These findings represent a major step forward toward the practical application of nanomaterials in civil engineering.This study reports on the development of a cementitious composite incorporating electrochemically exfoliated graphene (EEG). This hybrid functional material features significantly enhanced microstructure and mechanical properties, as well as unaffected workability; thus, it outperforms previously reported cementitious composites containing graphene derivatives. The manufacturing of the composite relies on a simple and efficient method that enables the uniform dispersion of EEG within cement matrix in the absence of surfactants. Different from graphene oxide, EEG is found to not agglomerate in cement alkaline environment, thereby not affecting the fluidity of cementitious composites. The addition of 0.05 wt% graphene content to ordinary Portland cement results in an increase up to 79%, 8%, and 9% for the tensile strength, compressive strength, and Young's modulus, respectively. Remarkably, it is found that the addition of EEG promotes the hydration reaction of both alite and belite, thus leading to the formation of a large fraction of 3CaO·2SiO Therefore, cement composites incorporating graphene oxide (GO), i.e., the most easily processable graphene derivative, have been intensively investigated in recent years.The cement–GO composites have emerged as materials with improved mechanical properties Table . Even in2+ in calcium hydroxide.Mechanical properties of cementitious composites should be assessed by means of their microstructure and composition of cement paste hydrates. As demonstrated in previous studies, the reinforcing mechanism of GO is attributed to the chemical reaction between GO nanosheets and cement hydration products.While the potential application of graphene oxide in cement composites has been extensively studied in recent years, similar research involving graphene, to date, remained limited. Some attempts have been made to fabricate cement composites with the addition of graphene nanoplatelets.Here we present an unprecedented cementitious composite incorporating few‐layer thick graphene nanosheets. We have devised a technologically simple yet efficient method for manufacturing a cement–graphene composite based on the use of graphene obtained via electrochemical exfoliation (EE) of graphite foil. Prior to being mixed with cement, the graphene aqueous suspension is dried and wiped through the set of sieves. Contrary to previous reports describing the generation of cementitious materials incorporating graphene derivatives, the preparation of our composite does not require special treatments or the use of surfactant to obtain uniform dispersion of graphene within cement matrix. We first evaluated the dispersion of graphene in cement alkaline environment and its effect on the consistency of cement mortar. Then we investigated the mechanical properties of produced composites to select the most advantageous content of graphene. The 0.05 wt% addition of graphene results in the highest increase of mechanical properties, namely, 79%, 8%, and 9% for tensile strength, compressive strength, and Young's modulus, respectively. Finally, we characterized the microstructure and composition of our mortars. As revealed by detailed characterization, the hydration rate of calcium silicates is remarkably increased in cement–graphene composites. The addition of electrochemically exfoliated graphene (EEG) accelerates the nucleation and growth of C‐S‐H phase, thus resulting in significant improvement of composites' strength. Subsequently, the fluidity of the processed composites remains unaltered compared to reference mortars, indicating that the concrete alkaline environment does not jeopardize the uniform dispersion of graphene within cement matrix. These results open new possibilities for the practical applications of graphene in high‐performance cementitious composites.22.1EEG has been prepared under anodic conditions by using a simple electrolytic cell.2 substrate by spin coating 200 µL of 1 mg mL−1 dispersion in dimethylformamide. FigureAtomic force microscopy (AFM) is employed to gain in‐depth insight into the surface morphology of EEG flakes. Toward this end, EEG flakes were deposited on SiO2.2FigureThe chemical composition of the as prepared EEG flakes was unraveled by X‐ray photoelectron spectroscopy (XPS) and compared with the starting material, i.e., the graphite foil, as well as with commercially available graphene oxide. Such analysis was performed aiming at following the changes in the chemical composition of the material during EE. 2.32+, Na+, and K+ and anions such as OH− and (SO4)2−. Ca2+ and OH− ions originate from calcium hydroxide, Na+ and K+ ions exist in sodium and potassium oxides present in cement, while sulfate ions can originate from gypsum, mixing water, or aggregate.2, NaOH, KOH, NaCl, CaCl2, and NH3 x H2O2, NaOH, and KOH.+ and K+ ions in the solution with pH value below 13.2+ ions. Toward this end, we added 100 µL of saturated Ca(OH)2 solution to 3.0 mL of aqueous dispersion of EEG or GO dispersion with concentration of 0.07 mg mL−1. Indeed, the addition of Ca(OH)2 results in rapid and notable flocculation of GO, conversely not affecting graphene dispersion .We first investigated the dispersion of EEG and GO in alkaline environment and their influence on the consistency of cement mortar. The solution present in concrete pores is highly alkaline due to the presence of cations, including CaFigureWe then attempted to manufacture several cement mortars made out of Portland cement Type I with various loadings of EEG and GO ranging from 0 wt% to 0.1 wt%. Noteworthy, the water‐to‐cement and sand‐to‐cement ratios remained 0.5 and 3.0, respectively. Moreover, in order to provide extensive benchmarking for samples with EEG, we have produced additional CEM I mortar incorporating 0.05 wt% of graphite flakes (GFs). However, as in the case of EEG, we have reported negligible effect of the graphite flakes addition on the consistency of cement mortar . Given a similar amount of water required to obtain the standard consistency for both types of cement (see Table S1 in the Supporting Information), as well as the same depth of plunger penetration noted for reference samples in consistency test, this phenomenon should be explained in terms of chemical composition of CEM I and CEM II . CEM II modified with granulated blast furnace slag contains less CaO than CEM I. This fact, combined with our consistency measurements, provides evidence that chemical cross‐linking of GO by calcium ions may account for the reduction of cement mortar fluidity, as revealed by previous studies.2.4FigureWe then performed compressive and tensile strength tests of all produced mortars at the age of 28 d. In order to obtain the most accurate results not affected by the end restraint effect, we have investigated mechanical properties on cylindrical samples with the height‐to‐diameter ratio of 2 .Figure On the other hand, with the introduction of 0.05 wt% of GO, the compressive strength is considerably decreased by 11% Figure b. This fFigureFurthermore, since it is of crucial importance to improve concrete performance in tension, we have then determined the tensile strength of produced mortars by means of a direct tension test Figure5a. HerSignificantly, the remarkable tensile strength values we achieved cannot be directly compared with those reported in the literature,Clearly, since large pores have been monitored in cement–GO composites Figure f, the teSimultaneously, CEM II composites incorporating EEG demonstrate slightly enhanced tensile strength  and, surprisingly, decreased compressive strength . The effect of GO addition on CEM II mortars strength and Young's modulus is, however, comparable with that noted for CEM I .2.5In order to understand the origin of the improvement of mechanical properties, as well as the underlaying strengthening mechanism of EEG, we investigated the microstructure and composition of produced cement mortars using various characterization techniques, including thermogravimetric analysis (TGA), X‐ray powder diffraction (XRD), Fourier transform infrared spectroscopy (FTIR), scanning electron microscopy (SEM), and Brunauer–Emmett–Teller (BET) adsorption–desorption analysis.2, and calcium silicate hydrate, C‐S‐H phase. C‐S‐H amorphous phase, accounting for ≈50%–70% of cement structural components, is considered as the strongest phase in hardened cement paste and plays the major role in developing macroproperties of cementitious composites. Ca(OH)2 with its hexagonal plate‐shaped crystals, being a considerably weaker phase, is responsible for providing high alkalinity of concrete environment, thereby protecting steel reinforcement against corrosion. However, high amount of calcium hydroxide may contribute to leaching, carbonation, alkali aggregate reaction, or sulfate attack, thus deteriorating the durability of concrete structures.Fundamentally, mechanical properties of cementitious composites originate from microstructure and composition of cement paste hydrates, in particular, calcium hydroxide, Ca(OH)2 content is of significant importance in characterizing cementitious materials, we first performed thermogravimetric analysis to probe Ca(OH)2 amount in all produced mortars at the age of 28 d. TGA has been shown to be a powerful tool for estimating portlandite content in cement composites, considering the specimens weight losses in the dehydroxylation (≈400–550 °C) and decarbonation (≈600–800 °C) regions.mH2O, mCH, and mCO2 correspond to the molecular mass of H2O (18 g mol−1), Ca(OH)2 (74 g mol−1), and CO2 (44 g mol−1), respectively.Since the determination of C‐S‐H and Ca(OH)on ≈400–5 °C and dFigure2 content is drastically reduced, achieving the lowest value of 13.07% for CI‐G‐0.05 samples, while an average value for reference samples is 16.37%. It should be pointed out that the lowest portlandite content occurred in samples exhibiting the highest tensile strength. On the other hand, we also report the slight decrease of Ca(OH)2 amount in GO composites, i.e., from 16.37% to 15.17% for CI‐GO‐0.05 samples. Since these specimens possess the lowest compressive and tensile strength, these results indicate the opposite trend in Ca(OH)2 amount compared to graphene composites.2 presence in hardening cement matrix. First, at initial contact with water, cement grains start to dissolve, thus releasing, among others, calcium and hydroxide ions from their surface. At the end of dormant period, when the supersaturation point of Ca2+ ions is reached, Ca(OH)2 starts to precipitate from the saturated solution. Second, Ca(OH)2 is produced as a result of hydration reactions  and dicalcium silicate 2CaO·SiO2 (belite)To fully explain our observations, it is particular important to attain a deep understanding on the origin of Ca(OH)Equation  of two m2. Alite with its fast reaction and high release of heat is responsible for early hydration, thus contributing mainly to early‐age strength of concrete, whereas the reaction of belite being markedly slower releases less heat and contributes to long‐term strength.Noteworthy, although these two reactions result in the same hydration products, they differ considerably in hydration rate, heat, and the amount of produced Ca(OH)Since the formation of portlandite is a highly complex process extensively varying at the different stages of ongoing cement hydration, we have then performed XRD and FTIR analysis on samples cured for 28 d, in order to gain further insight onto the course of cement hydration reactions and their products. We have focused our attention on four representative samples made out of CEM I, namely, CI‐R, CI‐G‐0.05, CI‐GF‐0.05, and CI‐GO‐0.05.Θ = 26.6°.2 appear at 18.0° and 34.1°.2 amount in CI‐G‐0.05 samples as revealed previously by TGA. Basically, as C‐S‐H gel is an amorphous phase, it cannot be detected during XRD analysis. However, the hydration degree may be estimated on the basis of remaining peaks of alite and belite at the positions of 28.7° and 29.4°.Figure 2 at 3650 cm−1,3 at 874 and 1414 cm−1,−1 represents the stretching vibration of Si—O bonds of C‐S‐H gel revealing, in particular, higher formation of C‐S‐H phase in cement mortar reinforced with graphene and, simultaneously, restrained C‐S‐H growth in cement–GO composites.In the case of FTIR, the peaks attributed to —OH of Ca(OH)2 adsorption–desorption isotherms measurement at 77 K . CI‐R composite displays a specific surface area of ≈9 m2 g−1, while the specific surface area of CI‐GO and CI‐G is calculated with the BET model as 17 and 29 m2 g−1, respectively. Significantly, the results indicate the increased development of the highly porous phase, i.e., the C‐S‐H gel, resulting from the addition of EEG. Moreover, the average pore diameter calculated with the Barrett–Joyner–Halenda model for CI‐R, CI‐GO, and CI‐G amounts to ≈3.9 nm .Interestingly, the evaluation of the specific surface area of cementitious composites has been shown being a powerful tool for assessing the C‐S‐H phase development, which is strongly associated with the degree of cement hydration.2 formation at different stages of cement hydration n Figure d. In thiIn the case of GO, Figure Significantly, we also observed that no major differences occurred in the diffraction patterns and FTIR spectra of reference samples and samples with graphite flakes, pointing toward similar mineralogical compositions of both materials. This provides unambiguous proof that the addition of graphite flakes, in contrast to EEG, does not affect the microstructure and hydration products of cement composites, acting only as dispersed reinforcement.Figure2 can be easily observed in plain, GO, and graphite composites, while they do not occur in graphene mortar. The visible densification of microstructure provides a rational explanation of the remarkable performance of cement‐EEG mortars in mechanical properties tests.Our observations are further confirmed by SEM images . As in CEM I composites, the lowest value of Ca(OH)2 occurred in samples with the highest tensile strength. Nevertheless, the value changes only from 11.21% in reference samples to 10.55% in CII‐G‐0.05 specimens. Interestingly, the addition of GO results in intense formation of calcium hydroxide, achieving a value as high as 13.8% for CII‐GO‐0.05 samples. These observations are also supported by XRD  and FTIR  analysis. Moreover, XRD patterns reveal a slight decrease in alite and belite amount for samples with graphene, while FTIR spectra show a similar trend in C‐S‐H phase formation as in CEM I.In addition to our findings in terms of Portland cement Type I, structural characterization of CEM II composites was performed using the same techniques previously described. The Ca(OH)2 content after 1, 3, and 7 d of cement hydration. In Portland cement with the addition of granulated blast furnace slag, Portland cement components start to hydrate first. At this initial stage, the hydration of CEM II strongly resembles that of ordinary Portland cement.2 content for reference sample is noted, a fingerprint of the presence of granulated blast furnace slag. Basically, when Portland clinker reactions are in progress, granulated blast furnace slag is then activated by alkalis and starts to react with calcium hydroxide released by Portland cement. As a consequence, additional amount of C‐S‐H phase is formed.Particular attention should be paid to Ca(OH)3In summary, we have demonstrated a novel cementitious nanocomposite incorporating graphene that features remarkably enhanced mechanical properties and microstructure. We first established a technologically simple and efficient method for producing a cement–graphene composite, which can be easily employed in industrial scale applications. At the core of our approach is graphene obtained by electrochemical exfoliation of graphite, dried and wiped through the set of sieves, prior to mixing with cement. The preparation method described allows manufacturing a nanocomposite without the use of surfactants or any special treatment to obtain the homogenous dispersion of graphene within cement matrix. In this regard, our composite significantly outperforms previously reported composites with graphene derivatives. Electrochemically exfoliated graphene proved not to aggregate in alkaline environment and thereby does not reduce the fluidity and workability of cement mortars. We then investigated the mechanical properties of produced composites. Remarkably, the addition of 0.05 wt% of graphene to ordinary Portland cement results in significant enhancement of tensile strength up to 79% and, simultaneously, slight increase of compressive strength and Young's modulus by 8% and 9%, respectively. As revealed by further structural characterization, graphene promotes the hydration reactions of calcium silicates, thus resulting in much more intense formation of C‐S‐H phase, as well as regular and compact microstructure. Overall, such findings provide unambiguous evidence that our composite can boost the practical application of graphene in concrete technology. Moreover, the improved performance of cementitious composites incorporating graphene, in particular the significantly enhanced tensile strength, will allow designing lighter concrete structures with extended durability. Therefore, the consumption of concrete components may be reduced, thus mitigating the environmentally harmful impacts of concrete production.4Materials: Cement, sand, distilled water, graphene, graphene oxide, and graphite flakes were used in this study to fabricate cement mortar. In particular, two types of cement were employed, which were provided by Gorazdze Cement S.A.: CEM I 42.5R (Portland cement Type I) and CEM II/B‐S 32.5R‐NA (Portland cement Type II with granulated blast furnace slag). The chemical composition of the cement is presented in Table S2 (Supporting Information). The mechanical and other physical properties of the cement are reported in Table S1 (Supporting Information). The results of sand sieve analysis are plotted in Figure S11 (Supporting Information). Dispersion of graphene oxide sheets with concentration of 4 mg mL−1 was purchased from Graphenea Inc. Graphite flakes were purchased from Sigma‐Aldrich. Graphene was produced by electrochemical exfoliation of graphite foil  using a platinum wire  and aqueous solution of ammonium sulfate (NH4)2SO4 (Sigma‐Aldrich). To stop cement hydration, acetone (Carlo Erba Reagents) was used in a process of solvent replacement and liquid nitrogen in a process of freeze‐drying. Dispersions of graphene and graphene oxide in alkaline environment were investigated by using saturated aqueous solution of calcium hydroxide Ca(OH)2 (Sigma‐Aldrich).Electrochemical Exfoliation of Graphite Foil: Graphite foil cut into pieces of 2.5 cm x 6 cm was used as anode and platinum wire was used as cathode in an electrolytic cell. Both elements of the electrolytic cell were immersed in aqueous solution of (NH4)2SO4 with concentration of 0.1 m. A starting current of 0.4 A generated by the applied voltage of 15 V (ISO‐TECH IPS‐603 DC power supply) prompted the exfoliation of graphite foil. The exfoliation was conducted for ≈3 h, i.e., until the graphite foil was completely exfoliated. The resulting material was then rinsed several times and filtrated via polytetrafluoroethylene (PTFE) membranes (pores diameter of 5 µm). Finally, collected materials were dried at 60 °C for at least 48 h.Cement Mortar Preparation: Cement mortar was prepared by mixing cement, sand, water, and graphene/graphene oxide/graphite flakes. The water‐to‐cement ratio was kept at 0.5. The sand‐to‐cement ratio was also kept constant in all mixes and it was equal to 3.0. A plain mortar (labeled as R), five mixtures with EEG (labeled as G) and three mixtures containing GO were prepared for each type of cement (CI or CII). Graphene was incorporated in the cement mortar at five different percentages, namely, 0.01 wt%, 0.03 wt%, 0.05 wt%, 0.075 wt%, and 0.1 wt% (dosage by weight of cement). The content of graphene oxide was 0.01 wt%, 0.03 wt%, and 0.05 wt%. Cement mortar with a higher dosage of graphene oxide was not produced due to the reduction of fluidity and setting time, thus resulting in poor workability. Additionally, a reference mixture with 0.05 wt% dosage of GFs was fabricated using CEM I.For cement mortar containing graphene, dried graphene was wiped through the set of five sieves with the finest mesh sieve size being 250 µm. Graphene and cement were stirred sufficiently at low speed (≈140 rpm) using a hand mixer to obtain a homogenous dry mixture. Mixing procedure from PN‐EN 196‐1:2005 was then adopted. Water was added to cement with graphene and the mixer was immediately started at low speed for 30 s. After that, sand was steadily added during the next 30 s and the mixer was switched to high speed (≈285 rpm) for additional 30 s. Afterward, the mixer was stopped to remove all the mortar adhering the walls of the bowl and then mixing was continued at high speed for 60 s. For samples with graphite, graphite flakes were also added to cement prior to pouring water, whereas for composites with graphene oxide, the GO dispersion was added to cement simultaneously with water. All resulting cement composites were placed into steel cylindrical molds with the diameter of 60 mm and the height of 120 mm. Cement mortar was placed in the molds in a few layers and each layer was subjected to vibration on a vibration table for 1 min to ensure the compaction of the composite. All specimens were immediately covered by polyethylene foil to prevent loss of water. After 24 h, the hardened cement mortar samples were demolded and continued to be cured in water at 20 °C. All cylindrical samples were dried in the air for 24 h before performing mechanical tests.Characterization and Measurements: In order to investigate the dispersion of graphene and graphene oxide in alkaline environment, 3.0 mL of graphene and graphene oxide aqueous dispersions with concentration of 0.07 mg mL−1 were first prepared. Then 100 µL of saturated aqueous Ca(OH)2 solution was introduced into both dispersions.The consistency of fresh mortars was determined according to PN‐EN 1015‐4 by the plunger penetration method. The sample was placed in the vessel in two layers and each layer was subjected to vibration on a vibration table. The vessel was then put on the base plate under the plunger. The plunger was allowed to freely fall from the height of 100 mm above the vessel. The depth of the plunger penetration was determined as a consistency measure. Two measurements were performed for each cement mortar and the average was taken.−1. The compressive strength tests were performed using MEGA 3‐3000‐100 compressive strength tester (Form+Test Prüfsysteme) at a loading rate of 0.5 kN s−1. Five samples of cement mortar were prepared for each test and the average was taken. To obtain the compression stress–strain curve, linear polyester strain gauges with gauge factor of 2.13  were employed. Two pairs of strain gauges were attached to the two sides of the sample to measure axial and transverse strain. On the basis of the strain gauges measurements, the modulus of elasticity was calculated. According to PN‐EN 1992‐1‐1:2008 the Young's modulus was determined as a secant value between zero stresses and stresses equal to 40% of the average compressive strength.Mechanical properties tests were conducted on cylindrical samples  at the age of 28 d. The tensile strength was determined using ZD 10/90 tensile strength tester (Heckert) at a loading rate of 1.0 kN sInhibiting the cement hydration is essential to prepare cement composites' samples for microstructure and composition characterization by means of TGA, SEM, XRD, FTIR, or BET analyses. Two different methods for stopping the cement hydration in order to obtain the most satisfactory and precise results were employed: freeze‐drying to effectively preserve the composition of cement mortar for TGA, XRD, and FTIR analyses and solvent replacement method to avoid damage of pores and alternations to the cement mortar microstructure for SEM and BET.The microstructure of the fracture surface of hardened cement mortar was investigated by SEM using FEI Dual Beam 235 with the accelerating voltage of 5 keV incident beam energy. The pieces of crushed cement composites after the solvent replacement process were glued to a support with conductive carbon adhesive and the fracture surface was sputter‐coated with a thin layer of gold prior to SEM imaging.For TGA, XRD, and FTIR analyses, the crushed pieces of cement composites were, initially, grounded completely into fine powder and filtered through a 250 µm sieve to remove coarse grains of sand.−1. All experiments were performed under air atmosphere.TGA was performed using Mettler Toledo TGA/DSC 2 sensor using alumina crucibles. Three samples of ≈15 mg of each cement mortar were, first, kept isothermally at 30 °C for 30 min and then heated from 30 to 1000 °C at a heating rate of 10 °C minXRD was carried out on Bruker ASX D8 Advanced with Cu anode with Kα radiation (λ = 1.5418 Å). Diffraction patterns were collected at room temperature in the scattered angular range between 6° and 60°. An angular resolution and a typical counting time per each step were of 0.02° and 10 s, respectively.−1) by using a Perkin Elmer Spectrometer (Spectrum Two) equipped with ATR Diamond.FTIR spectra were obtained within the mid‐IR range (400–4000 cm−9 mbar and an Al anode as the X‐ray source (X‐ray radiation of 1486 eV). Spot size was of 400 µm. Pass energies of 200.00 eV for wide energy scans and 10.00–20.00 eV for scans were used.XPS analyses were performed on a photoelectron spectrometer with a basic chamber pressure of −10The morphology of electrochemically exfoliated graphene was investigated using a Veeco Dimension 3100 AFM with a Nanoscope IV control unit under ambient condition. Topographic imaging was carried out in tapping mode with the use of antimony (n) doped silicon cantilever.The specific surface area was measured using a Micromeritics ASAP 2050 surface area and porosity analyzer. Prior to the BET measurements, the samples were outgassed for 10 h at 100 °C. Adsorption isotherms were calculated for nitrogen adsorption at 77 K and pressures up to 1 bar.The authors declare no conflict of interest.SupplementaryClick here for additional data file.


Polarization-insensitive modulation, i.e., overcoming the limit of conventional modulators operating under only a single-polarization state, is desirable for high-capacity on-chip optical interconnects. Here, we propose a hybrid graphene-silicon-based polarization-insensitive electro-absorption modulator (EAM) with high-modulation efficiency and ultra-broad bandwidth. The hybrid graphene-silicon waveguide is formed by leveraging multi-deposited and multi-transferred methods to enable light interaction with graphene layers in its intense field distribution region instead of the commonly used weak cladding region, thus resulting in enhanced light–graphene interaction. By optimizing the dimensions of all hybrid graphene-silicon waveguide layers, polarization-insensitive modulation is achieved with a modulation efficiency (ME) of ~1.11 dB/µm for both polarizations (ME discrepancy < 0.006 dB/µm), which outperforms that of previous reports. Based on this excellent modulation performance, we designed a hybrid graphene-silicon-based EAM with a length of only 20 µm. The modulation depth (MD) and insertion loss obtained were higher than 22 dB and lower than 0.23 dB at 1.55 µm, respectively, for both polarizations. Meanwhile, its allowable bandwidth can exceed 300 nm by keeping MD more than 20 dB and MD discrepancy less than 2 dB, simultaneously, and its electrical properties were also analyzed. Therefore, the proposed device can be applied in on-chip optical interconnects. Silicon photonics based on the mature silicon-on-insulator (SOI) platform has aroused tremendous interest for building compact, high performance, and energy-efficient photonic integrated circuits (PICs), which has also promoted the rapid development of optical interconnects, especially for data center and telecom applications ,2,3,4. T2·V−1·s−1 at room temperature) , imaginary [Im(neff)]) of the hybrid graphene-silicon waveguide as a function of the graphene chemical potential (µc) for both transverse electric (TE) and transverse magnetic (TM) modes at λ = 1.55 µm where the modal calculation is based on the above-mentioned surface conductivity model of graphene. The calculation engine is available from a commercial software package  where the MD of the TM mode has a peak value at ΔH = −20 nm. Moreover, the polarization-insensitive property of the proposed EAM remains unaffected if ΔW and ΔH are kept within the ranges of  and , respectively, which gives a relatively large tolerance for fabrication. For the device length L, the variations of transmittance and MD of the TE and TM modes are nearly the same, showing excellent polarization-insensitive properties. In addition, a higher MD for this device can be achieved only by increasing the device length because its insertion loss is very low and increases slowly owing to the small absorption loss at the “ON” state for both polarizations. Finally, Next, to investigate the performance deterioration brought by fabrication imperfections in practice, we analyzed the fabrication tolerances of the width f3dB) and the energy consumption (E), which are associated with the total capacitance (C) and resistance (R) of the optical device, f3dB = 1/(2πRC), E = CV2/4 (V stands for the switching voltage between “OFF” and “ON” states) [C) is estimated to be ~326 fF. Thus, f3dB and E are derived to be ~6.1 GHz and ~7.8 pJ/bit, respectively. To further enhance the electrical properties of the proposed EAM, its total resistance and capacitance should be effectively reduced, especially for the capacitance, which can be realized by cutting down the overlapping area between two corresponding graphene layers, choosing the spacer/insulating layer with the lower dielectric constant, and increasing the thickness of the spacer/insulating layer. In addition, traveling-wave electrode designs could also be leveraged to enhance the modulation performance of the proposed EAM [Besides the above-analyzed optical properties of the proposed hybrid graphene-silicon-based polarization-insensitive EAM, its corresponding electrical characteristics are also critical, such as the 3 dB modulation bandwidth ( states) ,42. Here states) ,28,29. Bosed EAM .In summary, using multi-deposited and multi-transferred methods, we propose a highly-efficient, ultra-broadband, and polarization-insensitive EAM, the key part of which is a hybrid graphene-silicon waveguide formed by five inverted U-shaped graphene-silicon layers. In order to effectively enhance LGI, multi-deposited and multi-transferred graphene-silicon layers are employed such that light interacts with the graphene layers at high optical intensity in the interior of the hybrid graphene-silicon waveguide rather than at weak optical intensity in the cladding region, as commonly done. We then optimized the dimensions of each hybrid graphene-silicon layer based on the inverted U-shaped waveguide structure and achieve an ME of ~1.11 dB/µm and an ME discrepancy of lower than 0.006 dB/µm for both polarizations, which are significant improvements over previous results. By using this structure, we designed and analyzed a hybrid graphene-silicon-based polarization-insensitive EAM only 20 µm long. The MD and IL were calculated to be > 22 dB and < 0.23 dB at 1.55 µm, respectively, and the available bandwidth can be over 300 nm (from 1.367 to 1.668 µm) by simultaneously keeping MD greater than 20 dB and MD discrepancy less than 2 dB. The MD of the designed EAM can be further improved by increasing its length. The electrical properties of the EAM  were studied. Several methods were also proposed to enhance its electrical features. With high ME, ultra-broad bandwidth, polarization-insensitivity, compact footprint, and scalable performance, the proposed hybrid graphene-silicon-based EAM will find applications in the high performance hybrid silicon modulators for high capacity on-chip optical interconnects.


During the past decades, major advances have been made in both the generation and detection of infrared light; however, its efficient wavefront manipulation and information processing still encounter great challenges. Efficient and fast optoelectronic modulators and spatial light modulators are required for mid-infrared imaging, sensing, security screening, communication and navigation, to name a few. However, their development remains elusive, and prevailing methods reported so far have suffered from drawbacks that significantly limit their practical applications. In this study, by leveraging graphene and metasurfaces, we demonstrate a high-performance free-space mid-infrared modulator operating at gigahertz speeds, low gate voltage and room temperature. We further pixelate the hybrid graphene metasurface to form a prototype spatial light modulator for high frame rate single-pixel imaging, suggesting orders of magnitude improvement over conventional liquid crystal or micromirror-based spatial light modulators. This work opens up the possibility of exploring wavefront engineering for infrared technologies for which fast temporal and spatial modulations are indispensable. 2O3, an amorphous-Si layer and a rear reflector made of gold. The structure also features gate and drain electrodes to electrically tune the device. When constructed into a 6x6 pixel array the modulators allow the realization of a SLM for proof-of-principle high frame rate single-pixel imaging in the mid-IR.Ultrafast modulators and spatial light modulators (SLMs) for mid-infrared light have been realized by the use of an electrically-tunable graphene on a silicon-integrated metasurface. The modulators designed and built by Beibei Zeng and coworkers from Los Alamos National Laboratory, Duke University and Arizona State University in the US offer a depth of modulation of 90% and a modulation speed exceeding 1 GHz. The modulator consists of a sandwich structure consisting of an array of gold nanoantennas, a layer of graphene, atop a thin dielectric film of Al For example, solid-state infrared light modulators based on semiconductor quantum wells either have a small modulation depth or require sophisticated material growth and cryogenic temperatures7. In addition, state-of-the-art liquid crystal8 and micromirror9-based spatial light modulators (SLMs) suffer from limitations including slow modulation speeds of a few kilohertz and complex and expensive instrumentation. In this context, active metamaterials/metasurfaces composed of planar subwavelength resonators and reconfigurable functional materials/devices have provided an alternative approach10. In the microwave region, the integration of electronic devices  has resulted in reprogrammable coding metasurfaces13. In the far-infrared region, the integration of metasurfaces and compound semiconductors  has enabled the modulation of terahertz (THz) waves with high modulation depth and speed at room temperature17. However, such configurations are difficult to scale up to the mid- and near-infrared regions, where electronic devices fail to operate and the optical conductivity of conventional semiconductors can only be slightly adjusted19, resulting in limited dynamic modulations.Emerging infrared technologies lie at the core of advanced photonic research because of their great potential for numerous applications. Despite the tremendous efforts in the development of sources23. In addition, the compatibility of graphene with CMOS processing makes it an extremely promising material for cost-effective optoelectronics in high-frequency applications25. Most recently, hybrid graphene metasurfaces have shown unprecedented capabilities in developing electrically reconfigurable infrared optoelectronic devices47, although in these designs the modulation relies on a high voltage bias and the modulation speed is limited, and many studies have focused on a waveguide configuration31. In this paper, we present the experimental demonstration of hybrid graphene metasurface free-space mid-infrared modulators that enable a large intensity modulation depth of up to 90% and a high modulation speed exceeding 1 GHz over a broad bandwidth by tuning the Fermi level of graphene at a low gate voltage bias of ∼7 V. The low-voltage operation achieved by the design architecture of our device is essential for obtaining a high-quality modulation signal and an ultra-high modulation speed under ambient conditions. Furthermore, a prototype mid-infrared SLM is demonstrated to function as an electrically encoded aperture or mask, realizing single-pixel imaging with a high frame rate.In contrast, graphene has demonstrated exceptional electrical and optical properties and a widely tunable electro-optical response through the entire infrared region47, which also limits the modulation speed and the scope of applications. Operating at mid-infrared wavelengths, the high gate voltage is due to the thick dielectric spacer  between the metal ground plane and metallic resonator array32. When graphene is integrated into the resonator array, the metamaterial absorber is essentially a field-effect transistor (FET) structure in which the application of a gate voltage Vg tunes the graphene conductivity by varying the Fermi level μC according to33μC= ħv(πCVg/e)1/2. To lower the gate voltage, it is necessary to increase the capacitance C by decreasing the gate dielectric thickness . Although we cannot reduce the dielectric spacer thickness of a metamaterial absorber, an alternative solution is to replace a large portion of the dielectric spacer with a material that is conducting for the applied voltage bias but behaves as a dielectric in the mid-infrared range. This strategy is demonstrated by using the FET structure schematically shown in Fig. 2O3 layer to increase the capacitance. In this way, we are able to significantly reduce the gate voltage to a few volts to reach the graphene charge neutrality point and effectively tune the graphene conductivity, as shown in Fig. Ids and resistance Rds as functions of gate voltage Vg for a few different drain-source biases Vds.A major issue in frequency-tunable hybrid graphene metamaterial absorbers is the requirement of a high voltage bias34. That is, the incident light is trapped within the metamaterial cavity and eventually absorbed by the metallic structure, the lossy dielectric spacer, or both. Because the optical loss of a-Si is negligible in the mid-infrared region, the absorption mainly occurs via ohmic dissipation at the metallic resonator array integrated with graphene. These two facts give rise to the possibility of tuning the resonant dispersion of the nanoantenna array with the goal of satisfying the phase condition and thereby the absorption at a different wavelength by tuning the conductivity of the integrated graphene through the application of a gate voltage. In Fig. λ = 7.3 μm (reflection minimum ∼16%) to 8.3 μm (reflection minimum ∼5%) when the low gate voltage changes from +7 V to −3 V. The corresponding change of reflection ∆R(λ) = R+7 V(λ) − R−3 V(λ) and modulation depth M = |∆R(λ)|/max are plotted in Fig. λ = 8.5 μm and a modulation depth as high as ∼90% in this wavelength range.Following this strategy, our hybrid graphene metasurface structure is schematically shown in Fig. 36. The gate-tunable plasmonic modes are clearly shown in Fig. λ = 8.3 μm, corresponding to the reflection minimum in Fig. λ = 7.3 μm, as shown in Fig. λ = 7.3 μm, as shown in Fig. λ = 8.3 μm, as shown in Fig. The shifting of the plasmonic resonance dispersion can be visualized through near-field imaging of localized plasmonic modes in the hybrid graphene metasurface with different gate voltages15 with a resistance of either R0 = 2 kΩ or 4 kΩ. The cyan circles and green triangles in Fig. R0 = 2 kΩ and 4 kΩ external resistors, respectively, from which we can clearly observe the different modulation cut-off frequencies. Using the high-frequency circuit model27, the experimental data are fitted with a set of circuit parameters. The results are shown as the solid curves. Based on the same set of device circuit parameters, the intrinsic modulation is shown as the red curve in Fig. In addition to the high modulation depth, the hybrid graphene metasurface modulator enables an ultra-high modulation speed. Our characterization system, schematically shown in the inset to Fig. λ = 8.3 μm. By simultaneously turning on the functional pixels in the corresponding mask patterns, the red curve in Fig. λ = 7 μm and 8.5 μm due to their excellent modulation depth, in contrast to the lower imaging contrast at λ = 5.5 μm and 9.5 μm due to their reduced modulation depth  in the spacer of the metamaterial absorber. For the mid-infrared light, the a-Si layer can be considered as a dielectric material with negligible loss, but the semiconducting properties make it suitable to serve as part of the electrode for applying a gate voltage bias to tune the graphene conductivity, resulting in shifted/damped resonant dispersion of the nanoantenna array and thus frequency tuning of the metasurface absorption. The high-speed modulation is in part due to the low voltage operation and the significant reduction of parasitic capacitance from the electrodes.38. This work opens up the possibility of exploring wavefront engineering for infrared technologies for which fast temporal and spatial modulations are indispensable, and the demonstrated high frame-rate imaging approach holds great potential for many applications, such as the detection of transient thermal phenomena, real-time thermal imaging, security screening, navigation, and medical detection.The hybrid graphene metasurface modulator was then formed into a pixel array to serve as a prototype mid-infrared SLM. We showed that single-pixel imaging can be achieved by creating a series of spatial mask patterns, revealing orders of magnitude faster imaging speeds than those obtained using traditional liquid crystal and micromirror spatial light modulators. The speed of single-pixel imaging can be further enhanced by adopting a compressive single-pixel imaging algorithm39. We would like to emphasize that the prototype mid-infrared SLM can be scaled to operate at other frequencies, including terahertz, near-infrared, and even visible, since graphene has been proven to maintain the high optical conductivity and have opto-electronic responses at these frequencies due to the intra-band (terahertz and mid-infrared) and inter-band (near-infrared and visible) transitions42. The basic principle is the same, but when approaching shorter wavelengths, the concern is the increasing absorption in the amorphous silicon spacer layer, which can potentially reduce the modulation depth. Additionally, when scaling to operate at shorter wavelengths, the further reduced critical dimension of the metasurface structure and back gate array can potentially create fabrication challenges.We note that the optics in Fig. 2O3 gate dielectric grown by atomic layer deposition (ALD), two photolithography-defined 150 nm-thick gold strips serving as source and drain electrodes, and a transferred monolayer graphene grown by low-pressure chemical vapor deposition (LPCVD). Figure 43. Compared to the FET structure shown in Fig. 2O3 insulating layer and 3 nm/30 nm Ti/Au layers and a lift-off process. Graphene was then transferred and patterned to align with the electrode array, with graphene between electrodes removed by photolithography and O2 plasma etching, defining electrically isolated functional pixels as depicted by the optical micrograph in Fig. Figure 45. The copper foil was first loaded in a quartz tube and annealed at 1000 °C for 1 h while maintaining the H2/Ar pressure at 1 Torr under a 200 sccm flow. Graphene was then grown for 30 min by introducing 50 sccm CH4 at a total pressure of 1.7 Torr. The furnace was then rapidly cooled to room temperature. The growth conditions outlined above produced a large-area, continuous graphene film with very few multi-layer regions or cracks. To transfer the graphene layer to the substrate, PMMA (950 A4) was spin-coated onto one side of the graphene/copper/graphene foil and stored in vacuum overnight. The graphene on the backside was then removed with O2 plasma etching. The sample was then placed in an FeCl3-based etchant for 2 h, allowing the copper to completely dissolve. The remaining graphene/PMMA film was cleaned in HCl solution, rinsed with DI water, transferred onto the substrate, and put into acetone to remove the PMMA layer.Graphene was grown on a piece of copper foil  in a home-built LPCVD system36 equipped with a broadband infrared beam by difference frequency generation (DFG) to cover a wide mid-infrared range from 650 to 2400 cm−1. The raw data contain spectral information due to the use of broadband illumination, and we used Matlab programming to extract near-field images at specific wavelengths. The modulation speed was measured for individual pixels using the setup schematically shown in the inset to Fig. λ = 6.75 μm) was transmitted through a beam splitter (55/45) and focused onto the hybrid graphene metasurface modulator with a mid-infrared objective. The reflected light was collected by the same objective, reflected by the beam splitter, and focused into the TE-MCT detector using an off-axis parabolic mirror. A functional pixel of the device was applied with  a variable frequency sine voltage with offset 0 V and amplitude 1 V from a function generator integrated with an oscilloscope . The response of the TE-MCT detector was measured with the oscilloscope, determining the optical modulation signal of the device. Note that we were only able to measure the modulation frequency up to 1 GHz due to the limited frequency range of the function generator and TE-MCT detector and the low signal-to-noise ratio of our measurement setup. Reduction of the modulation cut-off frequency (at 3 dB) within the measurement range was achieved by inserting different series resistors  into the driving circuit15. The intrinsic modulation speed was then inferred using a high-frequency circuit model of the device. The spatial light modulation was characterized using the FTIR microscope by raster scanning the functional pixels—measuring the reflection intensity using a single-pixel detector as we turned only one pixel ‘ON’ or ‘OFF’ while all other pixels remained ‘OFF’. The imaging capability of the hybrid graphene metasurface SLM was demonstrated using the setup schematically shown in Fig. X) is projected onto the SLM through an infrared objective lens. The reflected light thus represents the image encoded by the SLM mask (forming the measurement matrix Φ) and is collected by the same objective and focused onto the MCT single-pixel detector (Y). For a proof-of-concept demonstration, a raster-scan measurement matrix Φ of the SLM mask was employed by sequentially turning on/off only one pixel for each measurement, resulting in a 36 × 36 identity matrix Φ. The original image X was then reconstructed by X = Φ−1Y, where Φ−1 is the inverse measurement matrix.Each pixel of the 6 × 6 pixel array was independently connected to one of the 36 channels  of a DAC evaluation board , which was then controlled by a microprocessor  using home-built Python codes. The reflection spectra of the individual pixels under different voltage biases were measured using a Fourier transform infrared spectroscopy microscope . The plasmonic resonance modes were visualized using a scattering-type near-field scanning optical microscope CE (CG) represents the capacitance between the top Au electrode (graphene) and the a-Si layer across the 750 nm-thick (6 nm-thick) Al2O3 insulating layer, RG and RS are the resistances of the graphene and a-Si, respectively, and R0 is the externally loaded resistance. The high-frequency modulation voltage from a function generator was applied to the device as the gate voltage Vg(ω), and the modulation voltage Vm(ω) at the functional pixel and the frequency-dependent modulation depth η(ω) of the device were derived asA simplified high-frequency circuit model of the device was developed, as shown in the inset to Fig.


The observed potential energy surfaces, representing the intermolecular interactions within the CH4/N2 gaseous mixtures and between these and the graphene layers, have been formulated by adopting the so-called Improved Lennard-Jones (ILJ) potential, which is far more accurate than the traditional Lennard-Jones potential. Previously derived ILJ force fields are used to perform extensive molecular dynamics simulations on graphene's ability to separate and adsorb the CH4/N2 mixture. Furthermore, the intramolecular interactions within graphene were explicitly considered since they are responsible for its flexibility and the consequent out-of-plane movements of the constituting carbon atoms. The effects on the adsorption capacity of graphene caused by introducing its flexibility in the simulations are assessed via comparison of different intramolecular force fields giving account of flexibility against a simplified less realistic model that considers graphene to be rigid. The accuracy of the potentials guarantees a quantitative description of the interactions and trustable results for the dynamics, as long as the appropriate set of intramolecular and intermolecular force fields is chosen. In particular it is shown that only if the flexibility of graphene is explicitly taken into account, a simple united-atom interaction potential can provide correct predictions. Conversely, when using an atomistic model, neglecting in the simulations the intrinsic flexibility of the graphene sheet has a minor effect. From a practical point of view, the global analysis of the whole set of results proves that these nanostructures are versatile materials competitive with other carbon-based adsorbing membranes suitable to cope with CH4 and N2 adsorption.We theoretically investigate graphene layers, proposing them as membranes of subnanometer size suitable for CH Graphene has often been investigated as a possible material for the separation of small gases  results as was explained in previous publications  from now on in this work  over the graphene sheet. The molecules were placed such that at least 5 Å were left in between the different molecules, the edges of the box and the graphene sheet to avoid instantaneous, strongly repulsive interactions at the start of the simulation. We then ran one NVE simulation to allow the system to relax to a physically viable conformation and used the resulting output as input for a first NVT simulation. After this NVT simulation, we kept the molecules that were adsorbed (see below) and deleted the ones that were not. Subsequently, we randomly distributed 100 new gas molecules (50 CH4 + 50 N2) over the remaining system from the previous simulation to run a new NVT simulation. This protocol was repeated until convergence of the amount of adsorbed molecules. As such, this protocol gives an indication of the amount of molecules that saturates the first adsorption layer and it allows comparison for the different intramolecular force fields and molecular models under study.Since the stronger affinity of graphene for methane than for nitrogen was clearly proven in this previous work, we expect that the graphene sheet could effectively separate this gas mixture. Therefore, in a first set of simulations we have adopted a protocol where we randomly allocated 100 gas molecules , while this fraction lowers in subsequent simulations. In the first simulations, the amount of molecules that is introduced in the system is not enough to fully saturate the graphene sheet and thus most molecules stay will adsorb onto the sheet. When more molecules are gradually introduced, the graphene sheet gets quickly saturated and more molecules will be forced to stay in gas phase in equilibrium with the adsorbed molecules. A closer inspection of the results for methane and nitrogen within the mixture reveals a changing methane/nitrogen ratio from one simulation to the next. Indeed, it can be seen that in the first simulations, a substantial amount of nitrogen molecules is adsorbed on the graphene sheet , while in subsequent simulations, they are steadily removed from the adsorbed layer. The methane molecules are adsorbed preferentially, but in the first simulations there are not enough molecules to completely saturate the graphene sheet allowing nitrogen molecules to adsorb as well. When more methane molecules are added, they occupy more and more space on the graphene surface allowing less space to nitrogen, which is forced to stay in gas phase. In the end, a situation is reached where about 70% of the molecules in the simulation are methane molecules and 30% are nitrogen, consistent over the four graphene sheets. Adding to this, the fraction of methane molecules that adsorbs is much larger than the fraction of nitrogen molecules, giving a first indication that indeed the graphene sheet is well capable of separating methane and nitrogen from each other.4/N2 mixture, this results is no surprise. Dreisbach et al.  is lower than when it is explicitly included (about 65%). As methane is dominantly present in the mixture, this result is clearly a reflection of the results for methane within the mixture. While the simulation predicts an adsorption of about 70% of the available methane , the flexible ones increases the prediction to up about 80% of the available methane. For nitrogen, the difference is smaller and anyway with less influence on the total result since the initial fraction is much lower. Keeping in mind that for the pure methane, the united-atom model predicted a lower uptake on the rigid sheet  a strong second adsorption layer is formed. As before, the assumed rigid sheet adsorbs slightly less nitrogen molecules than the flexible sheets.Looking at the simulations with the atomistic model, the behavior changes quite a bit: the first adsorption layer for nitrogen is smaller for the simulation with the most (350) molecules compared to the other simulations with 150 and 250 methane molecules. This is a consequence of the strong methane adsorption predicted by the atomistic model: the stronger competition allows less nitrogen molecules to enter the first adsorption layer. In the simulation with 150 molecules, the first adsorption layer is not yet completely saturated by methane and nitrogen is allowed space in the first adsorption layer. As the amount of molecules increases, less and less space remains available for nitrogen and the first adsorption peak of nitrogen decreases as a result. Already in the simulation with 150 molecules, there is a small second adsorption layer present, which increases strongly upon increasing the amount of molecules in the system. In the simulation with 350 nitrogen molecules, there are very little nitrogen molecules present in the first adsorption layer, but there is a relatively large second adsorption layer and subsequent gas phase.Due to the changing widths of the adsorption peaks—due to the internal graphene movement—in For nitrogen, it can be seen that the united-atom model predicts a converged adsorption within the investigated pressure range due to the strong methane competition. The rigid model predicts a slightly lower uptake and slower convergence than the flexible ones. For the atomistic model it is striking that, due to the very strong methane adsorption, the simulation with 350 molecules has very little nitrogen adsorbed on the sheet leading to a very low uptake. For this reason, we have fitted the Langmuir isotherm using only the first two simulation results as it is not intended to describe such a sudden drop in the uptake. Furthermore, the drop is caused by the methane more than by the nitrogen itself. The atomistic model predicts a very slightly lower uptake of nitrogen in the rigid simulation than when considering the flexible representations. As was expected, more nitrogen is adsorbed using the united-atom model than the atomistic model, because the methane is less strongly attracted to the graphene sheet in that case.Kumar and Rodríguez-Reinoso  have invFinally, we have calculated the selectivity as followsx is the mole fraction adsorbed of the specified molecule, while y is the mole fraction in the bulk of the specified molecule.where, Even though (Dreisbach et al., In this work, we have studied the influence of the flexibility of the graphene sheet on its separation ability of the methane/nitrogen gas mixture. The flexibility was introduced via three intramolecular force fields taken from the literature, two of which contained stretching, bonding and torsional terms, while a third one lacked the latter contribution. Two fields were specifically designed for graphene, while a third one was originally intended for use on carbon nanotubes. Furthermore we studied the different behavior of a united-atom model and an atomistic model during these simulations.In general, we have confirmed that graphene shows a strong preference for methane over nitrogen. Using both models, the methane posed a strong competition toward the nitrogen and pushed the nitrogen out of the first adsorption layer, forcing it to form a second adsorption layer or go into gas phase. This effect was found stronger for the atomistic gas models than for the united-atom models, which is expected to give a poorer representation as it does not takes into account orientation effects. The difference, which is at times large, calls for care when assuming united-atom models in these types of simulations.Concerning the treatment of flexibility, we showed that, indeed, it influences the behavior of the graphene sheet as a separator of the methane/nitrogen gas mixture. For the united-atom model, neglecting the flexibility of the graphene sheet leads to predict lower methane uptakes that if such flexibility is taken into account. The atomistic model on the other hand predicted a slightly higher methane uptake. Similar results were found for nitrogen, leading to the conclusion that for the united-atom model, a lower amount of gas molecules in general is predicted to be adsorbed if the sheet is supposed to be rigid. The atomistic model, on the other hand, predicts a larger general gas adsorption also on the considered rigid graphene sheet. Observing the combination of these results in the selectivity of the four models of the intramolecular motions of graphene, we find that the united-atom model in conjunction with the rigid representation predicts the lowest selectivity for the methane/nitrogen mixture in the largest portion of the pressure range investigated, while the atomistic model shows very similar behavior for the rigid and the flexible graphene models. Moreover, since graphene is, in fact, flexible, it is then clear that in order to safely disregard its internal movements an atomistic model is definitely required, the simplified united-atom model being absolutely inadequate. On the other hand, if the flexibility of graphene is explicitly accounted for by means of the appropriate force field -the field 2 m (Fthenakis et al., JV, IG, and AS ran the first principles calculations and developed the subsequent interaction force field. NF-L, AL, and MR developed and ran the theoretical and molecular dynamics calculations.The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.


Humidity sensors are a common, but important type of sensors in our daily life and industrial processing. Graphene and graphene-based materials have shown great potential for detecting humidity due to their ultrahigh specific surface areas, extremely high electron mobility at room temperature, and low electrical noise due to the quality of its crystal lattice and its very high electrical conductivity. However, there are still no specific reviews on the progresses of graphene-based humidity sensors. This review focuses on the recent advances in graphene-based humidity sensors, starting from an introduction on the preparation and properties of graphene materials and the sensing mechanisms of seven types of commonly studied graphene-based humidity sensors, and mainly summarizes the recent advances in the preparation and performance of humidity sensors based on pristine graphene, graphene oxide, reduced graphene oxide, graphene quantum dots, and a wide variety of graphene based composite materials, including chemical modification, polymer, metal, metal oxide, and other 2D materials. The remaining challenges along with future trends in high-performance graphene-based humidity sensors are also discussed. Th. Th86]. 2 . . 111]. FHumidity sensors based on GO materials have been widely studied due to the simple, low-cost, and large-scale preparation of GO, and high proton-conductive sensitivity to water molecules. Based on the change of the proton conductivity, GO humidity sensors generally operate by detecting the capacitance or impedance signals ,117. Bi Generally, noble metal  interdigitated patterns prepared by the lithographical technique are used as electrodes to form GO based humidity sensors. However, Au and Ag materials are relatively expensive, and the lithographical preparation needs professional instruments and a complex process. The laser direct writing technique, which is a noncontact, fast, single-step fabrication technique with no need for masks, postprocessing, or a complex clean room, and is compatible with current electronic product lines for commercial use, has been employed to fabricate energy storage devices, electrically conductive circuits, sensors, as well as self-powered integrated devices ,132,133.Recently, we reported the facile preparation of humidity sensors based on an rGO/GO/rGO structure on a flexible PET sheet through laser direct writing with a semiconductor diode laser a 131]. . 131]. A2 carbon network. A transparent humidity sensor made of rGO stripes driven by convective self-assembly exhibited a reversible response to humidity in the range of 10–70% RH . . 2 and Wy sensor b, and shy sensor c. The ely sensor d,e. Theyuctivity . BP mateuctivity . HumiditVarious effective approaches have been devoted to developing graphene-based humidity sensors with high sensing performance. First, humidity sensors based on pristine graphene show high sensitivity with a detecting limit as low as 1 ppm and fast response, however, it suffers from low selectivity and relatively slow recovery. Chemical modification on the graphene surface may be a possible way to improve the selectivity and recovery. Second, graphene oxide-based humidity sensors possess high sensitivity and fast response in the high humidity range, but it is difficult for them to detect low humidity conditions of less than 5% RH. The swelling effect and relatively low stability possibly exist at high RH ranges, and the capacitive or impedance working mode generally needs relatively complicated circuits. Third, reduced graphene oxide provides an effective strategy to construct resistive humidity sensors with simple and easy fabrication and measurements and low power consumption, but it also reduces the available functionional groups for adsorbing water molecules, thus decreasing the sensitivity and sensing humidity range. Fourth, various kinds of materials, such as chemicals, polymers, metal, metal oxide, and 2D materials, have been incorporated into graphene oxide or reduced graphene oxide film to improve the humidity sensing performance, and some of them show excellent performances, such as wide sensing RH range, high sensitivity, small hysteresis, and fast response, but they may suffer from reproducibility and long-term stability.


We focused on the use of low-concentrated ozonized water (OZW) as the biofilm formation restraint system for DUWL. Here, we examined the effects of low-concentrated OZW on the growth of bacteria and related biofilm formation and harmfulness to dental unit components (DUCs) in vitro.To evaluate the bactericidal effects of OZW on biofilms in DUWL and DUC Candia albicans solution (control microbe) were treated with OZW for 1 h with gentle agitation before static culturing for 96 h in Reasoner's 2A liquid media. The control solutions were 0.1% cetylpyridinium chloride (CPC), chlorinated tap water (TW), and phosphate-buffered saline (PBS). Adenosine triphosphate (ATP) amounts of the microbes were measured and the biofilms of these microbes were observed using scanning electron microscopy (SEM). Moreover, surfaces of DUC soaked in OZW and TW were observed by SEM.Low-concentrated OZW (0.4 mg/L) was generated using an OZS-PTDX generator. Heterotrophic bacterial biofilms in old DUWL tubes and The OZW reduced ATP levels in microbes to 50% compared to TW and PBS treatment, although CPC reduced it below detection limits. SEM observation revealed deformation of microbes cultured with OZW, whereas no changes were seen on DUC surfaces.Low-concentrated OZW is bactericidal against heterotrophic bacteria biofilms and it is not harmful to DUC, suggesting that it might be useful in preventing DUWL contamination. Bacterial contamination of dental unit water lines (DUWL) has emerged as a problem, especially in the United States rom DUWL , 8.4–107 cfu/mL at the first clinical activity of the day 2 cfu/mL by Centers for Disease Control and Prevention (CDC), the United States. Opportunistic pathogens such as Pseudomonas and Acinetobacter were also detected in DU water In addition to the oral microorganisms getting accidentally sucked and contaminating DUWLs, biofilms were reported as a new source of contamination in the 1990s, as reviewed by O'Donnell In addition to the recommendation by the CDC, the acceptable limit of heterotrophic bacteria according to the American Dental Association (ADA) is less than 500 cfu/mL in DUWL water Therefore, an easily-handled inexpensive system to suppress biofilm formation is desirable for any general dental clinic. We selected low-concentrated ozonized water to develop a system for the prevention of microbial contamination in DUWL. This system has been used for sanitary and washing purposes of companion animals, since low-concentration ozone disappears after a short time. Ozone exists at about 0.005 ppm in the atmosphere Based on the above background, the aim of this study was to investigate the bactericidal effects of low-concentrated ozonized water on microorganisms and biofilms involved in DUWL contamination and their harmfulness to the components of DUWL.22.1Ozonized water was generated from TW using a small generator, OZS-PTDX . The ozonized water generators generate ozone by surface discharge control technology 2.22.2.1Candida albicans ATCC 10231 strain  was aerobically cultured in brain heart infusion  liquid media at 37 °C for 12 h. For later experiments, C. albicans was re-suspended in test water at 108 cfu/mL.2.2.2Planktonic heterotrophic bacteria were obtained from water in a three-way syringe of DU (DUWL-W). For morphological observation, both 50 mL of the mixture (DUWL-W: OZW = 1:9) and 5 mL of DUWL-W were shaken gently and in parallel at 25 °C for 1 h, then 5 mL of each solution was filtered using a nano-percolator/SEM pore .Biofilms of heterotrophic bacteria were obtained from the inner surface of DUWL tube (DUWL-T) previously used for turbine hand-pieces of DU. DUWL-T was provided for this study in wet conditions after exchanging for a new one in the DU. After DUWL-T was cut by 1-cm length and divided vertically into 2 pieces, 6 pieces of 3-cm length were soaked in 30 mL of the test water and gently shaken at room temperature for 1 h. Bacteria present in the DUWL-T were cultured in Reasoner's 2A (R2A) liquid medium  at 32 °C for 96 h. Heterotrophic bacteria grown from biofilms were suspended in the media by vigorous mixing.2.3The components of DUWL used in this study  were pro2.4Ozone concentration was colorimetrically measured using the diethyl-p-phenylenediamine (DPD) method 2.5−16 to 3.0 × 10−11 mol of ATP). Background ATP present outside the bacterial cells was removed using adenosine phosphate deaminase.The amount of adenosine triphosphate (ATP) present in the bacteria was used to test the vital activity of bacteria. ATP amount was measured using a Lucifer HS kit and Lumi-tester C-110 , which employs luciferin-luciferase reaction 2.6C. albicans treated with each test water sample, planktonic heterotrophic bacteria in DUWL filtered by nano-percolator/SEM pore, biofilms on the inner surface of used DUWL-T, and unused DUWL components made of brass, stainless steel, ethylene-acryl-nitrile rubber, polyurethane, and polyvinylidene difluoride.The samples observed by SEM were DUWL-T was immersed in 3-methyl-butyl for 1 h as a replacement after dehydration with an ascending ethanol series , and it was sublimated with the critical point dryer . These steps were omitted for other samples.The surfaces of each sample were treated with osmium by vapor deposition method using a vacuum evaporator . The morphology of the microorganisms and surface condition of each component was observed using a field emission scanning electron microscope .2.7t-test (Microsoft Excel) after using an F-test (Microsoft Excel) to analyze the variance between groups. P-values less than 0.05 were considered statistically significant.Statistical processing was performed using Student's 33.1The ozone concentration in OZW changed over time , with a 3.2C. albicans treated with OZW reduced to half of the negative controls  for 1 min To test the bactericidal effects of OZW on planktonic microorganisms, we used albicans A and hetWL water , 8. Secowith CPC C. The la8 cfu/mL of bacteria were observed (data not shown). Heterotrophic bacteria in 5 mL of DUWL were filtered and visualized under SEM (Heterotrophic bacteria from three-way syringe were also in planktonic conditions, mainly detached from biofilm formed in DUWL. When DUWL was used for heterotrophic bacterial culture on R2A agar plate before this study, more than 10We confirmed the bactericidal effect on biofilms formed by heterotrophic bacteria inside DUWL. Because biofilms consist of layered structures These results collectively suggest that OZW may reduce the growth of heterotrophic bacteria in DUWL through continuous treatments. However, the application of OZW may be limited to unused DUWL-T, because it is unclear if OZW can destroy biofilm structure.Finally, we determined the effect of OZW oxidation on main DUWL components. The DUWL components seemed to tolerate low-concentrated OZW. Significant deterioration and abnormal findings on the surface were not observed either by naked eye inspection or SEM . HoweverThe ozonized water used in this study could reduce the growth of microorganisms without apparently harming DUWL tubes. The advantages are that there is little risk of the tube deteriorating and organic substances eluting due to the strong oxidizing power of ozone, and that microorganism adhesion is reduced due to the smoothness of the tube surface. Thus, it is reasonable to consider that the risk of unknown adverse effects for patients from using low-concentrated ozonized water is low.5Low-concentrated OZW at 0.4 mg/L is bactericidal to microorganisms involved in DUWL contamination without obvious damage to the components of DUWL. Therefore, low-concentrated OZW can be used to prevent bacterial contamination in DUWL. However, the limitations of this study include the use of static conditions for DUWL using low concentrations of OZW.Further studies using simulated DUWL with periodical clinical conditions over a long period are needed.Low-concentrated OZW demonstrated significant advantages—it effectively reduces bacterial biofilm contamination and causes less damage to DUC. Therefore, establishment of the ozonized water supply system could be a promising solution for decontamination in DUWL and may reduce the risk of hospital-acquired infections in dental practice.Keisuke Okubo: Conceived and designed the experiments; Performed the experiments; Analyzed and interpreted the data; Contributed reagents, materials, analysis tools or data; Wrote the paper.Takashi Ito: Conceived and designed the experiments; Analyzed and interpreted the data.Shogo Takashiba: Conceived and designed the experiments; Analyzed and interpreted the data; Wrote the paper.Yasuyoshi Shiota: Performed the experiments; Analyzed and interpreted the data; Contributed reagents, materials, analysis tools or data.Yusuke Kawata: Performed the experiments; Contributed reagents, materials, analysis tools or data.Tadashi Yamamoto: Contributed reagents, materials, analysis tools or data; Wrote the paper.Okayama University.This work was supported by Organization for Research Promotion & Collaboration in The authors declare no conflict of interest.No additional information is available for this paper.


In recent years, the use of the multilayer graphene sheets has been considered more than the single-layer due to the cost-effectiveness and the possibility of mass production. But this type of graphene has some kind of structural weakness due to the weak physical link between its layers. Then, in order to strengthen, many structural modifications are proposed by various techniques to manage the mechanisms at interlayer distances. In this study, the focused ion beam irradiation method has been examined to cross-link and strengthen multi-layer graphene sheets with the help of the molecular dynamics simulation technique. Then, uniaxial and transverse tensile tests were performed to check the mechanical properties of obtained cross-linked multilayer graphene sheets. The results of this research can be considered for the creation of a new class of graphene structures. Such structures could be implemented as a membrane in water desalination or as a storage foam in hydrogen or carbon dioxide storage. With their placement in the graphene sheets and the connection with a surfactant, three-dimensional structures with intermediate spaces and a high surface area could be constructed. Recently, researchers have succeeded in controlling the intermediate space of graphene layers and filling the distance with epoxy and completely fixing it without using the columnar agent and only by controlling the moisture content. This controlled intermediate space between the graphene plates can selectively pass through water molecules and prevent the transfer of ions in the water7. Such a structure is an ideal membrane in the field of desalination and water purification. In the meantime, we see theoretical structures that have not been synthesized yet, such as pillared graphene, which was introduced by Dimitrakakis et al., in 20088, a structure consisting of single layers of graphene bonded by carbon nanotubes. By the simulations, it has been proven that the doped type of this material with lithium cations can store an appropriate amount of hydrogen that has been determined for applied uses under environmental conditions. Researchers are working to create such structures to transform the energy storage process in the near future.The increasing population and global concerns for safe water supply, the global warming due to the burning of non-renewable energy sources and the scarce resources from fossil fuels, all have made water and energy two of the most challenging issues of the present century. In order to overcome the leading environmental crises, there has been an increasing effort to discovery, design and invent more advanced materials to store renewable energy and provide potable water. In the last decade, graphene and its derivative structures have attracted the attention of researchers in various fields of science, due to their exceptional properties. Graphene-related studies are not unique to the single layer of this material and include a wide range of multilayered, composite, and hybrid structures of this material. One of the many types of these structures, which has a lot of potential in many fields, is multi-layered graphene with transverse bonds. In this unique material, the existence of transverse covalent bonds, in addition to strengthening the structure of graphene, is a three-dimensional structure with high porosity and high specific area. Until now, several factors have been proposed to create graphene framework structures, such as boronic acids11. In 2014, Wu et al., In a study conducted by the molecular dynamics simulation technique, introduced a method called “Nano joining of graphene sheets by ion beam irradiation”. They showed that two overlapping graphene sheets can connect with the carbon ion beam. The two mechanisms observed for this connection, have been introduced the formation of coordination deflection and the trapping of ions between the plates. Among the most influential parameters on the binding properties, are beam energy and irradiation dose. Optimizing theses can results in better mechanical properties of the binding region than pure graphene. The optimum energy level and the dose of the beam found to be 40 eV and 1.06 × 1015 cm2 respectively12. In a similar simulation, argon beam is also used to joining graphene sheets. In this study the mechanism has been proposed as a main factor for the connection, is the creation of dangling bonds during the irradiation process. The tendency of these bonds to saturate and atomic reorganization in the area of the defect causes the formation of the covalent bonds between the sheets. It has been proven that the temperature increase is a beneficial factor that accelerates the saturation process and develops the process of reorganization. In addition, in order to emphasize on the role of irradiation, another experiment was conducted in which the possibility of graphene joining has been studied by applying a large force above the sheets instead of irradiation, and it has been shown that without the irradiation it isn’t possible to join the graphene sheets. In the end, the graphene plane’s chirality effect on the bonding process has been studied, and it has been determined that the plane with the same chirality result in a better connection13. In another study, the effect of focused ion beam irradiation on the graphene stack has also been investigated. It has been observed that a focused ion beam can form covalent bonds between graphene sheets in the stack. In the case of the mechanism of the phenomenon, it is stated that under the irradiation process, large amounts of carbon atoms are removed from the surface of the first layer. By continuing the process of irradiation, carbon atoms with dangling bonds brought around the nanopore, these carbon atoms tend to saturated with other carbon atoms. Thus, the saturation of carbon atoms with dangling bonds in different layers leads to the formation of cross-link between different layers14. In 2017, Wu et al. tried to actually test the joining of graphene sheets by ion irradiation. they used 40 eV of nitrogen ion for this purpose, and then used Raman spectroscopy characterization and atomic force microscopy to test the bonding of the layers, the results indicate the probability of connecting the pages in the overlap area15.One of the techniques that can be used to connect graphene sheets is the irradiation technique, which has already been applied in theory and in practice to connect carbon nanotubes16, storage and absorption of gas such as hydrogen and carbon dioxide20, lithium-ion batteries21, catalytic processes22 and adsorbents23.In this study, by means of molecular dynamics simulation and exploiting the idea of bonding the layers by the irradiation process and the controlling of the distance between the graphene layers by moisture content, a method for constructing three-dimensional structures of graphene using focused ion beam irradiation proposed. The irradiation of the graphene structure in addition to creating bonds between the individual planes, also can cause the connection between graphene stacks in the border regions. The strength of this three-dimensional structure perpendicular to the graphene sheets is calculated with a range from about 4 to 6 GPa, with respect to the diameter of the beam, which is significant in comparison with other substances. It seems that the use of this technique along with the method of controlling the distance between graphene layers by moisture content, pillared structures of graphene can be obtained with an appropriate strength, which can have many applications in the fields of water and energy. This structure is a suitable candidate for use in membrane filtration systems for water24. Also, all images were made by Ovito software25, and all manipulations and modifications to the structures were done with this software too. The graphene building was created with the commands related to the custom lattice. The graphene lattice was defined in such a way that the distance between the carbon atoms in the graphene honeycomb structure is equal to 1.42 Å26. The distance between the planes in the base state of the multi-layered graphene structure was 3.4 Å, and the graphene plane dimensions in all simulations were 3 × 3 nm.All simulations of this study were done with the help of large-scale atomic/molecular massively parallel simulator (LAMMPS) software package27. The energy of the collision particle was set to 40 eV12. Due to the high kinetics of the system and preventing system failure and explosion, the system time step was set to 0.1 fs. Due to the excessive temperature rise in the system during ion irradiation, graphene temperature was set up by a nose hoover thermostat algorithm in NVT ensemble at 300 K and the temperature damping parameter was 100 times of the time step. The microcanonical ensemble (NVE) was used for ion beam, otherwise, the velocity of the particles would be overshadowed and diminished. The potential used for this simulation is a hybrid potential of CH-Airebo28 and Tersoff-ZBL29 potentials. The Airebo potential models the formation and failure of covalent bonds between carbon atoms in the graphene structure, and the Tersoff-ZBL potential also models the interaction between silicon ions and graphene atoms27. The angle of the collision is 90 degrees to the graphene plane and it is perpendicular to the sheets.In the ion beam irradiation simulation, the silicon element was selected as an ion particleAtoms around the layers were fixed to prevent the movement of the layers and the collapse of the initial spacing of the planes during the irradiation process. Applied ensemble to the fixed atoms was NVE. In Fig. , a schemThe total time of the simulation process was 1.4 million time steps. Prior to the ion collision, graphene was placed at a temperature of 300 K for 200,000 time steps or 20 picoseconds in order to stabilize the original structure in the NVT ensemble. The total amount of time that graphene exposed to the beam was 100 ps. The total dose of collision ions was considered to be 1,000 to ensure the formation and completion of the pore within the graphene structure. Also, between every two collisions, 1000 time steps were taken to give the system an opportunity to be stabled. Finally, after completing the collision process, before exporting the data from the configuration, the graphene structure for the next 20 ps was placed in the NVT ensemble to stabilize, then the configuration data file was extracted for loading in tensile mechanical testing. The irradiation simulation according to the nature of the random collisions of particles was repeated for 5 times, in order to increase the accuracy of the measurements in the same conditions, in the mechanical test by averaging the data.34. Here also, this value is changed from 1.72 to 2 Å. It was observed that the selection of values other than 2 Å results in discrepancies in the results and plots of strain-stress testing in different time steps. By correction of the minimum cut-off radius, the problem of non-conformance in the data was solved. The time step for mechanical simulation was selected to 0.1 fs. The strain rate in order to accelerate the simulation process was selected to 0.01/ps. The high strain rate causes a difference of less than 5% in the final graph, but the simulation speed increases by about ten times. In all strain-stress test simulations, before starting the stretching process, the structure in order to stabilize placed in an NVT ensemble at a temperature of 300 K for 20 ps. The output data extraction rate was every 1000 time steps. Also, as mentioned earlier, each of the graphs is the result of averaging of five tests in completely identical conditions to ensure the accuracy of the results. In the stress calculations, due to the irregular geometry created after the irradiation process, the atomic volume of carbon (8.789 Å3) and its multiplication in the number of atoms are used to compute the total system volume.To investigate the mechanical properties of the structure, the strain-stress test was carried out in two directions, perpendicular to the graphene sheets in the direction of the z-vector and parallel to the graphene sheets along the x-vector. The potential used in stress and strain test is Airebo potential. In various researches, in order to correct the potential, the amount of minimum cut-off radius has been modifiedTo investigate the effect of irradiation on a multi-layered graphene sheet, 18 different states were investigated, which has been expressed by three main parameters. These three main parameters are: beam energy, beam diameter and the distance between the graphene layers. For the ion beam, two energy levels of 40 eV and 160 eV were considered to study the effect of increasing energy or velocity of particles, on the formation of covalent bonds between graphene layers. The next parameter is the diameter of the beam, for this quantity three different sizes were considered: 0.5 nm, 1.0 nm, and 1.5 nm. The last parameter is the distance between the graphene layers. This distance also has three sizes of 0.34 nm, 0.65 nm and 0.90 nm.An overview of the results shows that the increase in the beam energy from 40 to 160 electron volts has no considerable effect on the final strengths. Increasing the beam diameter from 0.5 to 1.5 nm and the gap from 0.34 to 0.65 increases the strength from 50 to 100%, respectively. Increasing the gap from 0.34 and 0.9 nm, for all three sizes of pores, reduces the strength of about 33%. Details are presented in the following sections.In Fig. et al. as a reference32  to evaluate the effect of the formation of cross-links between the graphene layers. Before the process, a strain-stress test was conducted in the “zigzag” direction of the single layer graphene (SLG) sheet in order to validate and compare the output chart in this test with a result from a similar study by Chu In all next strain-stress tests, the strain rate was set to 0.01/ps. As shown in Fig. As previously stated, by changing the irradiation parameters, states are created in which the planes have no connection with each other, or at least there is no bond between the two graphene layers. In this study, all states with at least two non-bonded sheets lacked the necessary requirements of the strain test and therefore no mechanical test was performed on them. These states are depicted in Fig. In Fig. As is observable in Fig. For comparison, two other tensile tests were performed in addition to the previous studies. In the first test, pure multilayer non-exposure graphene was subjected to mechanical testing, in order to compare the physical force between the graphene sheets and the states in which the sheets formed a covalent bond Fig. . In thisHowever, the formation of pores by irradiation process creates bonds between the graphene sheets and reinforces the strength perpendicular to the sheets, in reverse, presence of these pores can weaken the strength in a parallel direction with the sheets. In order to investigate how to change the amount of ultimate tensile strength by changing the size of the pores, three configurations with 0.5 nm, 1.0 nm, and 1.5 nm pore diameter, and a spacing of 0.34 nm, were selected from the configurations obtained during the irradiation process, and The strain-stress test was performed for them in the direction of the x vector “zigzag”. As can be seen in the diagram Fig. , the incIn addition to all the tests carried out, the irradiation process for joining and strengthening, in which two stacks of graphene were flanked side by side, were also examined. A graphene film contains large boundary regions which are formed by adjoining graphene sheets and stacks. The study of the effect of the ion beam on these areas can also be important. For this purpose, with the same conditions as previous, a similar test was conducted in the boundary areas between the two five-layer stacks of the graphene. Variable parameters in this test are the diameter of the beam, the spacing of the layers and the boundary size between two graphene stacks. The diameter of the beam for this test was 1.0 nm and 1.5 nm. Sheets in the stack were placed at a distance of 0.34 and 0.65 nm to each other, and the size of the boundary regions between the two stacks of graphene were 0.24 nm and 0.49 nm which were created by removing one row and two rows of carbon atoms from the integrated graphene structure, respectively Fig. .Figure 8According to the graphs obtained from the strain-stress test on the eight existing states Fig. , increasFor comparison, a pure five-layer graphene stack without porosity was subjected to the tensile test Fig. . The ultIon beam energy, have the least effect on the final structure, and the difference between the two structures created by different levels of energy is not tangible. An increase in the diameter from 0.5 nm to 1.5 nm has a significant effect on the increase in the volume of the intermediate links between the graphene sheets.As the distance between the graphene layers increases, the ultimate tensile strength is reduced and increased by increasing the diameter of the collision beam from 4 GPa for a beam with a diameter of 0.5 nm to about 6 GPa for a beam with a diameter of 1.5 nm.Performing a transverse tensile test after ion beam irradiation on a ten-layer graphene stack, for all simulations, the lower layers contain more density of links, because of transferring of carbon atoms from the upper layers to the lower layers. Indeed, the lower layers receive more carbon atoms from the higher layers, which increases the probability of bonding between the underlying layers. The removal of the first layer with low bonds density improved the graphene strength by about 40%. Also, removal of the five upper layers has increased the ultimate tensile strength from about 4 GPa to about 8 GPa, which indicates the more density and strength of bonds in deeper layers of graphene than the surface layers.Increasing the size of the pores significantly has reduced the graphene strength so that by increasing the pore size from 0.5 nm to 1.0 nm, about 20% of the graphene strength is reduced, and increasing it from 1.0 nm to 1.5 nm resulted in a strong reduction (about 60%) in the strength. However, in the worst case, i.e. porous graphene with a 1.5-nanometer pore, the ultimate tensile strength is about 30 GPa, which is far superior to many other materials.Finally, nano welding of graphene sheets was proposed via ion beam irradiation. By increasing the boundary size from 0.24 nm to 0.49 nm, the ultimate tensile strength has been halved. Increasing the diameter of the beam from 1.0 nm to 1.5 nm, with a boundary size of 0.24 nm, has not had a significant effect on the graph, but in the case of a 0.49 nm of boundary size, the increase in beam diameter contributed to the increase in strength and left a notable increase in the strain-stress curve. Meanwhile, the change in the distance between the graphene sheets has had the least effect on the ultimate tensile strength. As a fantastic result, in the best case, the irradiation process improves the ultimate tensile strength by 1000% in the boundary area between two graphene stacks, which is about one-tenth of the integrated graphene structure.Bonding graphene layers by irradiation process and controlling of the distance between the graphene layers by moisture content, a method for constructing three-dimensional graphene structures was proposed. By presenting this new idea, the following results are achieved:Graphene Nanopore Creation