in the 50–380K temperature range. Devices having various nominal diameters, whichrange from 5 to 50 μm, show the same electron transport characteristics.Thus, the fabrication of smaller devices is only limited by the lithographyresolution and the fabrication process presented in this work.r = 3.8. The electrical transport measurementsacross the vertical heterostructure at high temperature (above 300K) are well described by the DSB model, which gives Φ01 = 0.58 ± 0.01 eV, Φ02 = 0.65 ± 0.02 eV, n1 = 1.033 ± 0.002, and n2 = 1.036 ± 0.003 from the measurements at room temperature(293 K) at the Gr/C60 and Au/C60 interfaces, respectively. At lowtemperature (50 K), the measurements are well described by FN tunneling,which gives Φ01 = 0.54 ± 0.07 eV and Φ02 = 0.60 ± 0.05 eV for the Gr/C60 and Au/C60 interfaces.The energy barriers extracted from the two models are very similar.These findings reveal that the charge transport is injection-limitedin the temperature range from 50 to 380 K. The current is dominatedby TE at high temperature (above 200 K) and limited by FN tunnelingat low temperature (below 100 K). The charge transport at intermediatetemperature (from room temperature down to 100 K) is possibly describedby a hybrid process where thermally excited charge carriers can tunnelthrough the potential barriers into the LUMO level of C60. From theelectrical measurements, direct tunneling is predicted for thin C60layers of maximum ∼4 nm.Impedance analysis shows that C60 is fully depleted and thus thatthe voltage drops linearly over the entire structure. The static dielectricconstant of C60 extracted from the geometrical capacitance is ϵFinally, this study shows thatCVD graphene can be used as a topelectrode in C60 thin film vertical devices. The fabrication processcan potentially be applied to other multilayer hybrid van der Waalsheterostructures where the charge transport across neat Gr/OSC interfacesdefine the functionality of the devices. For instance, tunneling mechanismscould be exploited in high-frequency devices such as organic hot electrontransistors, permeable-base transistors, and organic flexible memories.48C60(99.9%) powder was purchased from Sigma-Aldrichand evaporated without further treatments. CVD graphene was grownon copper foils using an in-house automated setup and following thegrowth protocol which can be found elsewhere.2 (335nm) substrate by photolithography and under ambientconditions, as described in the Supporting Information. The chip includes two distinct groups of devices: (i) 119 Au/C60/Gr Vertical Stacks . The chip overview is given in Table S3. In both Vertical Stacks and Graphene Bridges, the C60 thin film is interposedbetween a bottom gold and a top graphene circular electrodes. In the Graphene Bridge architecture, graphene is laterally contactedin order to measure its resistance . The chip includes deviceshaving graphene electrodes of various nominal diameter, i.e., 5, 10,15, 20, 25, 30, and 50 μm, while the Au bottom electrodes are2 μm larger. In short, Au electrodes were prepatterned and depositedby e-beam physical vapor deposition (EBPVD) and lift-off. Then, theC60 thin film was thermally evaporated  and patterned by lift-off. Finally, the graphenesheet was wet transferred on the chip and patterned into the top electrodecircular shape by reactive ion etching (RIE).Devices were fabricated on a Si(525 μm)/SiO–6 mbar)in the dark, using the Keithley 4200 semiconductor parameter analyzer.The I–V characteristics ofthe Vertical Stacks were measured in the −10to +10 V voltage range, with sweep rate of ca. 1 V/s, steps of 50mV, and internal averaging of 20 ms. The Au electrode was connectedto ground. The graphene resistance was measured in GrapheneBridge devices by sweeping the voltage in the −50to +50 mV voltage range, with sweep rate of ca. 20 mV/s, steps of1 mV, and internal averaging of 20 ms.The electrical measurementswere performed under various conditions, i.e., in air, at room temperaturein the dark, and under vacuum , under vacuum (∼10–6 mbar) andin the dark. The electronics consisted of an AdWin Gold II ADC-DACand a Femto DDPCA-300 current-to-voltage converter. MATLAB scriptswere used for the data acquisition. The I–V characteristics of the Vertical Stacks were measured in the −10 to +10 V voltage range, with sweeprate of ca. 150 mV/s, steps of 10 mV, and with internal averagingof 60 ms. For the charge transport analysis, the backward and forward I–V sweeps for positive and negativebiases were considered, respectively. The full I–V sweeps (forward and backward) for representative devicesmeasured at room temperature (293 K) and at low temperature (50 K)can be found in Figure S12.The temperature-dependent –6 mbar), andthe oscillator level was set to 100 mV in the in the 40 Hz–1MHz frequency range; 201 data points were acquired. The open/shortcompensation method was applied to the measurements after the dataacquisition and following the Agilent impedance measurement handbook.49 Devices for the open/short compensation wereembedded on the chip for this scope.An Agilent4294a precision impedance analyzer was used to measurethe impedance of three representative devices per area. Python scriptswere used to control the instrument. The measurement was performedin the dark, under vacuum (∼10–1). Raman spectra were collected witha 532 nm excitation wavelength under ambient conditions. The laserpower and integration time of 0.1 mW and 120 s were set for the device,while graphene spectrum on SiO2 was acquired using a laserpower and integration time of 1 mW and 60 s.A WITec Alpha 300R confocal Ramanmicroscope was used with a LD 100× objective  and a 300 mm lens-based spectrometer  and then through ion-induceddeposition . The cross section was then cut in a 30kV gallium ion beam at an ion current of 47 nA. After cutting, thecross section was polished in sequential steps, decreasing the ioncurrent down to a minimum of 790 pA.50 for the FN modeling and for the extraction ofthe graphene resistance, (ii) lmfit51 modelfor the DSB modeling, and (iii) impedance.py52 for the impedance spectroscopy analysis.Python scripts weredeveloped for the modeling, fitting and visualization of the data.The main libraries that were implemented are (i) scipy curve_fit