An equilibrium between the growth rate and the etching rate is obtained, and beyond a critical time, the total number of the carbon atoms and the edge carbon atoms do not change. Using an optimal first-order etching chemistry, a log-mean pore-size of 5.0 ± 1.7 (number of missing carbon atoms), and a pore-density of 3 × 1013 cm−2 was achieved. A high throughput calculation route for estimating gas selectivity from ensembles of thousands of nanopores was developed. The optimized result yielded H2/CO2, H2/N2 and H2/CH4 selectivities larger than 200, attributing to elongated pores generated by the competitive etching and growth. The approach of competitive etching during the crystal growth is quite generic and can be applied to a number of two-dimensional materials.A robust synthesis methodology for crystallizing nanoporous single-layer graphene hosting a high density of size-selective nanopores is urgently needed to realize the true potential of two-dimensional membranes for gas separation. Currently, there are no controllable etching techniques for single-layer graphene that are self-limiting, and that can generate size-selective nanopores at a high pore-density. In this work, we simulate a unique chemical vapor deposition based crystallization of graphene on Cu(111), in the presence of an etchant, to generate a high density (>10 Especially, the development of nanoporous single-layer graphene films has been intensively pursued for energy-efficient separations11. Single-layer graphene is atom-thick, is chemically and thermally stable in a non-oxidative environment, is stable in air at least up to 200 °C12, and is one of the strongest materials known (elastic modulus of 1 TPa)13. However, the synthesis of nanoporous graphene lattice with a narrow pore-size-distribution (PSD) and a high pore-density is not trivial14. Since the pristine graphene lattice is impermeable, one needs to controllably etch the lattice to generate gas selective nanopores. The lattice etching needs to be carried out in a controlled manner to open only those pores that are selective to the gas molecule of interest, and to avoid large non-selective pores. Here, selectivity could be size-selectivity (molecular sieving)15, shape-selectivity (entropic selectivity)16, and adsorption-selectivity17. Among these, size-selectivity has the potential to yield the highest separation selectivity, where a complete blockage of larger gas molecules can be achieved.Nanometer-thick nanoporous membranes have received much attention owing to their potential to yield large molecular flux18. H2 permeance greater than 10−7 mol m−2 s−1 Pa−1 was achieved from a pore-density of 1010cm−2, while the membrane yielded H2/CH4 selectivities in the range of 6–25. More recently, hydrogen sieving from hydrocarbons (CH4 and C3H8) using plasma- and ozone-based etching of single-layer graphene was achieved with H2 permeance reaching 2 × 10−6 mol m−2 s−1 Pa−1 and with H2/CH4 and H2/C3H8 selectivities exceeding 30 and 200, respectively19. Competitive diffusion rather than competitive adsorption determined gas selectivity with H2/C3H8 mixture separation factor exceeding the corresponding ideal selectivity. Interestingly, molecular simulations and transport calculations have shown that nanoporous graphene lattice, hosting size-selective nanopores at a high enough pore-density, exhibits ultrahigh gas permeance24. For instance, with a pore-density of 1012cm−2, a H2 permeance larger than 10−5 mol m−2 s−1 Pa−1  has been predicted by several molecular dynamics (MD) simulations22. Jiang et al. calculated energy barriers of 0.22 and 1.60 eV for H2 and CH4, respectively, from a hydrogen-functionalized pore-10 (pore made by etching 10 atoms) using the density functional theory (DFT)15. Their study employed a high pore-density of 5.3 × 1013 cm−2. The low energy barrier for H2 (0.22 eV) combined with the high pore-density corresponds to a H2 permeance of 10−5 mol m−2 s−1 Pa−1 . Using classical molecular dynamics (MD) simulations for the translocation of H2 across an unfunctionalized pore-10, Sun et al. estimated a H2 permeance of 10−3 mol m−2 s−1 Pa−1, and a H2/N2 and H2/CH4 selectivities greater than 10 and 100, respectively22. Therefore, the presence of small nanopores (close to pore-10) in graphene at a high pore-density can enable both high selectivity and permeance.Recently, we demonstrated that low-density of intrinsic defects in graphene can yield an attractive separation performance26. Therefore, one of the biggest bottlenecks in the realization of graphene membranes is the synthesis of nanoporous graphene possessing size-selective nanopores at a high pore-density. Herein, using a kinetic Monte Carlo (kMC) algorithm, we report a promising crystallization methodology capable of incorporating a high-density of size-selective, elongated nanopores in graphene in a scalable manner. This is achieved by subjecting a nanoporous graphene lattice  to the chemical vapor deposition (CVD) conditions in the presence of an etchant  surface. For etching, a first-order etching chemistry with etching rate proportional to the number of nanopore edge atoms was applied. A high-throughput transport calculation from the ensemble of nanopores indicates that attractive separation selectivities  at a high pore-density (>1013 cm−2) is indeed possible by such approach.The conventional pore etching techniques  suffer from a tradeoff between the PSD and the pore-density, especially in the sub-nanometer pore regime, attributing to a much higher reactivity of the pore edges compared to that of the basal planeant Fig. . The intet al., the sites on the Cu surface corresponding to the underlying hcp and the fcc cites are the most probable sites for the adsorption of carbon atoms29. Accordingly, all deposition and surface diffusion events were restricted to the hcp and fcc sites. This routes makes sure that pristine graphene can be formed on the Cu surface , a Cu surface of the dimension 26.5 nm × 26.5 nm was created which contained 12480 Cu atoms. In this fashion, three atomic layers of Cu were created Fig. . As per ace Fig. . As carb4 as a growth precursor  at 1000 °C on the bare Cu(111) surface  than the activation energy for the free diffusion of the precursors (0.5 eV)28. Therefore, once the bond is formed, there is a low probability of a bond cleavage.The kMC algorithm used CH2 . The high nucleation density in this study is attributed to a high CH4 deposition rate implemented in this study primarily to reduce the computational time. The grains grew progressively as a function of CH4 deposition steps  when the lattice was subjected to the competitive growth and etching. Keeping the growth rate was fixed, the etching rate was varied. Etching was ensued using a hypothetical etchant following a first-order etching chemistry with respect to the number of carbon edges. A limitation on etching chemistry was imposed such that the etchant could only etch graphene lattice from the pore-edge terminations, while the basal plane etching was not allowed. This is true for mild etchants such as carbon dioxide in the CVD conditions  were chosen such that the etching rates were close to that of the growth rate. In experiments, the etching rates can be varied by varying the partial pressure of the etchant gas. Applying the kMC algorithm led to nanoporous structures with the PSD and the mean pore-size dependent upon the etching rate constant . Then, the lattice was decorated with defects Fig.  with a mm−2 Fig. . Experimons Fig. 31. The ent Figs. .Figure 3−1 led to a lattice resembling the pristine graphene lattice, indicating that the nanopores shrank completely , 5.8 ± 2.7 (20000 s−1), 9.7 ± 7.5 (25000 s−1), and 16.4 ± 24.8 (30000 s−1). Extremely high pore-densities were obtained, reaching up to 1014 pores cm−2 . The number of carbon atoms reduced by 14.7, 28.6 and 36.7% at the etching rate constants of 20000, 25000 and 30000 s−1, respectively. Accordingly, the pore-density increased to 4.2 × 1013, 9.9 × 1013, 10.9 × 1013, and 8.4 × 1013 cm−2 at the etching rate constants of 18000, 20000, 25000 and 30000 s−1, respectively . The ratio became close to half (0.45) at the etching rate constant of 30000 s−1. It is noteworthy that the ratio became constant after a period of time, indicating an equilibrium between the etching and the growth of graphene lattice at the pore edge. This is quite attractive for regulating the structure of the nanoporous graphene.The total number of carbon atoms as a function of time highlights the overall porosity of graphene. To understand the size and shape evolution of pores, the ratio of the number of edge atoms with that of the total atoms was probed Fig. . At the −1 and 20000 s−1), the lattice may still be attractive for the gas separation.The generated nanopores can be characterized by the number of missing carbon atoms as well as their shape. While, the number of missing carbon atoms may indicate nanopore area, for membrane application, it is not necessarily a true indicator of the edge-to-edge gap especially if the pores are elongated. This is important because even the pores with a large number of missing carbon atoms can be attractive for molecular sieving if they are elongated or slit-shaped. Therefore, to characterize the pore-shape, the ratio between the number of missing carbon atoms and the number of edge atoms was probed. Here, the edge atoms are defined by those carbon atoms that are connected by less than 3 carbon atoms. Generally, when the number of missing carbon atoms increases, the ratio tends to increase Fig. . However21. Briefly, the gas transport takes place in five key steps, (i) adsorption of the gas molecule on the surface of graphene on the feed side, (ii) diffusion of the molecule towards the pore and association with the pore on the feed side, (iii) translocation of the molecule through the pore from the feed side to the permeate side, (iv) dissociation of the molecule from the pore and subsequent diffusion away from the pore on the permeate side, and finally (v) desorption of the molecule from the graphene lattice on the permeate side. Typically, the rate-limiting step is the pore translocation step where the molecule arriving at a transition state in the center of the pore experiences an energy barrier from the electron cloud at the pore-edges. Therefore, the overall transport is activated with activation energy corresponding to the potential energy difference between the adsorbed state and the transition state. Our experimental measurements of gas transport through the intrinsic defects of graphene confirmed this, and a temperature activated gas transport was observed19. Accordingly, assuming the translocation step as the activated step and a rate controlling step, a mathematical model was derived to estimate the gas pair selectivity (section S1). Overall, the gas pair selectivity, αAB, from an ensemble of N pores, each with an area of Oi can be written as Eq. Ht.s.,A,i−gas,A refers to the potential energy difference between two states of gas A when crossing pore i. These two states are transition-state of the gas during the pore-translocation when the molecule is in the cross-section of graphene pore, and the state in the gas phase. Similarly, ΔHt.s.,B,i−gas,B refers to the potential energy difference between the corresponding two states of gas B. T is the temperature. R is the universal gas constant.The gas transport across the nanoporous graphene was modeled adopting the adsorbed-phase transport model proposed by Strano and co-workersHt.s.,A,i−gas,A for more than a thousand pores generated in this study by traditional methods  is not feasible. Therefore, we developed a MATLAB based high-throughput algorithm probing the interaction of gases  with the nanoporous graphene lattice and extracting the energy of the molecules at the transition state. Energy calculations were carried out considering the van der Waals interactions  with the step size of 0.02 nm. At each z-step, the position and the rotation of the gas molecule were optimized by calculating the minimum potential energy. The adsorption energy was assigned to the minimum potential energy along the z-steps. The transition-state (inside the nanopore) was assigned to the minimum potential energy at z equal to zero.The LJ interactions between different types of atoms were evaluated using the Lorentz-Berthelot mixing rule. A cut-off distance of 1.2 nm was applied to the LJ pairwise potentials. To be close to the realistic conditions of the pore-structure generated in the CVD process, the pore edges were hydrogen terminated. The parameters used in the simulation are listed in Table 2, N2, CH4 and CO2 respectively, on pristine graphene, which is in good agreement with the literature33. The corresponding adsorption heights of He, H2, N2, CH4 and CO2 on pristine graphene were 0.29, 0.31, 0.33, 0.35 and 0.43 nm, respectively. The ensemble of pores was simulated in the presence of these gas molecules to generate a dataset of ΔHt.s.,A,i−gas,A, from which the gas pair selectivity was calculated  estimated here is highly attractive for the gas separation.The model yielded adsorption energies of 1.9, 4.0, 8.5, 11.2 and 14.5 kJ/mole for He, Hted Fig. . Attract13 cm−2 and sub-nanometer sized pores is extremely attractive for molecular separation. Our unique approach, involving the use of an etchant during the CVD of nanoporous graphene with CH4, leads to an equilibrium between the growth and etching of the edge atoms. As a result, beyond a critical CVD time, the total number of carbon atoms and edge atoms in the nanoporous graphene do not change as a function of time. Resulting nanoporous graphene structures show a high promise for gas separation with H2/CO2, H2/N2 and H2/CH4 separation selectivities greater than 200 and the pore-density greater than 1013 cm−2 at the etching rate constant of 18000 s−1. The realization of such membranes could be highly interesting for application in pre-combustion and post-combustion carbon capture, hydrogen purifications, and helium recovery from natural gas.The development of nanoporous graphene hosting pore-density larger than 104 deposition rate, growth temperature, etc. The highly porous lattice of graphene is also expected to find applications in band-gap engineered optoelectronic devices.Overall, the competitive etching method introduced here overcomes the tradeoff between the mean pore-size and the pore-density typically observed with the conventional graphene etching methods . The overall approach has a significant potential to be optimized with respect to etching rate constant, CH4 on the surface of Cu, (b) dehydrogenation of CH4 forming radicals, (c) surface diffusion of the radicals, and (d) the subsequent nucleation and growth of the graphene domains. When graphene lattice is already present on the Cu surface, the nucleation event is not required, and the growth precursors can directly attach to the existing lattice edges. Herein, we introduce an etchant in the CVD reactor along with the growth precursor (CH4) to tune the structure of nanoporous graphene  adsorption of CHene Fig. . A mild cts Fig. . The pur4 on Cu(111) surface was calculated according to the adsorption model where the deposition rate is dependent on the surface coverage  is the fractional surface coverage. In this simulation, CH4 can only be deposited on the fcc and hcp layer as mentioned in above section. Accordingly, there are 24960 sites for the deposition of CH4. It has been assumed that the gas molecules can be deposited into every site has equal probability and every deposition is successful (noted as CH4(s)). ra can be modeled by the Knudsen’s formalism34, where the rate is dependent on the partial pressure of CH4 in the gas phase (noted as CH4(g)). As per Eq. P is the CH4(g) partial pressure, and m is the molecular weight of CH4. T is the deposition temperature, and kB is the Boltzmann constant, and A is the area of deposition.The physisorption of CH4(g) is adsorbed on the Cu(111) surface, Cu surface atoms catalyze the dehydrogenation reaction leading to the formation of growth precursors. The activation energies for the successive dehydrogenation steps (CH4(s) → CH3(s) → CH2(s) → CH(s) → C(s)) on Cu(111) have been calculated by Zhang et al. using DFT including the Arrhenius dependence, and are listed in Table 27. The rate of the dehydrogenation step, i(s) species  can be written as a first-order kinetics (Eq. i(s) radicals on the surface, w0 is the frequency factor calculated from the Eyring equation and is equal to 2kBT/h, where h is the Plank’s constant. i(s). The overall rate of the process, rd, can be expressed as a summation of the dehydrogenation steps for all the species (Eq. After CHon the Cu11 surfac2 is immediately removed from the surface (adsorbed inside the Cu lattice or released in the gas phase), and that it does not play a role in CVD. Consequently, the dehydrogenation step is irreversible and the concentrations of the intermediate radicals (CH3(s), CH2(s), CH(s)) are reduced. In this scenario, the surface diffusion of intermediate radicals can be neglected. Therefore, the crystallization of graphene proceeds by the surface diffusion of C(s), and subsequent bond formation.For simplification, we have assumed that Het al. calculated the surface diffusion energy considering the interaction with the neighbor  and the next to next neighbor  using the ab-initio calculations and expressed the activation energy for a particular type of diffusion as a function of nne,i and nnne,i28. In this case, for a carbon which is present in hcp site, it can have maximum three fcc neighbors and six next to next hcp neighbors and vice-varsa ,i is the number of carbon radicals associated with the activation energy of Eact,sd,i. vi is the number of directions the molecule can diffuse while taking into account the degree of freedom for the diffusion which is equal to the number of empty neighbor atom present.Hopping of carbon radicals from one adsorption site to another is marked as a single diffusion step. The single diffusion step is restricted to the nearest empty sites . So, carbon on an hcp Cu site can jump to the nearest fcc sites and vice-varsa. Based on the number of carbon radicals close to the particular diffusing radical, the activation energy for diffusion will be different due to the interaction with the other carbon radicals. A large number of neighboring radicals would imply that the diffusing radical will need higher energy to arrive at the transient state. For example, Gaillard rsa Fig. . The gen2 can be used as an etchant which neither nucleates defects on graphene basal plane nor etches the grain boundaries ,edge is the number of edge carbon atoms on the surface.The etching is modeled with a first-order etching rate assuming a single etching rate constant. For example, COies Fig. . The firet. al.35. Once a CH4(g) molecule is adsorbed on the Cu(111) surface, it gets dehydrogenated sequentially to form a carbon (C) radical. Next, the C radicals diffuse to the graphene nanopore edge and contribute to the edge growth. In parallel, nanopore edge atoms are etched by the etchant. According to this model, the deposition of CH4 on Cu occurs with a constant rate , the next deposition of CH4 takes place, and again the Δts count starts from zero.The evolution of nanoporous graphene was simulated with the kMC algorithm to study system evolution at a large time-scale. The kMC algorithm was adopted from the study of Lucas on the Cu1 surface Δta Eq. . In this Δta Eq. . If the . ts Eq. . After a4, validating the kMC model for the graphene growth  was demonstrated by introducing only CH11 was deet al., where it was found that the sites on the Cu surface corresponding to the underlying hcp and the fcc cites are the most probable sites for the adsorption of carbon atoms29.In our model, carbon atoms are only bound to the hcp and fcc layer of the copper lattice. This assumption follows from the calculation of Riikonen 2 show that the grain-boundaries are resistant to etching . The zigzag edge has been shown to be more energetic than the armchair edge due to the confinement of POverall, the assumption made in the kMC algorithm are:Supplementary Information


Graphene is mainly implemented by these methods: exfoliating, unzipping of carbon nanotubes, chemical vapour deposition, epitaxial growth and the reduction of graphene oxide. The latter option has the advantage of low cost and precision. However, reduced graphene oxide (rGO) contains hydrogen and/or oxygen atoms hence the structure and properties of the rGO and intrinsic graphene are different. Considering the advantages of the implementation and utilization of rGO, voltage-dependent electronic transport properties of several rGO samples with various coverage ratios are investigated in this work. Ab initio simulations based on density functional theory combined with non-equilibrium Green’s function formalism are used to obtain the current–voltage characteristics and the voltage-dependent transmission spectra of rGO samples. It is shown that the transport properties of rGO are strongly dependent on the coverage ratio. Obtained results indicate that some of the rGO samples have negative differential resistance characteristics while normally insulating rGO can behave as conducting beyond a certain threshold voltage. The reasons of the peculiar electronic transport behaviour of rGO samples are further investigated, taking the transmission eigenstates and their localization degree into consideration. The findings of this study are expected to be helpful for engineering the characteristics of rGO structures. Graphene is an atomic-thick layer made of carbon atoms in honeycomb lattice structure . Graphen−1 [13C solid-state nuclear magnetic resonance [The reduction of GO to rGO is a low cost method for obtaining graphene; however, the main drawback is that the rGO is not an exact graphene sheet but a relatively imperfect honeycomb crystal with controllable oxygen and hydrogen content. Moreover, the atomic structure and the chemical formula of GO and rGO is still under debate , 21, 22.−1 . Similar−1 . Zhang e−1 . Similar−1 . A simil−1 . On the −1 . Similaresonance . Xu and esonance . Among tAlthough the conductance properties of various rGO configurations are investigated in the literature as summarized above, the conductance characteristics are estimated using the obtained band gaps and DOS characteristics which show zero-bias behaviour. However, it is recently shown that the conductance characteristics of functionalized graphene structures have to be evaluated by taking the voltage-dependent transmission spectra and voltage-dependent transmission eigenstates into consideration  which is®) program based on the transiesta method is utilized for obtaining the transport properties of the rGO structures considered in this paper.There are various methods used for analysing the characteristics and behaviours of nano-scaled structures in the literature such as tight-binding calculations , semi-emAs it is stated before, rGO structures with more than 25 % coverage are expected to have both hydroxyl and oxygen contamination while the conducting rGO sheets with 25 % or less coverage contain hydroxyl groups only. Self-consistent energy calculations show that for rGO structures with more than 25 % coverage, chemisorption energy is 60 and 30 meV lower per hydroxyl and oxygen, respectively, compared to configurations having same type of contaminations of the same coverage ratios . Moreove® with DFT method using the following parameters: mesh cut-off is 150 Ry, electron temperature is 300 K, k-point sampling is  and the exchange-correlation functional is local density approximation (LDA) which gives accurate results for carbon-based structures [−1. After the optimization step, six periods of scattering region for each rGO sample are built with the optimized structures. The reason for selecting six periods of scattering region is that it is enough to avoid electrode interactions while providing a balanced simulation cost [ The unit cells of the considered rGO structures are built in ATK with the approximate oxygen and hydroxyl group positions . These gion cost . As a reion cost .Fig. 1Co® for 21 applied voltage points in the 0–1 V range as shown in Fig. I–V characteristics is that at zero-bias; rGO with 50 % coverage shows non-conducting characteristics, while rGO samples with coverage ratios of 25 % and lower clearly show conducting characteristics as shown inside the circle. These results are in accordance with [I–V behaviour in this range shows peculiar properties as explained below.Firstly, the current–voltage variations of the intrinsic graphene sheet and the rGO samples with different coverage ratios are obtained in ATKI–V behaviour of the rGO sheet with 25 % coverage shows a linear trend implying that the transmission eigenchannels of 25 % coverage rGO is maintained in the 0–1 V range. Another important result is that rGO samples with 6.25 and 12.5 % coverage ratios show a negative differential resistance (NDR) characteristic between 0.2 and 0.4 V. Beyond 0.4 V of applied voltage, rGO with 6.25 % coverage ratio shows non-conducting characteristics, while the rGO with 12.5 % coverage has a constant current behaviour. On the other hand, it was shown in previous studies, that rGO with 50 % coverage would show non-conducting characteristics at zero-bias [I–V behaviour obtained in this study also supports this prediction for applied voltages around zero-bias. However, as the applied voltage increases, rGO with 50 % coverage also shows a conducting behaviour as it can be seen from Fig. The ero-bias . The I–VI–V characteristics of ballistic nano devices stem from the voltage-dependent variation of their transmission spectra [I, of a nanodevice is given as in Eq. (T(E) is the transmission spectrum, µR and µL are the electrochemical potentials of the right lead and the left lead, respectively, q is the elementary charge and h is the Planck’s constant. fL(E) and fR(E) are the Fermi–Dirac functions of the left lead and the right lead, respectively [It is well-known that the  spectra , 53. Thes in Eq.  where T characteristics of the intrinsic graphene sheet and rGO samples are also calculated as shown in Fig. The conductance–voltage  while rGO with 50 % coverage shows non-conducting behaviour near zero applied voltage as estimated by previous studies [I–V characteristics can be engineered to achieve specific functions in nanoscale circuits easily since the coverage ratios of rGO structures can be precisely controlled using reduction by laser light.Considering the reduction of GO is an effective method to implement conducting rGO structures, voltage-dependent transport properties of rGO samples with different coverage ratios are investigated in this study. Electronic transport behaviours of rGO samples with coverages of 6.25, 12.5, 25 and 50 % are obtained together with the properties of intrinsic graphene sheet for comparison. It is exposed that depending on the coverage ratio, rGO structures show different and peculiar electronic transport characteristics such as (i) rGO with 6.25 and 12.5 % coverages show NDR behaviour, (ii) the transmission of the GO sample with 6.25 % coverage ratio diminishes after the NDR voltage range, (iii) the  studies , it star


One of the main advantages of 2D materials for various applications is that they can be prepared in form of water‐based solutions. The high yield and cost‐effectiveness of this method make them of great interest for printed electronics, composites, and bio‐ and healthcare technologies. However, once deposited on a substrate, etching away these solution‐processed materials is a difficult task, yet crucial for pattern definition and thus device fabrication. In particular, the realization of micrometer‐sized patterns requires mesh and paste optimization when screen‐printed or solvent‐engineered and surface functionalization when inkjet‐printed, both usually involving additional postdeposition steps. These constraints are holding back the integration of these 2D materials in devices and applications. In this work, a method for the fabrication of micrometer‐sized well‐defined patterns in water‐based 2D materials is presented, with an extensive characterization of the films and patterns obtained. The method is ultimately used to create humidity sensors with performance comparable to that of commercial ones. These sensor devices are fabricated onto a 4′ silicon and polyethylene terephthalate (PET) wafers to create all‐graphene humidity sensors that are flexible, transparent, and compatible with current complementary metal–oxide–semiconductor (CMOS) and roll‐to‐roll workflows. Monitoring humidity has become significant in many fields from electronic‐skin‐based diabetes therapy to industrial manufacturing. For example, sensing humidity is required for the proper functioning of diabetes patches,The scalability and cost‐effectiveness to produce solutions of 2D materials make them of big interest for the industry. Additionally, their production and deposition technologies, generally carried out at room or moderate temperatures, dramatically ease the integration of 2D materials with flexible polymeric substrates such as polyethylene terephthalate (PET) and polyethylene naphthalate (PEN) that are used massively as flexible electronic substrates. Indeed, recent developments of the deposition technologies of 2D solutions for flexible electronics have demonstrated vacuum filtration,Moreover, manufacturing devices integrated with liquid‐exfoliated 2D materials should be compatible with the existing approaches, especially with complementary metal–oxide–semiconductor (CMOS) technology, which is used ubiquitously for device processing, with no prospects to change in a short span of time. For a fast and successful integration of liquid‐exfoliated 2D materials into the semiconductor foundries, it is necessary to make use of existing CMOS techniques adapted to 2D materials, before specific processes may be gradually introduced into the production lines. Similarly, the emerging R2R processing is set to disrupt the way electronic devices are manufactured, especially fabrication of transparent and flexible electronics. This technique is already employed intensively to produce conductive patterns using Ag‐ and Cu‐based inks. The new properties achievable with the range of 2D inks available, suitable for cost‐effective production, will produce a much wider variety of devices fabricated by R2R, driving down material and manufacturing costs. Here, we demonstrate a workflow compatible with CMOS and R2R flexible technologies that allows 4′ wafer‐scale integration of both graphene and graphene oxide inks, specifically all‐graphene humidity sensor array. We employed water‐exfoliated graphene and GO solutions to create all‐graphene humidity sensors integrated with interdigitated graphene electrodes and thin GO films as sensing layer over wafer scale, with CMOS and R2R compatibility and room‐temperature processing techniques. This new technique allows the accurate placing and deposition of solution‐processed graphene and graphene oxide in a large scale without the etching process, overcoming one of the main issues of these solutions. These sensors have uniform response, good lifetime, transparency, and flexibility. In principle, this method can be conveniently adapted to any kind of 2D material dispersion to create patterns and devices with different functionalities with high yield over large areas.Figure2 substrate . Then, PR is exposed with an UV source along a digital input pattern mask to create a pattern, and is then developed. As a pattern, we demonstrate here an interdigitated electrode geometry, which is the most widely accepted geometry for the electrodes of a sensor as it enables a wide contact area between the electrodes within a limited sensor area. The interdigitated PR patterns were obtained for the electrodes of the humidity sensor and identification of conductivity and fine deposition of graphene, respectively, as shown in Figure N‐methyl‐2‐pyrrolidone (NMP) under mild ultrasonication (lift‐off process), resulting in the fine patterned graphene film  were patterned, as shown in 3O+) that behave as charge carriers.FigureBenefiting from the deposition and pattering methods demonstrated above, graphene humidity sensors were fabricated, and their performance was analyzed. The physical process giving rise to this response is summarized as follows: in GO, planes and edges contain oxygen functional groups, which gives GO its hydrophilicity and also their naturally insulating properties.2 wafer where we created devices of 1 cm2 as shown in Figure2, contacted with a thin wetting layer of Cr (5 nm) and thick layer (100 nm) of Au; in Figure 2 devices were created and contacted with commercial carbon paste, creating all‐carbon‐based humidity sensors. The performance of these devices before and after 2000 bending cycles (5 mm radius) is shown in Figure We mentioned before that for the successful incorporation of the 2D materials into the industries, the workflow needs to be adapted to their requirements such as the implementation of mass production and reproducibility (high yield). We note that all the processes carried out so far are compatible or have potential to be introduced into CMOS back‐end production line (BEPL). To assess this, we processed a 4′ Si/SiOS = (G90 − G40)/G40 = 0.2%,xG is the conductance at an x value of %RH, in our case 90% (top) and 40% (bottom) humidity. As opposed to other sensors, the range of sensitivity of our sensors is of interest for low power, Internet of Things (IoT) applications, where simple electronics are needed, which may not cover a range of signals in the range of thousands of microvolts. Moreover, the voltage supply employed in our device was 1 V, with an estimated current of 1 µA, which gives a power of 1 µW for normal performance. For comparison, the CMOS sensor needs 200 µA at 5 V supply, which translates into a powering need of 1 mW, i.e., three orders of magnitude more power required than in the case of our sensors.Finally, we assessed the sensitivity of our devices as 0.2%, calculated as 2 wafer with exceptional CMOS compatibility; we also employed this process to create all‐graphene flexible, transparent devices on PET substrate for the first time, which are also resilient to intense mechanical stresses. We strongly believe that this technique may become a cornerstone when integrating solution‐processed 2D materials into real‐world, CMOS‐based applications in the future, and we forecast massive potential for applications where cost‐effective and disposable electronics are needed.In conclusion, this work reports a combination of techniques that allowed to accurately create patterns and deposit solution‐processed 2D materials such as graphene and GO. We would like to emphasize that, after parameter optimization, this process should be compatible with any type of 2D materials, opening up ways to construct functional yet cost‐effective, flexible, fully functional electronics with these materials. Several transfers can be performed to increase conductivity and film uniformity. Stacks of different materials and films could potentially be made using this technique, and at least four depositions are necessary to create graphene patterns that are conductive in our work. In the polymer‐free process, the final 2D materials are never in touch with a polymer, reducing the well‐known polymeric contamination. The potential of the process is further demonstrated by comparing the performance of four different devices and showing that their performances are very similar and correlate very well with that of a commercially available sensor. Finally, the scalability of this process was shown creating a whole 10′ Si/SiOLithography: PMGI SF6 (from MicroChem Corp.) was spun at 4000 rpm, baked at 180 °C for 360 s. After then, photoresist S1813 (MicroChem Corp.) was spun at 4000 rpm and baked at 120 °C. Resists (S1813 (top)/PMGI (bottom)) deposited on desired substrates were exposed in a laser‐writer (Durham Magnetoptics Ltd) by UV light along the digital input pattern mask to create a pattern, and exposed S1813 resist was developed in an alkaline‐based solvent (MF‐319 from MicroChem Corp) as well as at the same time, the PMGI layer underneath was etched by the same solvent along the S1813 resist pattern. Then, the S1813 resist was stripped off with acetone and IPA.Graphene Solutions and Films: The solution was created with a concentration of 15 mg mL−1 for graphite powder from Sigma‐Aldrich and 5 mg mL−1 of sodium cholate from Sigma‐Aldrich mixed in 800 mL of de‐ionized water (DIW). This solution was mixed with a high shear mixer from Silverson for 60 min at 4500 rpm. They were then centrifuged for 60 min at 1500 rpm and decanted. This solution was additionally dipped in an ultrasonic bath for 90 min, and further centrifuged for 90 min and then decanted to remove unwanted thick graphene flakes. About 10 mL of this solution was dispersed in 60 mL of DIW and vacuum‐filtered on a PTFE membrane from Millipore, which was finally used for the IDT.IPA‐Assisted Direct Transfer: The sample with the PMGI pattern was placed on a hotplate at 90 °C, covered with IPA, and immediately after the membrane was placed on top of the sample, with the graphene facing down. After ≈300 s, all the IPA was evaporated and the membrane was removed. This step was repeated after desired films were deposited.Characterization of Patterned Graphene and GO: Raman mapping was performed using a RM1000 from Renishaw with a 532 nm laser source and air‐cooled charge‐coupled detectors. The spot size was ≈1 µm and the pitch was 2 µm to ensure no overlap between scans. The number of graphene layers was extracted by the method that was developed by Paton et al.Humidity, Temperature, and Conductance Measurements: RH was measured using a HIH‐4000‐003 humidity sensor from Honeywell; temperature was measured using a TMP36GT9Z‐ND temperature sensor from Digi‐Key Electronics. They were both powered with 5 V using a 2400 power source from Keithley, and voltage drops were recorded with a 34 401 multimeter from Agilent. Graphene/GO sensors conductance were recorded with an HM8118 LCR bridge from Rohde and Schwarz at 1 kHz and 0.5 V. The data were collected using LabView software from National Instruments.The authors declare no conflict of interest.SupplementaryClick here for additional data file.


Hydroxyapatite (HA) has been widely used in fields of materials science, tissue engineering, biomedicine, energy and environmental science, and analytical science due to its simple preparation, low-cost, and high biocompatibility. To overcome the weak mechanical properties of pure HA, various reinforcing materials were incorporated with HA to form high-performance composite materials. Due to the unique structural, biological, electrical, mechanical, thermal, and optical properties, graphene has exhibited great potentials for supporting the biomimetic synthesis of HA. In this review, we present recent advance in the biomimetic synthesis of HA on graphene supports for biomedical applications. More focuses on the biomimetic synthesis methods of HA and HA on graphene supports, as well as the biomedical applications of biomimetic graphene-HA nanohybrids in drug delivery, cell growth, bone regeneration, biosensors, and antibacterial test are performed. We believe that this review is state-of-the-art, and it will be valuable for readers to understand the biomimetic synthesis mechanisms of HA and other bioactive minerals, at the same time it can inspire the design and synthesis of graphene-based novel nanomaterials for advanced applications. Hydroxyapatite (HA) is one of the important inorganic components of the human teeth and bone, and it has shown wide applications in the fields of biomaterials science , tissue However, it is difficult to apply the pure HA for bone tissue engineering and biomedical implant due to its relative weak toughness and tensile strength. To solve this problem, the combination of HA with other kinds of bioactive tough materials, such as Ti alloys , polymerGraphene has attracted great attention in the fabrication of various novel nanomaterials due to its unique two-dimensional (2D) structure and innovative biological, electrical, mechanical, thermal, and optical properties ,24,25,26Previously, a few review papers on the biomimetic graphene-HA hybrid composites for biomedical applications have been reported ,31,32. FInspired by nature, biomimetic strategies have been widely utilized for the facial synthesis of various bioactive nanomaterials without using high temperature, high pressure, and hazard chemicals ,34. Typi2+ and HPO42− mediated the nucleation and growth of HA crystals on various inorganic or organic templates with the surface groups of –COOH, –NH2, and –OH [2+ along the polymer chain by electrostatic and coordination (Ca2+–CONH–) interactions when CaCl2 was added, and then the addition of (NH4)2HPO4 solution promoted the deposition of PO43− ions onto the polymer chain via their electrostatic interactions with –R3N+ groups and Ca2+. Due to the PCBAA-mediated nucleation and growth, HA nanocrystals could be formed quickly (0.5 h). With the increasing mineralization period to several days, HA nanocrystals were further adsorbed and assembled to form rod-like HA crystals. Their study provided new ideas for studying the mechanism of protein-inspired biomimetic mineralization due to the similar chemical structure of PCBAA to proteins.It is possible to directly create HA within a short term (about 30 min) through in-situ biomimetic chemical synthesis, in which the presence of diluted Ca and –OH . By usin and –OH . It was  and –OH  and poly and –OH ,41 as nu and –OH , and fouIn order to mimic the mineralization of HA in biological environment, SBF was developed as an organic-free salt buffer to replace the in-situ chemical synthesis . With thTo accelerate the nucleation and mineralization process, the buffer solution with 1.5-folder ionic strength to SBF (named as 1.5× SBF) has also been widely utilized for biomimetic synthesis of HA crystals ,21,37,44In this section, we would like to present the biomimetic synthesis strategies of HA on both graphene supports and graphene-based nanohybrids.2) on support materials are crucial for the biomimetic synthesis of HA crystals as introduced above in the cases in Part 2. To overcome the problem of hard mineralization towards pure graphene support, GO and chemically modified GO surfaces were utilized for in-situ biomimetic synthesis and mineralization [The surface groups -functionalized GO (GO-Car), as shown in lization b, while lization c. In anoBased on the above case studies, it can be concluded that the functionalization  of graphene supports is the most effective strategy to create biomimetic graphene-HA materials, which showed controllable structure, high mechanical properties, and high biocompatibility that suitable for biomedical applications.In this section, the biomedical applications of biomimetic graphene-HA materials in drug delivery, cell culture, bone repair and growth, biosensors, and anti-bacteria fields are introduced and discussed further.In recent years, great interest has been focused on the drug delivery and release due to increasing demands for cancer therapy . GraphenTo introduce biomolecules into the drug delivery systems, GO-HA nanocomposites have been widely applied in drug delivery due to their unique affinity with biomolecules with a large adsorption capacity and pH-controlled release in various biomedical applications ,64. For 2+ ions were introduced to bind with the remaining epoxy and hydroxyl functional groups of GO sheets through the electrostatic interactions. The PO43− ions were slowly released from the CPDS molecules and then reacted with Ca2+ ions to form HA nuclei through the electrostatic interactions  carbodiimide (EDC) and N-hydroxysuccinimide (NHS) reagents to promote the nucleation and growth of HA on multi-functionalized GO surface (GO-carnosine/HA) . ImproveBiomimetic graphene-HA nanohybrids are excellent scaffolds for the adhesion and growth of various cells due to their strong mechanical properties, high biocompatibility, and unique structures ,70,71,72Previously, Kim et al. demonstrated that biomimetic HA can be formed on GO surface for improving the adhesion and proliferation of osteoblast cells . To enhaRecently, Wang et al. demonstrated the -COOH group-medicated biomimetic synthesis of GO/HA scaffolds a, which By using the mineralization of polymers, biomimetic GO-HA nanocomposites have also been utilized for cell culture ,77. For Due to the high potential of HA for bone tissue engineering and the strong mechanical strength of graphene support, biomimetic graphene-HA materials exhibited wide applications in the fields of bone repair and regeneration ,79,80.3− ions [Wang et al. reported collagen (Col)-mediated biomimetic synthesis of HA on GO support for the formation of bioactive Col/GO-HA composites, which exhibited improved hydrophilic and mechanical properties, as well as the ability to promote the bone repair by testing the osteoblastic cells . Besides3− ions . The mes3− ions . The in In addition, biomimetic graphene-HA nanohybrids can be utilized as regeneration medicines for bone repair ,86. For Due to its unique properties, such as an excellent biocompatibility and a good adsorption ability, HA has become an attractive material for the fabrication of biosensors .2−) by cyclic voltammetry (CV) and square wave voltammetry measurements [2−, respectively.Previously, Pang and co-workers developed a simple and sensitive method for detecting luteolin by using graphene nanosheets (GNs)-HA composite as modified electrode materials . In theiurements . They fo2 was added into the as-prepared GO solution with stirring for 1 h and H3PO4 solution was then added into the mixed solution, which was aged in an air oven at 85 °C for 24 h to obtain RGO-HA nanocomposites, as shown in Recently, Gao et al. designed an electrochemical sensor based on RGO-HA for the oxidation of hydrazine . To fabrIn another study, Alam and co-workers developed a selective chemical biosensor based on RGO-HA nanocomposites, and further investigated the electrochemical detection of Bis-phenol A (BPA) by current-voltage (I–V) measurement . The fabIn addition, based on enzymatic recycling amplification and metal NPs, graphene-HA based electrochemical biosensors have been reported. For instance, Bharath and co-workers fabricated a glucose sensor based on GOx-modified GCE . The as-Antibacterial effects are crucial for evaluating the potential of implant biomaterials. Graphene-based nanomaterials have been widely utilized for long-term antibacterial applications ,95,96.Staphylococcus aureus and Escherichia coli (E. coli). In a further study, they investigated the biomimetic mineralization of HA on PCL-modified GO substrate [Staphylococcus aureus and Escherichia coli indicated that the formed GO/PCL-HA also exhibited a high performance for killing both bacteria due to the release of Mg2+ and Zn2+ from the biomineralized HA.Murugan et al. demonstrated the biomimetic synthesis of HA on GO/CNF nanohybrids for the formation of functional GO/CNF-HA composites, which revealed high mechanical strength and similar characteristics to natural bone . It was ubstrate . The ant2) to improve their antibacterial activity. For instance, Xie and co-workers reported the biomimetic mineralization of CS-functionalized GO (GO/CS) for the formation of HA-like octacalcuim phosphate (OCP) [E. coli and S. epidermidis with bactericidal ratios of 94% and 91%, respectively. It is clear that the addition of AgNPs to the scaffold enhanced the antibacterial activity of OCP-GO/CS materials. It is interesting that both OCP-GO/CS and GO/CS revealed relative antibacterial activity towards E. coli and S. epidermidis, as shown in E. coli, meanwhile 89% and 90% against S. epidermidis, respectively. The excellent antibacterial activity of the created CBB-Ag-OCP-GO/CS composites was ascribed to the synergistic effects of AgNPs, GO, and CS.Biomimetic GO-HA nanohybrids can also be conjugated with other bacteria-killing components (such as AgNPs and SiOte (OCP) . The minStreptococcus mutans [Besides the material effects for antibacterial activity, drugs with antibacterial activity can also be adsorbed onto/into GO-HA nanohybrids for inhibiting the growth of bacteria. In a typical study, Gholigegloo and co-workers reported the biomimetic synthesis of HA on GO-carnosine conjugates for loading an antibacterial drug, ICG, to enhance the antibacterial effects against s mutans . It was 2, ZnO, and SiO2 with graphene-HA hybrids for improved antibacterial activity, and the synthesis of multi-functional materials for simultaneous performances , and the development of novel synthesis methods for graphene-HA materials.In summary, we presented recent advances in the biomimetic synthesis of HA crystals on various graphene supports, and further demonstrated and discussed the potential biomedical applications of biomimetic graphene-HA nanohybrids in the fields of drug delivery, cell culture, bone repair and regeneration, biosensors, and antibacterial materials. Previous studies indicated that the functional modification of graphene with various groups, polymers, biomolecules, cells, and others can mediate the biomimetic formation of HA, as well as improve the mechanical strength, bioactivity, and biocompatibility of hybrid materials. A lot of case studies have proven that both graphene supports and biomimetic HA crystals exhibited their unique effects towards improved biomedical applications. We believe that this overview will be helpful for readers to understand the design, synthesis, and mechanisms of biomimetic minerals on graphene supports and explore the biomimetic materials in biomedical engineering, nanotechnology, materials science, analytical science, as well as energy and environmental science. In our opinion, future studies on the following topics could be studied in depth, for instance, the design of 3D porous graphene-HA based materials for cell growth and drug delivery, the combination of NPs like TiO


Palladium nanoparticles display fascinating electronic, optical and catalytic properties, thus they can be used for various applications such as sensor fabrication. Conducting polymers such as polyaniline have also been widely used in sensor technology due to its cost effectiveness, versatility, and ease of synthesis. In this research, attention was given to unify the exceptional properties of these two materials and construct palladium nanoparticle coated polyaniline films to detect hydrogen and moisture. Electrochemical polymerization of aniline was carried out on gold sputtered epoxy resin boards. Polyaniline film was generated across a gap of 0.2 mm created by a scratch made on the gold coating prior to electrochemical polymerization. A palladium nanoparticle dispersion was prepared using sonochemical reduction method and coated on to polyaniline film using drop-drying technique. Polyaniline only films were also fabricated for comparative analysis. Sensitivity of films towards humidity and hydrogen was evaluated using impedance spectroscopy in the presence of the respective species. According to the results, polyaniline films exhibited an impedance drop in the presence of humidity and the response was significantly improved once palladium nanoparticles were incorporated. Interestingly, polyaniline only films did not respond to hydrogen. Nevertheless, palladium nanoparticle coated polyaniline films exhibited remarkable response towards hydrogen. UV absorption of the Pd nanoparticle suspension after the sonication was compared with the initial solution containing ethylene glycol, PVP and Pd arising from PAni film Fig. .Fig. 10AAccording to the impedance data obtained, the PAni sensor exhibits capacitive behavior hence this sensor can be categorized as capacitive type humidity sensor Fig. . PAni fiIn 1–10 MHz region, some variation of impedance with humidity can be seen in PAni film, however the respective variation was marginal in contrast to that of Pd incorporated PAni film. However, a clear dependence of impedance on frequency with a direct correlation can be seen in Pd incorporated PAni film within the 1–10 MHz region Fig. .Fig. 12aAccording to the Fig. 2 = 0.97) implying that it is more suitable for sensor development in comparison to the sensitivity and linearity at 2 MHz (R2 = 0.94)  Fig. c.Similar to the humidity sensing experiment, PAni films exhibit a capacitive behavior, hence the possible sensing element can be categorized as a capacitive type sensor.2 environment may be the reason for such a deviation [2 was observed. Thus, Pd incorporated PAni film exhibited much superior performance towards H2  for linear curve fitting results over 0.90 for all frequencies. However, the sensitivity drop with the increasing frequency must also be taken into consideration in such an instance. Even though the impedance varied in a narrower range in presence of H2 at higher frequencies (13–15 MHz), a substantial variation in impedance was found as the sensor was exposed to the H2 gas. In detail, at 13 MHz, impedance decreased around 1/5th of its original value in the presence of H2 (11%), and it decreased less than half at both frequencies of 14 MHz and 15 MHz   Fig. . Unlike ies Fig. b.2 gas. The capability of Pd to adsorb hydrogen and its ability to break H–H bond may be the cause for the impedance drop in the presence of H2 [The results evidently revealed that Pd nanoparticles were the key component in detecting Hce of H2 , 31. The2 sensing properties. Hence, the presence of humidity in H2 might be the reason for such observations. Moreover, Pd only also exhibits hydrogen sensing activity and Pd incorporated PAni film shows significant sensing performances towards hydrogen. Fast and easy fabrication and cost-effectiveness would justify the candidacy of Pd incorporated PAni towards sensing both humidity and hydrogen.This study has shown that PAni is a suitable material for the detection of humidity. Incorporation of Pd to PAni increased the sensitivity for humidity. Importantly, PAni film alone did not exhibit H


Graphene has been publicized as the game changing material of this millennium. To this day, scalable production leading to exceptional material properties has been difficult to attain. Most methods require harsh chemicals, which result in destroying the graphene surface. A method was developed, exploiting high speed elongational flow in a novel designed batch mixer; creating a distribution of pristine few to many layer graphene flakes. The method focuses on exfoliating in a molten polyamide 66 (PA66) matrix, creating a graphene reinforced polymer matrix composite (G-PMC). The process revealed that high speed elongational flow was able to create few layer graphene. Graphite exfoliation was found driven in part by diffusion, leading to intercalation of PA66 in graphite. The intercalated structure lead to increases in the hydrogen bonding domain, creating anisotropic crystal domains. The thermal stability of the G-PMC was found to be dependent to the degree of exfoliation, PA66 crystal structure and composite morphology. The aim of this research is to characterize uniquely produced graphene containing polymer matrix composites using a newly created elongational flow field. Using elongational flow, graphite will be directly exfoliated into graphene within a molten polymer. Graphene has been publicized as the leading edge material of this millennia. In the past two decades, the 2-D material has been considered for use in applications, such as thermal management , environAn alternative approach would be to cultivate an exfoliation method, for which the matrix is not a sacrificial phase but rather becomes part of the end use application, for example, a graphene contained reinforced polymer nanocomposite. In order to do that, the matrix must impart significant shear forces to exfoliate flake graphite to graphene nanoflakes. Thus, it is important to study scalable polymer processing techniques that are optimized for efficient graphite exfoliation.2, exfoliation was further improved through bubble formation. The bubbles were able to generate an additional extension mechanism, upon expansion of CO2 to further exfoliate the graphite. Those that are specifically designed for elongational flow start with pre-expanded graphite, attempting to reduce the size of graphite with inherent defects [Direct exfoliation has been known to providing an economical, top-down approach in converting the graphite to graphene . The pro defects . There a defects  uses dua defects  producesSome of the earliest experimental and theoretical work in fluid flow of two immiscible liquids was pioneered by Taylor in his experiments with droplet suspension breakup in a dilute medium of dissimilar viscosity . Taylor p to a droplets critical deformability; where μD is the viscosity of the dispersed phase and μC, the viscosity of the continuous phase (1).Taylor found the mechanism for breakage of droplets to be related to apparent differences in kinematic viscosity ratio defined by crit was needed; where Cacrit < Ca [d the droplet radius (2).The critical deformability was found to work but only up to a maximum viscous ratio of 4. Similar results were confirmed by Rumscheidt and Mason of this critical viscosity ratio to deformation ; suggestrit < Ca . The capp > 1 in Associating elongational flow to simple shear flow figure, Grace found the critical capillary number to be smaller for the case of elongational flow. Leaving for simple shear, the critical capillary number for breakage approaching a near infinite value as the viscous ratio exceeds The aim of this research is to characterize uniquely produced graphene containing polymer matrix composites using a newly created elongational flow field. Using elongational flow, the drive is to exfoliate graphite into graphene directly within a molten polymer. This work is an expansion of a method to create graphene reinforced polymer matrix composite by Nosker et al. . In usin3 s−1.In this work, we use elongational flow to exfoliate 35 wt % mined graphite within polyamide 66 (PA66) to create graphene containing polyamide 66 nanocomposites (G-PA66). The components are subject to 10, 20, 30 and 45 min of elongational flow mixing. The mixed components are then characterized for their changes in structure, morphology and thermal properties. Using this single step approach to convert graphite directly to nanographite and graphene, it is suggested that a molten elongational phase can apply unique hydrodynamics on the graphite flake to produce exfoliated few, many layer graphene and nanographite at shear strain rates exceeding 103, Asbury Carbons, Asbury, NJ, USA) was used as the exfoliating species. Polyamide 66  was used as the high temperature polymer in this study. Prior to processing, both components were dried in an oven to eliminate water. PA66 was dried at 85 °C under vacuum and graphite was dried at 350 °C for 4 h.Natural flake graphite  was performed using a Philips X-Pert Powder X-ray Diffractometer  with a Cu (λ = 1.54 Å) K-alpha source at 45 kV/40 mA and scanned from 4° to 70° 2ϴ. The samples were scanned with a step size of 0.02° 2ϴ to detect graphitic and nylon characteristic peaks. Extrudates were notched and cryogenically fractured to produce a pristine surface. Tm), crystallization (Tc) and glass transition (Tg) temperatures of the PA66 component was determined, to include enthalpy of crystallization ΔHc, enthalpy of melting ΔHm and % crystallinity in PA66. Thermal properties of PA66 and G-PA66 samples were determined using a Q1000 Differential Scanning Calorimeter . PA66 and G-PA66 samples weighing 10 mg, were sectioned from the extrudate and subject to a heat, cool, reheat method at 10 °C/min from 0–300 °C under a nitrogen atmosphere. The effect of processing on the melting (Morphology of G-PA66 samples was investigated using a Sigma Field Emission Scanning Electron Microscope (SEM)  with Oxford EDS.In order to better investigate the nanoflakes produced during this in situ exfoliation method, G-PA66 nanocomposites were granulated using a Spex 6700 Freezer Mill  to cryogenically impact samples and form a fine powder for viewing using the SEM. The fine powder was ultra-sonicated in isopropyl alcohol (IPA) for 5 min, then the suspension was drop coated on to a lacey carbon grid for transmission and characterization in a JEOL 2010F Transmission Electron Microscope  operating at 200 kV for High Resolution Transmission Electron Microscopy (HR-TEM).K = 0.9 represents the shape factor, λ is the incident radiation wavelength, β is the line broadening at half the maximum intensity of the detectable peak and ϴ is the Bragg angle in degrees.In our assessment of the crystalline characteristics of the X-ray diffraction results, the Scherrer Equation (3) was used to calculate the crystalline domain size of the phases in PA66, G-PA66 and Graphite. The starting graphite is analyzed to having a (002) lattice spacing of 0.334 nm, indicative of the stable hexagonal phase . After eIn The results of the curve fit for the diffusion model shows that the linear equation is second-order with respect to mixing time. Second-order approximations accounts for a linear driving force (LDF) for diffusion a graphite intercalant, from Budzianowski . In his Diffraction patterns of PA66 under extensional mixing are shown in XRD patterns of PA66, relative to the extensional mixing with graphite, appear in Extending the time of mixing from 30 min to 45 min with graphite, produces a distinctive change in the G-PA66 lattice structure. G-PA66 shows a lattice increase for both the (100) and (010)/(110) by +0.0030 nm and +0.0037 nm, respectively. This material also changes with a reduction in the hydrogen bonded (100) domain and an increase in the intersheet (010)/(110) domain. The change in the van der Walls (010)/(110) direction trends parallel to the lattice changes in graphite (002). The reduction in (100) d-spacing is likely due to disruptions in the hydrogen bonded network due to graphene and nanographite lattice mismatches, that may increase available surfaces for binding [It is likely that the lattice parameters in the hydrogen bond and the van der Waals direction follows not with the degree at which the graphene is reduced in size but with the (002) lattice spacing between graphene layers beyond 30 min of mixing. The stacking faults in exfoliated graphite invoke changes to the lattice parameters in PA66, which lead to parameter increase. These faults are the results of created turbostratic graphite , consistTg). Transition temperatures for 10, 20, 30 and 45 min of exfoliation mixing of PA66 & 10, 20, 30, 45 min G-PA66 were found. The glass transitions (Tg) were found to occur at 60, 56, 57, 53, 56, 57, 51, 50 °C respectively. For all the values listed in m = 197 J/g [m were adjusted for percent PA66, since G-PA66 samples only have 65 wt % PA66.Thermal properties for PA66 and G-PA66 after 10, 20, 30 and 45 min of elongational mixing appear in  197 J/g . The valTm. As the mixing time increases, Tg, Tm, Tc and ΔHc all decrease. However, there is no consistent trend with ΔHm results in In Tc) and percent crystallinity for 10 and 20 min of mixing; symptomatic to heterogeneous nucleation of PA66 in graphene/nanographite presence. As the exfoliation time continues beyond 30 min, there a reduction in Tg, Tm and Tc. These properties are much lower than all the other materials, suggesting a drastic change occurs at 30 min and beyond, in its structure.In Tm and Tc. The reduced crystallization previously mentioned, suggests the suppression of crystallization to occur. The exfoliated graphitic nanoflake in the matrix is likely to cause a subdued rate of crystallization [Tc of 9 °C. The same results appear for the Tg, with 45 min of exfoliation resulting to a 6 °C reduction than the lowest Tg. The crystallization suppression is shown sensitive to the percent crystallinity calculated, which suggest lower Tc depression leads to lower crystallinity.Compared to pristine PA66 at 30 min, G-PA66 at the same exfoliating time, has nearly equivalent values of percent crystallinity but has a much lower lization ,35. ThisIdentifying how the crystal structure relates to the thermal stability of G-PA66 composite, we see that the hydrogen bonded lattice parameters play an important role. This is not represented in the neat PA66, since the differences between 20 and 30 min are nominal. G-PA66 mixed at 20 min; having the least modified hydrogen bonded lattice, has relatively higher thermal properties and percent crystallinity. The thermal property variations occurring through reduction of the hydrogen bonded lattice, suggesting there is a relation to the structure in the PA66 of the composite. From the increased crystallization temperature Tc and percent crystallinity, we know this is evidence of surface crystallization occurring in the G-PMC (10 & 20 min). The cause of crystallization temperature depression in G-PA66 mixed at 30 min might be explained by the exfoliated graphene acting as unstable crystals growth surfaces. These surfaces form disrupted hydrogen bonded sites; imparting themselves between the formed crystals, suppressing their crystallinity and thermal stability .In all stages of the exfoliation process, it is important to note that there exist both nanographite and many layer graphene flakes, even though the average graphite domain is calculated to be 14.8 nm. This is suggested, due to the fact that few to many layer graphene crystallites in the (002) direction are weakly diffracted. The weakly diffracted domains make the crystals unresolved from the background. That being said, the first thing to notice is the lack of large flakes of graphite present on the fractured surface. The lack of visibly large flakes suggest that nearly all the graphite were reduced in thickness but retained relatively large surface area in When looking at To observe the created graphene in the absence of the extrudate matrix, cryogenic milling was performed on the 30 min G-PA66. In What is interesting in What is difficult to find, is a distinct indication of the PA66 with the graphene in The first thing noticed from the HR-TEM analysis is that from the low magnification image, the process produces graphene and nanographite composite flakes of varying sizes Taking a step back and looking at the arrangement of the flakes, many layer graphene and nanographite are shown to be orientated in multiple directions in Tm and Tg. Due to these changes, the graphene exfoliate is found suppresses the crystallization in PA66, leading to a modified microstructure. In our novel process, elongational exfoliation in the molten phase led to the formation of pristine, large area nanographite and multilayer graphene flakes. These results are confirmed from HR-TEM and morphology in electron microscopy. From HR-TEM and XRD, the resulting process identified that the graphite was exfoliated, leading to the disruption of the graphene layers in the composite and a 76% reduction in the initial (002) domain size. In-situ, a misalignment is described by a rotation of the graphene layers and conversion to a turbostratic phase. The model of graphite exfoliation revealed that the process is diffusion driven, with intercalation of the PA66 within the graphite galleries. In addition to diffusion, the process is found to tear and fold the outer layer sheets in graphite creating instance for the formation of dangling bonds. The results of which, preclude to a solvent-less method to functionalization graphene. The morphology created changes in the hydrogen bonded structure, given by changes in the lattice and domain in the (100) and (010) crystalline plane. Changes lead to preferential crystal growth in the hydrogen bonded direction, resulting in changes to the The G-PMC created by the process, shows graphene to be created due high shear. The results show a single step approach to create as well as modify graphene. Further investigation into this work may reveal a better description of the interfacial properties of these two species and the type of bonding present. This may provide ways of modifying the interface functionality in a scalable process and further exploiting the capabilities of elongational flow. This in-turn drives down cost, which is an important variable in adaptation and scalability.


Electrochemical exfoliation is a promising bulk method for producing graphene from graphite; in this method, an applied voltage drives ionic species to intercalate into graphite where they form gaseous species that expand and exfoliate individual graphene sheets. However, a number of obstacles have prevented this approach from becoming a feasible production route; the disintegration of the graphite electrode as the method progresses is the chief difficulty. Here we show that if graphite powders are contained and compressed within a permeable and expandable containment system, the graphite powders can be continuously intercalated, expanded, and exfoliated to produce graphene. Our data indicate both high yield (65%) and extraordinarily large lateral size (>30 μm) in the as-produced graphene. We also show that this process is scalable and that graphene yield efficiency depends solely on reactor geometry, graphite compression, and electrolyte transport. Such materials are used for energy storage, structural composites, and advanced electronics9. Industry reports estimate that graphene demand will reach over 4100 tons/year by 202610.Graphene, i.e., a single layer of sp11. Although a wide range of graphene production methods have been studied, the difficulty of producing large quantities of graphene from flake graphite has been a persistent problem. Although bottom-up methods of growing graphene sheets on substrates through vapor deposition have been successfully applied in industry, these techniques are limited to large-area graphene but cannot produce large masses of graphene13. The most commonly used top-down methods for producing graphene from graphite suffer from a longstanding quality vs. quantity tradeoff14. There are two large-quantity methods of producing graphene from graphite: (i) The oft-used modified Hummers’ method involves extensive oxidation16, but the resulting graphene oxide (GO) nanosheets are defect-laden and electrically insulating. Graphene oxide may be thermally or chemically reduced to yield “reduced graphene oxide” (rGO) nanosheets that retain some of their defects and oxygen-containing functional groups20. (ii) In contrast, pristine graphene is typically mechanically exfoliated from graphite, often in a liquid medium with sonication or shear mixing, but the actual product is a mixture of pristine graphene and unexfoliated graphite31. Pristine graphene production has been plagued with difficulties either in production rate or in the number of layers present in each graphene platelet. This has led to much confusion in metrics and nomenclature; both “graphene nanoplatelets” and “nanographene platelets” are commercially available (both many-layers). Several recent reports have described advances in scalable (better than linear) methods for production of few-layer pristine graphene from graphite21. However, this analysis invariably neglects the problem of separating the as-produced graphene from the unexfoliated graphite; such separation techniques are notoriously difficult to scale up32. If high-yield production methods can be attained, the separation problem becomes far less pressing.A variety of graphene family materials are a major target for industrial scaleup as high-value-add fillers and additives35. An applied voltage drives the ionic species in an electrolyte to intercalate into the graphite electrode and increase the inter-layer distance36. For example, in ammonium sulfate, the sulfate ions and water molecules migrate into the interstitial regions of the graphite and locally form gas bubbles , which forces adjacent sheets apart  Only graphite monoliths (as opposed to loose graphite powders) may be used as a source for electrochemical exfoliation because the graphite electrode must be continuous, electrically conductive, and connected to the external power supply. (ii) Most importantly, if electrochemical exfoliation actually occurs and the monolith is expanded to form graphene, then the monolith disintegrates and the electrochemical exfoliation process stops. The graphite electrode tends to fall apart as intercalation proceeds, and when graphite pieces become disconnected from the monolith , they can no longer participate in the intercalation process. The degradation Fig.  of the get al. attempted to address these challenges by placing the monolithic graphite electrode at the bottom of the exfoliation vessel, with the expectation that gravity would allow the disintegrated graphite particles from the monolith to remain connected by gravity, even as exfoliation progresses34; they reported a slight improvement in yield (based on sonicated dispersions in dimethylformamide), but graphitic particle retention is uncontrolled. Abdelkader et al. proposed (but have not attempted) a similar setup where a steady flow of electrolyte solution is continuously fed through the setup, allowing any exfoliated, floating graphene to be collected37. Wang et al. demonstrated that paraffin coating outside the graphite electrode could effectively confine the electrochemical exfoliation to a defined volume38. However, intensive sonication in N-Methyl-2-pyrrolidone and further centrifugation were used as post-processing in this case. We discuss the “yield” of these processes in further detail below.Liu 39. As shown in Fig. 4)2SO4). When a positive voltage (+10 V) is applied to the working electrode, the graphite flakes are continuously intercalated and expanded.Our method overcomes these challenges and provides a new approach for scalable, high-yield graphene production with large sheet size−) attack graphite at the edges and grain boundaries41; the oxidation at the edges and grain boundaries expands the graphite, facilitating the intercalation of sulfate ions (SO42−); the reduction of SO42− anions and oxidation of water generate various gases, such as SO2, O2, and others. The expanding gases separate the graphene layers, resulting in electrochemically exfoliated graphene (EEG). The counter electrode does not experience any exfoliation. This method is suited to a range of precursor graphitic materials, unlike prior electrochemical methods that require a monolithic graphite structure.Because of the compression of the container, the graphite flakes remain in electrical contact, even as they expand due to gas evolution between atomic layers of graphite flakes. 42. This improved wetting leads to improved ionic transport.The wetting of the graphite by the electrolyte may be hindered by graphite’s natural hydrophobicity since the reaction involves oxidation only at wetted graphite surface sites. We used HNOtes Fig. 42. This ). (Shear mixing time and speed were varied and are discussed below).In our batch reactor schematically depicted in Fig. S3−1, a broad G peak at 1580 cm−1, and 2D peak at 2680 cm−1. The shape of the 2D peak identifies the number of layers in graphene43. The G peak corresponds to the first-order scattering of the E2g mode of sp2 carbon atoms. The prominent D peak is from the breathing mode of the sp2 carbon atoms and activated by the existence of defects, including edges, functional groups (hydroxyl and epoxide groups), and structural disorders. The shape of the 2D peak identifies the number of layers in graphene43. The ratio of D peak to G peak (ID/IG) varies from 0.9 to 1.2 in different spots. Typically, an ID/IG ratio of 1.2 is very similar to chemically or thermally reduced GO (~1.1–1.5)44. The symmetric 2D peak of our EEG-2 is quite similar to previously reported few-layered graphene. Note that the shear mixing itself does not appreciably alter the Raman signature of the material . X-ray photoelectron spectroscopy (XPS) was used to probe the chemical composition of the EEG-2  analysis reveals a similar trend  to study the sheet size and thickness of EEG-3. Water-suspended EEG-3  was drop-casted on mica directly for AFM Fig. . The heiial Fig. . X-ray pG-2 Fig.   and this buckypaper is termed as “EEG-5”. This buckypaper showed an electrical conductivity of 2.1 ± 0.2 × 10ods Fig. . We alsoods Fig. .In order to probe the morphology of the produced EEG-2 and EEG-4, scanning electron microcopy (SEM) was performed. Figure 36. This is somewhat misleading, because after the yield is determined, additional sonication and centrifugation (i.e. separation) is needed to produce graphene such that only a fraction of that material actually becomes few-layer graphene. In contrast, we define our yield as the EEG-1 produced after electrochemical exfoliation and sedimentation, with no additional separation steps that reduce the yield; our characterization data above indicate that all of the as-produced EEG-1 can be considered graphene . Our prototype batch reactor results in yields of 38% (only 15% if the graphite is not pre-treated with HNO3). Below, we show that this number can reach 65% and possibly higher.Yield is one of the most critical metrics for scalable production. Note that the field of graphene production has suffered from a lack of standardization on this point. For example, prior efforts have calculated yield as simply the mass liberated from the graphite electrode divided by the original graphite mass), and we assume that the electrical resistance of the Pt wire is negligible. As R increases, the average voltage drop from the Pt wire to any given graphite location increases and the efficiency (η) is lowered. (ii) Similarly, ion diffusion into the interior may be the limiting factor; for isothermal ion diffusion-limited reactions, the efficiency η scales with R−1 (regardless of geometry or reaction order) because the Thiele modulus scales linearly with R45. Note that this problem of ion diffusion into the interior is exacerbated as graphite is converted to graphene and the tortuosity increases inside the electrode. .The working electrode (graphite flakes compressed in a permeable container) can be modeled as an electrically conductive cylinder Fig.  where thS1With these limitations in mind, the electrochemical reactor was redesigned Fig.  to incre). These graphite flakes would be in contact with an in-line screw conveyer that doubles as the working electrode, and are continually electrochemically exfoliated as the graphite moves through the reactor. The average exfoliation residence time τ in the continuous reactor is controlled by the graphite flow rate Q and the geometry  of the reactor as τ = AL/Q. This is directly scalable. The flowing graphite must be at a high enough concentration to ensure electrical contact, such that the flow is a dense slurry. This is only one embodiment of the possible continuous electrochemical reactors that uses this compression technique, and many potential designs exist that could lead to efficient continuous graphene production. Additional scaleup factors include: Number, geometry, and material of electrodes and current collectors; parent graphite particle size, and electrolyte distribution systems.In addition to batch processing, a continuous electrochemical reaction scheme is proposed in which flake graphite is conveyed through a screw reactor and electrolyte is forced under pressure through a porous pipe or ridged mesh Fig. . These gWe successfully demonstrated that graphite powder can be electrochemically exfoliated into graphene in a permeable and expandable container as a monolithic working electrode. Raman spectra and other characterization data suggest a high yield of graphene with large sheet sizes (>30 µm). In addition, various sources of graphite flakes/powder can be utilized as the parent material using our approach. We also demonstrated design principles for graphene reactors, including geometric effects on efficiency and electrolyte transport; these design principles indicate that scaleup of continuous graphene production can be accomplished without compromising yield or quality.The datasets generated during and/or analyzed during the current study are available from the corresponding author on reasonable request.4)2SO4) aqueous solution as electrolyte. A flat piece (~10 cm × 10 cm) of graphite foil serves as the anode. The cathode and anode are connected to a power supply. 10 volts are applied to the platinum electrode immersed in dialysis bag filled with graphite flakes during electrochemical expansion, with a compliance current of ~0.5 A. Electrochemical expansion time varies from 1–24 hours. The EEG is then washed via centrifugation with distilled water for five times. The EEG can be distinguished from flake graphite by its matte black color and high porosity. The yield is calculated by freeze dried weight of EEG-2 collected over the initial weight of graphite placed in the tubing. After washing, the supernatant was transferred into a 400 mL Erlenmeyer flask, and the flask was filled with 300 mL of DI water. After settling for 10 minutes, the supernatant was transferred into a beaker, labeled EEG-1 for further characterization and treatments. The sediment was unexpanded graphite.5 g of flake graphite (Sigma-Aldrich) is compacted in dialysis tubing  with a platinum wire inside as the current collector.  The dialysis tubing is clipped on both sides to maintain the integrity. This assembly serves as the cathode. The compact cathode is immersed in 3 L of 0.1 M ammonium sulfate  (Bruker Dimension Icon) operated in tapping mode and a scanning electron microscope (SEM) (JEOL JSM-7500F instrument) at 5 kV accelerating voltage. The EEG-1 and EEG-3 was dropcasted on freshly cleaved mica surface for AFM imaging. The EEG-2 and EEG-4 was pressed on carbon tape for SEM imaging. X-ray photoelectron spectroscopy (XPS) was conducted using an Omicron XPS system with Mg X-ray source and the data analysis was performed using CasaXPS software . The EEG-2 was pressed on carbon tape for XPS. The Raman spectra were obtained using a ThermoScientific DXR Raman confocal microscope with 900 g/mm gratings for excitation lasers with a wavelength of 532 nm. The EEG-2 and EEG-4 was pressed on a piece of glass slide for Raman spectroscopy. For bright-field optical microscopy, the EEG-1 and EEG-3 sheets were suspended in DI water and deposited on a glass slide. X-ray diffraction (XRD) patterns were collected on Bruker D8 Discover diffractometer fitted with LynxEye detector in a Bragg Brentano geometry at 40 kV, 40 mA with CuKα (λ: 1.5418 Å) radiation. The EEG-2 was pressed on a piece of glass slide to obtain the XRD patterns with a scan rate of 1 s per step and a step size of 0.02°. Thermogravimetric Analysis (TGA) was performed in a TA Instruments Q50 TGA to determine the mass loss of the EEG-4. ~10 mg of each sample was heated up in a ceramic pan from room temperature to 1000 °C at a rate of 1 °C/min in air. The electrical resistance of the EEG-5 was measured using the four-point probe method. The four-point probe head  was mounted on a resistivity measurement stand . The spacing between the probe tips was 1.5875 mm. The current was passed to the sample through the outer probes using a Keithley 6221 AC and DC current source. A Keithley 2000 multimeter was used to measure the voltage across the sample. The sheet resistance and electrical conductivity of the samples were calculated using the measured voltage values.Supplementary Information