The resulting seven hybridized Li-C orbitals that support nine possible bond pairing amplitudes. The gap equation is solved for all possible gap symmetries. One band is weakly dispersive near the Fermi energy along Γ → M where its Bloch wave function has linear combination of dxy character, and is responsible for dxy pairing with lowest pairing energy in our model. These symmetries almost preserve properties from a two band model of pristine graphene. Another part of this band, along K → Γ, is nearly degenerate with upper s band that favors extended s wave pairing which is not found in two band model. Upon electron doping to a critical chemical potential μ1 = 0.22 eV the pairing potential decreases, then increases until a second critical value μ2 = 1.3 eV at which a phase transition to a distorted s-wave occurs. The distortion of d- or s-wave phases are a consequence of decoration which is not appear in two band pristine model. In the pristine graphene these phases convert to usual d-wave or extended s-wave pairing.A study of possible superconducting phases of graphene has been constructed in detail. A realistic tight binding model, fit to ab initio calculations, accounts for the Li-decoration of graphene with broken lattice symmetry, and includes  SrTiO3 and related substrates2. With the enormous research activity focused on graphene in recent years, it is not surprising that graphene-based superconductivity has become an active area of research. Very recently superconductivity up to 1.7 K has been reported3 in magic angle bilayer graphene, which will buttress activity on two dimension superconductors and especially the related type that we discuss here.Two dimensional superconducting phases have become of great interest since the discovery of the high temperature superconducting (HTS) cuprates and subsequent finding of Fe-pnictide and -chalcogenide HTSs. Interest was re-invigorated by the discovery of superconductivity onsets up to 75 K in single layer FeSe grown on C6Ca and C6Yb4. With the many remarkable properties of graphene, it has been anticipated that doping by gating or by decorating with electro-positive elements, thereby moving the chemical potential away from the Dirac points, might induce superconductivity. However, graphene decorated with alkali metals has three valence bands with one weakly dispersive band near Fermi energy. Due to this flat band, there are additional available states around the Fermi level and the required pairing potential is reduced.Superconductivity has been known for some time in intercalated graphite compounds such as et al. based on angle-resolved photoemission spectroscopy5 (ARPES). Scanning tunneling spectroscopy (STS) was applied by Palinkas et al.6 to graphene suspended on tin nanoparticles, who concluded that superconductivity is induced in the graphene layer. Evidence of superconductivity in Li-decorated few layer graphene at 7.4 K has been reported by Tiwari and collaborators7. Low temperature mobility of K and Li atoms on graphene was observed by Woo et al., and suggest that mobility may persist at lower temperatures8, which would provide new challenges for theory.Discussion of superconductivity in doped graphene has been primarily within theoretical models, as we review below, but some encouraging data have been reported. Experimental evidence for a superconducting gap in Li-decorated monolayer graphene around 6 K has been reported by Ludbrook 9. Repulsive electron-electron interactions were modeled by Nandkishore and collaborators11. Beginning from pristine graphene, varying the chemical potential leads to dominant chiral singlet et al.12 a triplet f-wave state has been proposed to arise from next-nearest neighbor interaction with chemical potential near van Hove peak13. Both chiral and conventional p-wave states in graphene have been discussed14, with the many pictures raising various possibilities but little of a certain nature.Various mechanisms of pairing have been proposed. Uchoa and Castro-Neto modeled pristine and doped graphene with electron-phonon coupling or plasmon mediated in mindet al. predicted15 based on Eliashberg theory that decoration by electron donating atoms such as Ca and Li would make single layer graphene superconducting, with modest critical temperatures in the 1–8 K range. In somewhat related work, Wong et al. have predicted16 from an ab initio treatment a critical temperature around Tc = 14 K for carbon nanotubes, which was increased to above 100 K for a certain type of carbon ring.More specific predictions have begun to appear. Profeta 6 for illustration, each cell site has six C atoms in a hexagon with an alkaline atom lying above the center of the hexagon. The C π orbitals and alkaline atom’s s-orbital hybridize to give seven “molecular” orbitals. For two dimensional graphene-like structures effects, differences in nearest neighbour hopping integrals affect the band structure near the important Dirac point, which is folded back to the Γ point of the shrunken graphene superlattice investigated by Hou et al.17 and Long-Hua et al.18. For such systems not even the full analytic tight-binding band structure has yet been reported. The intent here is to extend study of this system, with representative LiC6, from the underlying electronic structure to investigation of the possible superconducting phases.Expectations of adjusting the chemical potential include gating, but the main focus has been on decoration of graphene by electropositive atoms, viz. alkalis or alkaline earths. Charge migration from such decorating atoms to the graphene layer will affect the C-C bonding, leading to contraction or expansion of the graphene hexagons that are centered by the decorating atoms, thus breaking the symmetry of C-C hopping integrals around the honeycomb loop. This asymmetric graphene layer will be referred to in this paper as “shrunken graphene”. Taking LiCThe organization of the paper is as follows. In Sec. II the interacting seven orbital model Hamiltonian is presented. The exact band structure of the normal state of this shrunken graphene system is described in Sec. III. Perturbation theory is applied to obtain the band structures in analytic form. Applying the Hubbard model and minimizing free energy of the superconductor state, we obtain in Sec. IV the gap equations and approximate critical temperature. These equations are solved analytically to establish the possible pairing symmetries and other properties of the superconducting states. A summary is provided in Sec. V.6, as illustrated in Fig. hz = 1.85, somewhat smaller than the value 1.93 Å obtained by Profeta et al.15. The nearest Li-C distances are h = 2.40. Since the Li 2s orbital energy is higher than the C 2pz orbital, charge transfer occurs. It is calculated that 0.685e from Li transfers to the six C atoms equally19. The positive Li ion and negative C ion provide a relative Coulomb (Madelung) shift in site potentials of the two atoms.Because the unit cell contains several atoms with important specific aspects, we provide many of the details of the expressions that can be obtained analytically. LiCa1 = 1.425, while the bond length of nearest neighbor C atoms in different hexagons is slightly larger at a2 = 1.426. For Ca instead of Li, this difference should be larger, hence we keep these lengths distinct. The hopping integral between short-bond carbons is t1, with that between stretched carbon sites is denoted 6 compared to graphene. Symmetry breakdown leads to the opening of a small energy gap at the Γ point.The attractive interaction between Li and C ions after charge transfer contracts the Li-C distance and reduces the C-C bond lengths in the Li-centered hexagon to A1, A2, A3, B1, B2, B3 and Li, as illustrated in Fig. The lattice then becomes a two dimensional hexagonal Bravais lattice with seven atomic sites. These will be labeled as HN and HP denote the non-interacting and interaction Hamiltonians respectively. In these expressions α and β run over Ai, Bi and Li. Here σ on subsite α of ith lattice site, and μ0 and tiα,jβ is the hopping integral from the α site of ith cell to the β site of jth cell. We denote the on-site energy by εα.Here U) Hubbard model, which allows a variety of phenomenological values to be chosen and studied. It is largely for this reason that we provide substantial detail of the underlying, non-interacting C-Li lattice and electronic structure. The interactions that we study are introduced in Sec. IV.The interaction stated above corresponds to an extended , this band plays a major role in the formation of superconductivity in graphene. Also, one can observe that the band structure is four-fold degenerate at the charge neutral Dirac points. Solution of the Schrödinger equation for pristine graphene in the mini-BZ has another advantage: the Bloch-wave symmetry character of each branch can be distinguished. The Bloch coefficients of the branch labeled by Eγ are of s-wave character, Eα and Eβ are of the form d ± id -wave i.e. Exact analytical solutions for pristine graphene wherein next neighbor hopping integrals are taken into account are presented in Supplementary Materials Eqs entclass1pt{minimaLiC6 lattice bases length Eg = 0.36 eV. Another gap can arise at the Γ point because of symmetry breaking of on-site energies εA ≠ εB, seen from Eq. εA − εB|. In Li decorated graphene that we consider here, all carbon on-site energies are equal so this type gap does not arise.Decoration of graphene with metals reduces symmetries that lead to removal of bands degeneracy in some regions. While decoration causes expansion and contraction of bonds length in three inequivalent directions in the honeycomb lattice i.e. s-wave or chiral d ± id-wave and there are no mixed states, when symmetries in hopping integrals are broken by decoration, Bloch functions are linear combinations of all these phases, Eq. While for folded but pristine graphene Bloch wave solutions are pure ELi,0(0) and the lowest valance band, Esh,6(0), are mutually affected. The energies of these bands are, with E0(0) ≡ E+, E6(0) ≡ E−,Li-C hopping adds a perturbation term to the shrunken graphene Hamiltonian. Obtaining exact dispersions from Eq. M point. We define a hopping integral symmetry breaking index, wt open a small gap Eg at Γ, which does not depend on N(0) changes with decoration.There are two critical points in the pure graphene band structure which are affected by decoration and become important: the charge neutrality Dirac points folded at the Γ point, and the van Hove singularity at the 6 was fit to the DFT band structure, with results illustrated in Fig. A1 subsite chosen as central site labeled by 0 and B1 subsite in adjacent hexagon considered as second neighbor while just slightly longer than the first neighbors atoms B2 and B3 in same hexagon, this neighbor labeled by n = 2 and so on the next neighbors are labeled. In Fig. nth neighbor has been shown by m = 4 neighbors. Li to C hopping integrals are very small with respect to those of C-C and Li-Li, so we keep only the near neighbor Li-C hopping amplitude.The seven band tight binding model of LiC15. The fitted hopping amplitudes and on-site energies are presented in Table 6 with pristine graphene in ref.23, it is observed that Li decoration only slightly changes the pristine graphene band structure. These changes are due to electron transfer from Li to graphene, which changes the pristine on site εpristine = 0 to εA = εB = εc.Since Li is small with respect to alkaline earths such as Ca, the pristine band structure is less affected by decoration by lithium than by calcium, as can be seen in Fig. 6 presents a multiband system in which three bands cross the Fermi level. We presume singlet pairing that can be both intraband and interband in nature. We adopt a local viewpoint in which pairing occurs between electrons on carbon atoms. Seven hybridized Li-C orbitals, support nine possible bond pairing amplitudes in real space. Figure s = 1 is for particles and s = −1 for holes, and Em are the normal state eigenvalues. The m and n are band indexes. The band pair order parameter m-th and n-th bands in LiC6. Also, j-th band. Possible order parameter symmetries in Eq. LiCentclass1pt{minimaThe linearized gap equation, obtained by minimizing the quasiparticle free energy with respect to nearest neighbor order parameters, isij〉 has been dropped for brevity. The A3×3, B3×3, C3×3, and D3×3 matrices, given by Eq. This equation can be written in matrix form asmentclasspt{minimag0 = g1, the Bloch wave coefficients appearing in Eq. C3×3 = A3×3 and D3×3 = B3×3. Eq. For folded six band pure graphene V1 = V2 = V3 = Vsy where sy subscripts indicates each of the s, dxy or 12. These three solutions preserve symmetry of the two band unit cell as illustrated in Fig. n0 = (Vsy 0 −Vsy) and Φn1 = (Vsy −Vsy 0) that break symmetries of pristine graphene two band model. Inserting these solutions into Eq. g0. In the following section the superconducting gap equation has been solved for LiC6 and it is demonstrated how Li-C coupling influences superconducting phases.For the case Vsy refers to one of the Vs, Self-consistent solutions of the gap equation Eq. asy, bsy, csy, and dsy by ad(s), bd(s), cd(s), and dd(s) respectively.In these expressions αsy coefficient appearing in the pairing amplitudes of stretched bonds as shown in Eq. By comparing the gap equations introduced in Eqs αsy = 1 as illustrated in Fig. κ = 1, asy = ccy and bsy = dcy from Eq. bsy > 0 then The six bands of pristine graphene support nine pairing amplitudes while in the two band model there are three possible pairing amplitudes along three different bonds. These two notions can be mapped onto each other only if 3 electron pairing states with island character and very high pairing potential; 3 states with higher electron pairing potential; 3 states with lower electron pairing potential; κ = 1, i.e. g1 = g0. Only for this case the matrix gap equation becomes Hermitian, then band order parameters takes the following form in terms of the band Green function and g0,All states are orthogonal except those with same subscript, viz. σ in the ith band with energy g1 = g0 but deviation from pristine leads to distortion of Green’s functions Here g0 in the (+) states is lower than g0 in the (−) states hence pairing in this modes are dominant. From Eqs Phases 15. Some have suggested phonon-mediated superconductivity in single layer graphene. Most prominently, Profeta et al.15 calculated on the basis of density functional theory for superconductors that decoration by electron donating atoms such as Ca and Li will make single layer graphene superconducting, up to 8 K for the case of Li. The ab initio anisotropic Migdal-Eliashberg formalism was used by Zheng and Margine24, who predicted a single anisotropic superconducting gap with critical temperature Tc = 5.1–7.6 K, in surprisingly good agreement with experimental reported superconductivity around 6 K in LiC65.The possibility of a superconductivity state in metal decorated graphene has been suggested theoretically by a few groups9 and also by Black-Schaffer and Doniach12. The possibility of a singlet p + ip phase pairing near the Dirac points between nearest neighbors subsites were suggested by Uchoa and Castro-Neto9. They worked in terms of a plasmon mediated mechanism for metal coated graphene, and discussed the conditions under which attractive electron-electron interaction can be mediated by plasmons.Using a phenomenological microscopic Hamiltonian in a nearest-neighbor tight-binding approximation, possible superconducting phases of pristine graphene have been discussed by Uchoa and Castro-Neto12. For the nearest neighbors pairing amplitudes iA site to its three nearest neighbors, it was observed that there are three states that minimize the free energy in various regimes of the parameters, which here have been denoted by Vs = T, dxy and d + id or s states tend to dominate. Their numerical calculation showed that d-wave solutions will always be favored for electron or hole doping in the regime M corresponding to 3/8 and 5/8 fillings i.e. d + id superconductivity, which breaks time-reversal symmetry, can be stabilized. In this regime d-wave superconductivity may arise from repulsive electron-electron interaction11.Singlet superconducting gap phases of pristine graphene have been proposed and discussed by Black-Schaffer and Doniachs-band and graphene π bands changes the graphene band structure. The s orbitals of Ca have more overlap with C orbitals than Li and lead to stronger and longer range interactions as well as increasing the doping level, effects that become detrimental to superconductivity. For this reason our emphasis here is on the Li decorated graphene.Although doping by a gate voltage is normally considered to change only the chemical potential but not the band structure, gating cannot be expected to push the Fermi energy to the van Hove singularity without altering the band dispersion. The most likely way to do this is by decoration with electropositive atoms, which has been our focus. We note that doping is essential, when graphene decorated, in addition to the expected charge migration from the decorating atoms to the graphene sheet, it is then necessary the interlayer state is partially occupied to induce superconductivity as happens in GICs. Hybridization of interlayer 6 and the usual two band model of C2 were compared. These two viewpoints coincide if the same pairing paradigms are considered. For pristine graphene with its two site cell, in real space picture electrons can pair with near neighbors in three inequivalent directions, Vsy quantities are the three vectors that belong to the irreducible representation of crystal point group Dh6 i.e. sy subscript stands for symmetries s, dxy and s-wave solution  along three different bonds constructs just one state while permutation of dxy solution  up to a minus sign constructs two nonorthogonal linear independent states viz.  and  which orthogonal linear combination of them are We review some of our main points. When graphene is decorated by Li, electron transfer from Li atoms to C contracts the Li-C distance and reduces the C-C bond lengths in the Li-centered hexagon. In this kekulé -type structure, hopping amplitude symmetries of all C-C neighbors are broken (our “shrunken graphene”). This model allows study of multiband effects on the superconducting phase diagram. To gain insight into our model, solutions of superconducting gap equation in both cases of folded bands otherwise pristine C6 includes six carbon subsites and nine different bonds that support nine possible nearest neighbor bond pairing amplitudes as illustrated in Fig. A similar procedure again can be applied to pristine graphene but now in enlarged six site unit cell. Unit cell of Cf, f, within which a pairing amplitude is localized within island hexagons and cannot propagate. For these island phases, numerical calculation of the electron pair potential energy g0 shows that g0 is large. This kind of solutions is a consequence of the six atom basis and does not appear for the two atom basis. Also, there are three solutions of the form In fact reduction of symmetry leads to increasing of the system free energy. After the orthogonalization procedure, one obtains three solutions Φ2, two normal bands are E± = ±t1|η0| which fold to six branches in mini-BZ of C6 i.e. γ are of s-wave character, α and β are of the form d ± id type, i.e. d-wave pairing emerges from the d-wave branches of the folded band structure s-wave pairing arises from the s-wave branch M, is responsible for dominant singlet superconductivity in chiral d ± id symmetry. Upon electron doping to the critical vHs at g0 in the d ± id phase decreases, beyond which density of states decreases. g0 increases until a second critical value of doping s-wave pairing occurs. Bloch states in higher conduction bands include combinations of s and f symmetries that favor extended s wave pairing. The multiband character is responsible for stabilizing singlet s superconductivity at high electron or hole doping.For pristine graphene C6 gap solutions with those of folded bands C6 at the same doping. Numerical results for pristine graphene gap equation performed in the nearest neighbor approximation in ref.12 have been extended by applying a more accurate tight binding model fit to the DFT band structure of pristine graphene23. Although a quantum critical point for zero doping reported by Black-Schaffer and Doniach12 at dimensionless coupling d- and s-wave solutions are degenerate. In the more realistic tight binding model we applied, this degeneracy is not observed at the Γ point, and the d-wave solution is dominant. This difference may be consequence of particle-hole symmetry breaking of valence and conduction bands. Also the van Hove singularity at the M point is moved from 0.25 doping for nearest neighbor hopping to 0.16 doping in the accurate model. The phase transition from d-wave to s-wave is shifted to 0.35 doping instead of the 0.4 doping reported for nearest neighbor hopping12. Numerical calculations for this more detailed model are illustrated in Fig. To understand how superconducting phases of graphene can be affected by decoration by Li, one can compare the LiCs band and C π band is small. This means nearest neighbor Li-C hopping is in the range When graphene is decorated by Li, around 0.68 electron per lithium atom transfers to neighboring C sites, viz. s- or d-wave character and their magnitudes are 6 they become mixed and vary with αsy appears in the pairing amplitudes. In terms of this coefficient we have classified superconducting phase symmetries into three groups. Eqs 6. There are three categories of solutions which have not appeared in complete form in the literature. The total of nine phases arise from spatial, and therefore hopping parameter, symmetry breaking.Li decoration of graphene changes not only the band structure but also the Bloch wave coefficients from those of pristine graphene. While pristine graphene Bloch wave coefficients have pure f, αsy = 0 identical to that of folded pristine C6. For the second category, αsy (denoted by α−) is negative, in the case of pristine α− = −2 as discussed. These three phases break the two site cell symmetry, and numerical calculation shows that the pairing potential g0 must be large to realize these phases. For the last category α+ is positive. Three phases which correspond to α+ > 0 include In the first category Φwt and α+, for s-wave phase is significant due to its spatial isotropic symmetry. In spite of the pristine nature this phase no longer preserves two band model symmetry. On the other hand, d-wave phases are hardly affected by doping and their superconductivity is more persistent against perturbation. The chirality or non-chirality of Cooper pairs in these phases is undetermined, however. As shown in Fig. In the limiting case of folded six band pristine graphene Tc and chemical potential μ0, for each of nine possible superconducting phases, Eqs C6 to find the corresponding pairing potential αsy coefficient. Smaller g0 means less Cooper pair formation energy is required. Figure Tc in terms of the pairing potential g0 for LiC6 in which μ0 = 0. For a given transition temperature Tc, by changing the chemical potential μ0 of LiC6 via gating, one can engineer the pairing potential g0. Figure g0-μ0 phase boundary diagram at Tc = 0.1 K. As illustrated in this figure, similarly to pristine graphene, decoration with Li atoms makes it is possible to change the dominant pairing and to have a symmetry-change phase transition from d to “distorted s-wave.” Changing μo up to μo−v ≈ 0.22 eV so that the distance between the Fermi energy and the saddle points decreases, leads to a decrease in g0. Continuously increasing μo up to 0.5 eV causes g0 to increase for both d-wave and “distorted s-wave” pairing, and after that a smooth decrease proceeds. For both symmetries at critical μo−c = 1.3 eV mixed state exist.At a given critical temperature μo−c = 1.3 eV, the flat band plays a primary role in formation of Cooper pairs with lowest energy. The Bloch wave function of this band consists of d and p character, therefore Γ12, Γ15, Γ45 and Γ48 in Eq. d–wave pairing is dominant. Beyond that, the uneven part of the “flat band” and also upper bands assume a major role. These bands consist of d, p, s, and f character Bloch wave functions (as defined in earlier sections) with a significantly low density of states. In this case Γ12, Γ15, Γ45 and Γ48 change their sign, hence s-wave pairing is favored.Up to g0 in the limit of pristine graphene is minimal for all dopings. Our calculations indicate that any perturbation of the flat band reduces Tc. The flat band can be perturbed through electron hopping from decorating atoms to carbon sites (wt. For fixed doping at wt = 0.94 as obtained for lithium decorated, in a variety of Li-C hopping between 0.3–0.4 eV, numerical calculation doesn’t show significant altering of pair interaction potential g0 in s- and d-wave phases. But, as one could see there is not an explicit behavior in a general coupling strength. A result is that a general aspect of superconducting pairing in LiC6 and pristine graphene is almost the same in the dxy phases due to robustness of the flat band against perturbation.Numerically we have demonstrated that electron pairing d-wave phases exist and are dominant symmetry of pairing in both pristine and Li decorated graphene. Pure s-wave phase does not appear in LiC6, and s-wave superconductivity in metal decorated graphene is disfavored because of spatially increased overlap for s-wave symmetry. These results show that while degree of doping plays a major role in the graphene superconductivity, perturbation effects of decorating atoms finally determine the phase diagram. Our work also provides a new type of classification of superconducting phases in LiC6-like nanostructures, and certain aspects of the formalism may be useful in modeling the recently observed superconductivity in magic angle bilayer graphene3.To summarize, our calculations indicate that Supplementary Materials


Fabrication of microsystems is traditionally achieved with photolithography. However, this fabrication technique can be expensive and non-ideal for integration with microfluidic systems. As such, graphene fabrication is explored as an alternative. This graphene fabrication can be achieved with graphite oxide undergoing optical exposure, using optical disc drives, to impose specified patterns and convert to graphene. This work characterises such a graphene fabrication, and provides fabrication, electrical, microfluidic, and scanning electron microscopy (SEM) characterisations. In the fabrication characterisation, a comparison is performed between traditional photolithography fabrication and the new graphene fabrication.  Here, the minimum achievable feature size is identified and graphene fabrication is found to compare favourably with traditional photolithography fabrication. In the electrical characterisation, the resistivity of graphene is measured as a function of fabrication dose in the optical disc drive and saturation effects are noted. In the microfluidic characterisation, the wetting properties of graphene are shown through an investigation of the contact angle of a microdroplet positioned on a surface that is treated with varying fabrication dose. In the SEM characterisation, the observed effects in the previous characterisations are attributed to chemical or physical effects through measurement of SEM energy dispersive X-ray spectra and SEM images, respectively. Overall, graphene fabrication is revealed to be a viable option for development of microsystems and microfluidics. Strategic leveraging of these physical properties has led to microsystems becoming ubiquitous in today’s modern world. Further to this point, microsystems have found applications in the following industries: automotive, e.g., seatbelt activation with accelerometers5 and airbag switches6; biotechnology7, e.g., polymer chain reaction8 and advances in DNA and protein separations8; and healthcare9, e.g., disposable blood pressure sensors10 and point-of-care diagnostic devices11.Microsystems, or microelectromechanical systems (MEMS), make use of distinct and advantageous physical properties, including piezoelectric effects12 and electrowetting-on-a-dielectric devices13.Point-of-care diagnostic devices are of particular interest—given their life saving potential—and such microsystems make use of microfluidic technologies. These microfluidic technologies enable microlitre and nanolitre fluid volumes to be actuated through microchannels, in the case of continuous-flow microfluidics, or between planar electrodes, in the case of digital  microfluidics14. Here, it is desirable to achieve fabrication of highly-conductive electrodes with micro-scale feature sizes15 and to utilise materials with favourable wetting properties 16. The need for low-cost materials that lend themselves to use in microfluidic devices has led researchers to explore the novel material graphene19. Graphene can be fabricated in specified patterns, with graphite oxide undergoing optical exposure and conversion to graphene with simple optical disc drives, as reported previously15. Preliminary studies have shown favourable wetting properties17. However, there is still a gap in the literature for deep explorations into the fabrication of graphene  and characterisation of graphene for use in microsystem and microfluidic electrodes .Microsystems, and its subset of microfluidic technologies, share a desire for low cost and rapid fabrication of constituent components, as traditional photolithography fabrication can have debilitating associated cost. Specifically, there is great interest in developing low cost fabrication processes for microsystems and this is the subject of much recent researchIn this work, the above gap in the literature is addressed. An exhaustive characterisation of graphene fabrication is performed—with consideration to application in microfluidic electrodes. First, a fabrication characterisation is presented. Here, a comparison is performed between traditional photolithography fabrication and the new graphene fabrication.  Here, the minimum achievable feature size is identified and graphene fabrication is found to compare favourably with traditional photolithography fabrication. Second, an electrical characterisation is presented. The resistivity of graphene is measured as a function of fabrication dose in the optical disc drive. The minimum fabrication dose required for saturation of resistivity is identified. Third, a microfluidic characterisation is presented. Here, the wetting properties of graphene are shown through an investigation of the contact angle of a microdroplet positioned on a surface that is treated with varying fabrication dose. Fourth, a scanning electron microscopy (SEM) characterisation is presented. The observed effects in the previous characterisations are attributed to chemical or physical effects through measurement and analysis of SEM energy dispersive X-ray spectra and SEM images, respectively. Ultimately, graphene fabrication is revealed to be a viable option for development of microsystems and microfluidics.20. To establish this minimum achievable feature size, a recipe for graphene fabrication is followed, as described in the Appendix of this manuscript. For this fabrication characterisation, a mask pattern is developed  graphene fabrication method, minimum achievable feature size is considered. Once established, this minimum achievable feature size can be compared to that of traditionally photolithography fabrication, commonly used in fabrication of microsystemsped Fig.  with decn = 5, through the optical disc drive to be ideal for distinguishing electrodes gaps. This conclusion is supported by Fig. It should be noted that the implementation of the recipe for graphene fabrication revealed five doses, g = 500 to 40 μm in Fig. g = 60 μm  than traditional photolithography fabrication, is a contender for fabrication of electrodes in microsystems (or in microfluidic devices). This is because graphene fabrication achieves micro-scale minimum feature sizes. It is shown in the next subsections that graphene fabrication can yield great utility for use in microfluidic devices.ρ, versus the number of doses through the optical disc drive, n. The error bars represent the standard deviation over three trials. The resistivity initially forms a high value (beyond the limit of our measurement) for n = 1, which indicates the insulative nature of the graphite oxide. The resistivity falls considerably for additional doses and reaches a steady-state resistivity, ρ0, after n = 12 doses, representing the final conversion of graphite oxide to graphene.Characterising the electrical properties of microsystems fabricated using the graphene fabrication method is of the utmost importance for use as electrodes in microsystems. Figure Given the ability to achieve (and even tune) the resistivity, graphene fabrication of electrodes is again shown to be a viable option for microsystems.16. To meet this desire, a detailed investigation of the wetting properties of the graphene fabrication method is performed through a microfluidic characterisation. The results of the microfluidic characterisation is presented in Fig. θ, versus the number of doses through the optical disc drive. For the microdroplets placed on the graphene surface, images of representative low, medium, and high contact angles are shown as insets for dose of n = 2, 4, and 9, respectively.  The contact angle increases monotonically from θ = 50° at a dose of n = 1 until it saturates to a contact angle of θ = 116° in the range of n = 7–9. This result is of particular importance for optofluidic applications16 within microfluidic devices, as graphene fabrication provides tunability of the contact angle over the contact angle range of 50° to 116°, without the application of any further hydrophobic materials , which is favourable compared to traditional photolithography fabrication.When fabricating electrodes for use in microsystems, particularly in microfluidic devices, it is desirable to select materials with favourable wetting properties such as a high contact angle An important clarification for the above characterisations is to distinguish what observations can be attributed to chemical changes  and what observations can be attributed to physical changes . This important clarification can be addressed through a scanning electron microscopy (SEM) characterisation, whereby chemical changes are observed through SEM energy dispersive X-ray (EDX) spectra and physical changes are observed through SEM images.n = 0 through to n = 13, and each sample undergoes spectroscopy to produce SEM EDX spectra.) Representative SEM EDX spectra are shown for n = 0, 1, and 13 as these samples correspond to, respectively, graphite oxide prior to optical exposure, graphene after initial optical exposure, and graphene after extensive optical exposure. There is a transition from n = 0  to n = 1  that results in large change in atomic percent of carbon and oxygen of respective atomic percentages of 65% and 35% (the mean of four points on the sample for n = 0) to atomic percent of carbon and oxygen of respective atomic percentages of 76% and 24% which is maintained through to n = 13  within an atomic percent standard deviation of 2%. . The overall trend follows previous observations of an increase and decrease in atomic percent of carbon and oxygen, respectively21. The large chemical change due to initial optical exposure mimics the large change in resistivity observed in Fig. To attribute graphene effects to chemical changes, SEM EDX spectra are presented in Fig. 22. The surface roughness, r, then can be found as the surface area divided by the top view area of interest. This surface roughness can be divided into two parts through r = 1 + Δr, where the unit value is the surface roughness of a smooth surface and the relative roughness, Δr, quantifies the increased roughness beyond the smooth surface. In Fig. r/Δr0, is shown versus dose, n.  It can be seen that there is an overall increasing trend for n = 0–6. The normalized relative roughness begins to stabilize in the range of n = 7–9 and this mimics the saturation of contact angle observed in Fig. 23.To attribute graphene effects to physical changes, an SEM image analysis is presented in Fig. A graphene fabrication process was characterised for use in microsystems and microfluidic devices. The graphene fabrication process was based on graphite oxide undergoing optical exposure through an optical disc drive and converting to graphene. Fabrication, electrical, microfluidic, and scanning electron microscopy characterisations of the graphene fabrication yielded micro-scale feature sizes, low resistivity properties, tunable wetting properties, and provided insight into chemical and physical graphene effects, respectively. The resistivity of graphene showed a close connection to chemical composition and the contact angle of microdroplets on graphene showed a close connection to physical composition of the surface. Overall the graphene fabrication process compared favorably to a photolithography fabrication process.The graphene fabrication procedure begins with a graphite oxide solution, made by mixing 100 mg of graphite oxide powder  with 27 mL of deionized water in a beaker placed into a bath sonicator (Fisherbrand M series) for 15 minutes. To prepare for exposure, a film of transparent plastic material in the shape of an optical disc is adhered onto a LightScribe disc. A volume of 17 mL of the graphite oxide solution is applied to the LightScribe disc with a syringe and left to air dry for 48 hrs. The mask pattern Fig.  designedFor the traditional photolithography fabrication, the desired mask pattern is made using image processing software with feature sizes as small as 5 µm. To make a negative photomask, the image file is exported from the image processing software as a portable document format file with a rendering resolution of 10,000 dots per inch and printed with black emulsion onto a film of transparent plastic material.To prepare the substrate, a 75 mm × 25 mm × 1 mm glass slide with a 50 nm copper coating is cleaned of any debris. The copper-coated glass slide then undergoes a spincoating process to coat with negative photoresist. The copper-coated glass slide is placed into a programmable spincoater (CEE 200X), and approximately 3 mL of negative photoresist (MICROCHEM SU-8) is dispensed manually onto the surface using a syringe. The programmable spin coater is set to leave a 30 µm layer of the negative photoresist through a process of 500 rpm for 25 seconds with acceleration of 100 rpm/s and 1300 rpm for 30 seconds at an acceleration of 300 rpm/s. The copper-coated glass slide then undergoes a soft bake process. The copper-coated glass slide is placed onto a hotplate (Corning PC351) for 90 seconds, and then immediately into an oven (Fisher Isotemp Vacuum Oven 282) for 5.5 min and then left to cool for 15 min. Once the copper-coated glass slide is cool, the negative photomask is placed ink-side-down onto the surface of the slide coated with SU-8 and loaded into an ultraviolet exposure chamber (KLOÉ UV KUB2) for an exposure of 100% energy for 7 seconds. The length of exposure time is found through


Acute exacerbation of idiopathic pulmonary fibrosis (AE-IPF) is an acute respiratory worsening of unidentifiable cause that sometimes develops during the clinical course of IPF. Although the incidence of AE-IPF is not high, prognosis is poor. The pathogenesis of AE-IPF is not well understood; however, evidence suggests that coagulation abnormalities and inflammation are involved. Thrombomodulin is a transmembranous glycoprotein found on the cell surface of vascular endothelial cells. Thrombomodulin combines with thrombin, regulates coagulation/fibrinolysis balance, and has a pivotal role in suppressing excess inflammation through its inhibition of high-mobility group box 1 protein and the complement system. Thus, thrombomodulin might be effective in the treatment of AE-IPF, and we and other groups found that recombinant human soluble thrombomodulin improved survival in patients with AE-IPF. This review summarizes the existing evidence and considers the therapeutic role of thrombomodulin in AE-IPF.  Idiopathic pulmonary fibrosis (IPF) is a representative subtype of idiopathic interstitial pneumonia. IPF is characterized histologically by a usual interstitial pneumonia (UIP) pattern and radiologically by the presence of typical honeycombing in high-resolution computed tomography images . The cliAlthough the pathophysiology of AE-IPF is poorly understood, acceleration of underlying fibroproliferative disorders might be involved ,10. Hist2) and fraction of inspired oxygen (FiO2) ratio (PaO2/FiO2 ratio) is one of the indexes used to classify the degree of respiratory dysfunction. In retrospective studies, PMX-DHP improved PaO2/FiO2 ratio [The incidence of AE-IPF is relatively low, and the number of clinical trials has thus been small. Therefore, evidence regarding the treatment of AE-IPF is limited. No randomized clinical trials have investigated AE-IPF; thus, international consensus recommendations are based on weak evidence . In geneO2 ratio  and 1-yeO2 ratio . DonahoeO2 ratio . Plasma O2 ratio . In the Optimal treatment for AE-IPF has not been defined, and AE-IPF mortality remains high. Thus, new treatment strategies are urgently needed. Thrombomodulin is a transmembranous glycoprotein expressed on the surface of vascular endothelium. It interacts with multiple proteins to block blood coagulation and inhibit inflammation . RecombiThe pathological findings of AE-IPF comprise DAD superimposed on the underlying fibrosis of IPF . Oda et Several lines of evidence indicate that anticoagulant therapy improves fibrosis and inflammation in lungs in animal models. Inhalation of anticoagulants, including heparin and PAI-1 inhibitor, suppressed development of bleomycin (BLM)-induced pulmonary fibrosis in mice ,34. ThesPathological analysis of transplanted lungs revealed that lungs of IPF patients who developed AE-IPF had very high levels of inflammatory cells, including neutrophils and lymphocytes . SeveralIncreased proinflammatory cytokines such as IL-6, and IL-8 were observed compared to stable IPF in the analysis of serum cytokine profile of AE-IPF ,41,42. MThe profibrotic roles of alternative activated macrophages were reported in animal models of lung fibrosis and IPF ,45,46,47Activated inflammatory macrophages produce proinflammatory cytokines and chemokines such as IL-8 that attract neutrophil recruitment and activation in the lungs of patients with AE-IPF. Lymphocytic activation might also promote inflammation. Moreover, sustained activation of profibrotic macrophages might be important in the development of irreversible fibrosis during the fibrotic phase of lungs with DAD.High-mobility group box-1 (HMGB-1) is a non-histone nuclear protein that promotes binding of transcription factors and regulates maintenance of nucleosomal structure. HMGB-1 is constitutively present in nearly all cell types and is passively released when cells are injured . ActivatPrevious studies reported that HMGB-1 mediates coagulation activity. HMGB-1 binds to RAGE on platelets, which leads to platelet aggregation . HMGB-1/HMGB-1 promotes tissue repair and regeneration. Through RAGE, it induces proliferation of smooth muscle cells in rats and humans ,68. FurtA study of serum HMGB-1 levels in IPF patients found no significant difference between stable IPF patients and control subjects ; howeverHamada et al. reported that HMGB-1 was highly expressed in alveolar epithelial cells and inflammatory cells in a murine bleomycin model . AnotherHMGB-1 might have a crucial role in AE-IPF pathogenesis through its proinflammatory, thrombotic, and profibrotic function and may therefore be a biomarker of AE-IPF and a potential target for AE-IPF treatment.Thrombomodulin is a transmembranous glycoprotein found on the cell surface of vascular endothelial cells encoded by chromosome 20 in humans. It comprises 557 amino acids organized into 5 domains, including the N-terminal lectin-like domain, 6 epidermal growth factor (EGF)-like domain, a serine and threonine-rich domain, a transmembrane domain, and a cytoplasmic domain . ThromboThrombomodulin can form a complex with thrombin, a terminal enzyme of the blood clotting process. The region including the fourth, fifth, and sixth EGF-like domains of thrombomodulin are responsible for thrombin binding . Thus, tAnother function of the thrombomodulin/thrombin complex is activation of protein C. Activated protein C inhibits coagulation by cleaving circulating protein C and inactivating procoagulant factors such as Va and VIIIa ,82. ActiThe N-terminal lectin-like domain of thrombomodulin has potent anti-inflammatory activity . MaruyamThrombomodulin might also exert an anti-inflammatory effect by regulating the complement system . The N-tIn an analysis of thrombomodulin in IPF patients, serum thrombomodulin levels were significantly higher in AE-IPF patients than in stable IPF patients . MoreoveSeveral recent lines of evidence show the therapeutic effect of thrombomodulin in animal models of lung inflammation and fibrosis. In a mouse model of acute lung injury, mice were intraperitoneally administered rhTM. rhTM treatment prolonged survival and improved lung histological scores. Levels of regulatory T cells and anti-inflammatory cytokines such as IL-10 were higher in the lungs of thrombomodulin-treated mice .Fujiwara et al. reported that rhTM improved lung inflammation and fibrosis by means of an anti-inflammatory and anticoagulant effect in several animal models in vivo. A murine BLM lung injury model and TGF-β1 transgenic mice treated with LPS showed less fibrotic response after treatment with rhTM. rhTM inhibited apoptosis of alveolar epithelial cells induced by BLM stimulation in vitro and reduced apoptotic cells in TGF-β1 transgenic mice. The study authors hypothesized that rhTM suppresses lung inflammation and fibrosis by suppressing coagulation abnormality and inflammation, and by its antiapoptotic action .Kida et al. investigated the role of HMGB-1 and thrombin in the fibrotic process and the inhibiting effect of rhTM against this process in vitro and in a BLM mouse model . HMGB-1 Several recent studies in Japan investigated the effectiveness of rhTM for AE-IPF . Patient2/FiO2 ratio and survival rate after AE-IPF onset were significantly better in the rhTM-treated group than in historical controls [(30%) . In addicontrols .p = 0.048). A univariate Cox proportional hazards regression model showed that rhTM treatment was a predictive factor for survival. Regarding adverse events, 1 patient in the rhTM group developed mild bleeding [We previously compared outcomes of 16 patients treated with rhTM and 25 patients treated with conventional therapy. The former received rhTM (0.06 mg/kg/day) for 6 days as initial treatment, in combination with corticosteroids. As compared with the control group, the rhTM group had a significantly higher survival rate at 3 months (40% vs 69%, bleeding . In an ebleeding .Serum HMGB-1 level might reflect AE-IPF disease activity and predict survival. Therefore, change in HMGB-1 level was analyzed in AE-IPF patients treated with rhTM. Hayakawa et al. reported that serum HMGB-1 level did not significantly change from day 0 to day 29 after AE-IPF in 7 patients . We examAdverse events related to anti-coagulant effect of rhTM were reported in few patients treated in the previous studies. Mild hemoptysis and hematuria were observed in one patient in our study populations ,101 and The results of these clinical studies suggest that rhTM treatment is beneficial in AE-IPF. However, the sample sizes of the studies were too small to conclusively show the effectiveness of rhTM in AE-IPF. In addition, most were single-center retrospective studies. To confirm the effectiveness of rhTM treatment, a multicenter prospective randomized controlled trial is ongoing in Japan.AE-IPF is an acute respiratory dysfunction characterized by alveolar epithelial cell and endothelial cell injury, inflammation, coagulation abnormality, and fibrotic deposition. rhTM suppresses inflammation mainly by binding to HMGB-1 and it also suppresses excess coagulation by forming a complex with thrombin and by activating protein C. Several recent clinical studies suggested the possibility that rhTM improve the prognosis of AE-IPF when it used in combination with corticosteroids.We reviewed the pathogenesis of AE-IPF, the therapeutic roles of thrombomodulin in AE-IPF, and evidence from clinical trials. Thrombomodulin is a promising treatment for AE-IPF because of its multiple anti-inflammatory, anticoagulant, antifibrotic effects. A multicenter prospective study to confirm the effectiveness of rhTM treatment is ongoing.


Herein, the nanohybrids of lanthanum (La) and selenium (Se) co-doped bismuth ferrites with graphene oxide were synthesized. A structural analysis from X-ray diffraction confirmed the transition of phases from rhombohedral to the distorted orthorhombic. Scanning electron microscopy (SEM) revealed that the graphene nano-sheets homogenously covered La–Se co-doped bismuth ferrites nanoparticles, particularly the (Bi0.92La0.08Fe0.50Se0.50O3–graphene oxide) LBFSe50-G sample. Moreover, the band-gap nanohybrids of La–Se co-doped bismuth ferrites were estimated from diffuse reflectance spectra (DRS), which showed a variation from 1.84 to 2.09 eV, because the lowering of the band-gap can enhance photocatalytic degradation efficiency. Additionally, the photo-degradation efficiencies increased after the incorporation of graphene nano-sheets onto the La–Se co-doped bismuth ferrite. The maximum degradation efficiency of the LBFSe50-G sample was up to 80%, which may have been due to reduced band-gap and availability of enhanced surface area for incoming photons at the surface of the photocatalyst. Furthermore, photoluminescence spectra confirmed that the graphene oxide provided more electron-capturing sites, which decreased the recombination time of the photo-generated charge carriers. Thus, we can propose that the use of nanohybrids of La–Se co-doped bismuth ferrite with graphene oxide nano-sheets is a promising approach for both water-treatment and water-splitting, with better efficiencies of BiFeO3.Photocatalysts are widely used for the elimination of organic contaminants from waste-water and H They also found that the Eu3+-doped (3%) samples showed an enhanced photocatalytic activity for RhB that degraded 97% within 50 min under UV-light . T. T2O3  showed the absence of spin orbit splitting [2 [The Fe-2p20.60 eV . The corplitting . The low3+ and Se+4 co-doped BFO with graphene oxide samples were examined, as presented in The UV-vis diffuse reflectance spectra of the pristine BFO and nanohybrids of Laex = 300 nm) spectrum was accomplished to observe the recombination rate of the photo-generated charge carriers of the as-synthesized nanohybrids of La3+ and Se+4 co-doped BFO with graphene oxide. The photocatalytic performance of semiconductor photocatalysts mainly depends on the recombination time of photo-generated charge carriers [3+ and Se+4 co-doped BFO with graphene oxide nanohybrids series. Asymmetric broad spectra in the range of 380–700 nm with prominent peaks identified in the positions of 425 and 455 nm were observed. Usually the band-gap of pristine BFO lies in green visible region; the emission peak at 455 nm corresponds to the transition of electrons from the valence band to conduction band [Photoluminescence emission  under visible-light obtained from a Xenon lamp (power = 300 W and filter wavelength = 420 nm) [2. The origin of the enhanced photocatalytic performance in BFO is its low recombination rate and enlarged surface area. As the surface area of the photocatalyst increases, the dye gains direct contact with the catalyst, making redox reactions easy to occur. By increasing Se-doping, the average distance between trap sites decreases, causing a large recombination rate [The photocatalytic performance of the nanohybrids of La 420 nm) ,67. The ion rate ,69,70. H3+ and Se+4 co-doped BFO nanostructures, as reported in the literature [By using the LBF–G nanohybrid as a photocatalyst, over 72% of CR was removed within the same length of irradiation time (90 min). For the LBFSe25–G nanohybrid sample, almost 50% of CR was degraded. Moreover, with varying the concentrations of selenium onto LBF–G (up to 50%), the degradation efficiency reached 79%. However, the LBFSe100–G nanohybrid sample showed 68% degradation efficiency for CR under visible-light irradiation, which suggests that optimal concentrations of co-dopant and graphene oxide significantly enhanced the photocatalytic activity of BFO. The data suggest that the presence of graphene oxide significantly affects the photocatalytic behavior of Laterature ,72.2O and CO2 as by-products after reacting with the hazardous organic compound. The mechanism started with the absorption of photons, then led to the creation of electrons and hole pairs, and, finally, ended with the production of superoxide radicals for the degradation of organic pollutant. The highest photocatalytic degradation efficiency of LBFSe50–G might have been due to the availability of extra electron-capturing centers due to the availability of the larger surface areas of graphene nano-sheets. These captured electrons converted O2 into O2•, where the H2O molecules were converted into OH•. Both radical species efficiently degraded the organic pollutant and converted it into a harmless product (CO2 and H2O). The stability of a photocatalyst is also important parameter for its industrial applications. As such, the chemical stability of the LBFSe50–G sample was observed its after repeating the photocatalytic experiment three times, as shown in These calculations show that the reduction of graphene oxide to rGO was very low. The low reduction meant the inefficient charge transfer and transportation of photo-generated charges. This inefficient charge transfer and charge trap increased the rate of recombination of charges and caused a decrease in photocatalytic activity. In these samples, the overall degradation was due to adsorption. The band-gap was decreased slightly after adding selenium onto La-doped BFO; however, after incorporation, the band-gap varied significantly from 1.84 to 2.09 eV. The band-gap value for the LBFSe50–G sample was 2.01 eV. The possible photocatalytic mechanism could be explained as: When the visible-light fell on the surface of the photocatalyst, the electron and holes were created. The electron-hole pair further produced H3+ and Se+4 co-doped BFO with graphene oxide were synthesized by the co-precipitation method. A structural analysis from X-ray diffraction confirmed the transition from the rhombohedral phase to the distorted orthorhombic phase. Moreover, the presence of Se-3d, Bi-4f, C-1s, O-1s, Fe-2p and La-3d constituents was also confirmed from the XPS examination. SEM micrographs revealed that the graphene oxide nano-sheets thoroughly covered the nanoparticles, which provided a broad surface area for the incoming photons of irradiated visible-light. The as -synthesized nanohybrids showed enhanced photocatalytic activity for CR under visible-light, where the LBFSe50–G sample exhibited the highest photocatalytic performance—up to 80%. This may have been due to the decrease of the recombination time for photo-generated charge carriers due to the availability of trapping sites provided by graphene oxide nano-sheets. Therefore, the results suggest that the nanohybrids of La3+ and Se+4 co-doped with graphene oxide nano-sheets can be promising photocatalysts for the removal of waste-water from industrial waste.The nanohybrids of La


However, the ultimate stretchability of graphene has not yet been demonstrated due to the difficulties in experimental design. Here, directly performing in situ tensile tests in a scanning electron microscope after developing a protocol for sample transfer, shaping and straining, we report the elastic properties and stretchability of free-standing single-crystalline monolayer graphene grown by chemical vapor deposition. The measured Young’s modulus is close to 1 TPa, aligning well with the theoretical value, while the representative engineering tensile strength reaches ~50-60 GPa with sample-wide elastic strain up to ~6%. Our findings demonstrate that single-crystalline monolayer graphene can indeed display near ideal mechanical performance, even in a large area with edge defects, as well as resilience and mechanical robustness that allows for flexible electronics and mechatronics applications.The sp The extraordinary mechanical properties of graphene are usually measured on very small or supported samples. Here, the authors develop a method to test a large area of graphene and show that even with edge defects it displays near-ideal mechanical performance. New methods were later developed to fabricate large-scale graphene monolayers, among them the most commonly used method is chemical vapor deposition (CVD), which can fabricate monolayer graphene towards mass production8. The mechanical robustness of them, however, is the basis for successful realization of their structural and functional performances. Theoretical investigations10 show that the Young’s modulus and ideal tensile strength of graphene are about 1 ± 0.1 TPa and 100–130 GPa, with maximum strains to failure up to ~13–19% and ~20–26% (with nonlinear elastic behavior at large strain12) along the armchair and zigzag directions, respectively. With the presence of defects  in large-area CVD-grown graphene, it is expected that the corresponding tensile strength shall be considerably reduced17. These understandings, however, remain to be assessed by direct experimental evidences.Since graphene was studied by Geim and Novoselov through mechanical exfoliation from the highly oriented pyrolytic graphite in 200420. These pre-cracked samples fracture in a brittle manner, featuring breaking stress significantly lower than the ideal strength of graphene. Still, direct measurement of the realistic strength of crack-free monolayer graphene or below the flaw-tolerance limit21 has not yet been made, which are more relevant for 2D material device applications based on high-quality single-crystalline graphene where cracks with lengths conforming to the applicability of linear elastic fracture mechanics do not show up (with only defects created during the growth and sample fabrication processes)23. Alternatively, indirect tests through nanoindentation into suspended graphene were earlier performed. Lee et al. reported that the intrinsic Young’s modulus of monolayer graphene can be indeed up to 1 TPa by the atomic forces microscope-based nanoindentation, and the intrinsic strength is up to 130 GPa, suggesting that graphene is the strongest material5. Consequently, the localized breaking stress measured by indenting graphene with grain boundaries is reduced by 15% from that of the pristine lattice24. However, these local tests measured only a small area of graphene under the indenter tip, and the strain distribution in the whole graphene membrane was highly non-uniform. This very sensitivity to the local atomic structures and prominent geometrical effects in indentation tests25 limit the technique in providing a reliable assessment of mechanical performance of large-area graphene with engineering relevance in practical applications, such as reinforced composites and electromechanical devices. Burst tests were recently carried out to directly measure the strength of suspended graphene under gas flow measurement, where a wide distribution of burst pressures was reported and attributed to wrinkles and defects26. As the engineering strength and strain to failure of a brittle material are controlled by the weakest point, a direct tensile test of large-area crack-free graphene membrane under uniaxial straining condition remains as the most efficient way for a quantitative study.Fracture properties of free-standing graphene with pre-crack, in forms of bilayer and multilayers, were experimentally measured in scanning electron microscope (SEM) and transmission electron microscope (TEM)31. Recently, twisted bilayer graphene around the first magic angle (θ ≈ 1.1°) has been reported for the realization of unconventional superconductivity and correlated insulator behavior33. Alternatively, by varying the interlayer strain with hydrostatic pressure, superconductivity in twisted bilayer graphene can be also achieved34. However, significant strain-engineering effects often require large elastic strain to be applied, especially for monolayer graphene with zero intrinsic bandgap35. In contrast, in the experimental studies reported in literature, the stretchability of finite graphene sheets was limited by the presence of point and line defects, resulting in a common tensile strain to failure of ~1%36, far below the strain level associated with notable strain effects. These facts clearly indicate that a sample-wide elastic strain level much higher than 1% is desired to endow graphene with realistic strain-tunable device applications27.Besides, mechanical straining of graphene has been investigated as it can actually induce lattice distortion and phonons, modify the chemical reactivity and magnetic characteristics, and modulate their electronic and optoelectronic properties37. We demonstrate their representative engineering tensile strength reaching ~50–60 GPa with sample-wide elastic strain up to ~6%, even in a large area with edge defects, displaying near-ideal mechanical performance, high resilience, and mechanical robustness.In this work, we circumvent the gripping problem by developing well-controlled sample transfer and shaping techniques, namely well control of the shape of free-standing CVD-grown graphene samples, which were transferred onto the tensile straining stage through a modified wet-transfer method, and explore the in-plane mechanical responses of suspended monolayer graphene by our previously developed in situ nanomechanical testing platform under SEM5 to the substrate as illustrated in Fig. −1) and G (~1579 cm−1) peaks is ~3, and the D peak (~1345 cm−1) is negligible, indicating the transferred sample is monolayer7 of high quality. The edges of the suspended graphene were characterized separately inside a high-resolution transmission electron microscope operating at 80 kV. As can be seen from Fig. To conduct in situ tensile tests, monolayer graphene samples were chosen and transferred onto a push-to-pull (PTP) micromechanical device through a polymethyl-methacrylate (PMMA)-assisted method as described in the Methods part. The graphene samples loaded on PTP devices were then cut into ribbon shape by focused ion beam (FIB) processing, with controlled size/area of the suspended part, as shown in Fig. 31.To study the elastic properties and fracture behavior of monolayer graphene, displacement-controlled tensile tests were performed on the suspended graphene monolayers inside SEM, such as a sample shown in Fig. E2D =~309 N/m (see Methods), which is comparable with the 2D Young’s modulus of 348 N/m predicted from ab initio calculations12. The stress/strain state in the sample is not entirely uniform in our measurements, differing from the center free-standing part with zero transverse stress to the both clamping ends with transverse strain restricted (see Methods). However, to extract elastic constants from the force-displacement data measured in the experiments, assumptions have to be made. Our continuum mechanics simulations show that the uniaxial stress assumption is more reasonable than uniaxial strain that yields E2D = 309 N/m, which applies for most area of the sample  Young’s modulus of the free-standing graphene sample can be calculated as E3D = ~920 GPa in the uniaxial stress assumption as well (see Methods), considering the actual width and gauge length of the tested sample. The measured Young’s modulus, along with some other sample results ranging from ~900–1000 GPa, is fairly close to the theoretical value of a pristine monolayer graphene crystal4, indicating the high quality of CVD-grown single-crystalline samples and that the presence of defects in the bulk region is negligible in graphene membrane tensile stiffness and modulus, even in large area.We further examined graphene sample by in situ tensile test until fracture. As shown in Fig. 18, our result clearly demonstrates much higher elastic stretchability and tensile strength for free-standing graphene without crack, suggesting the good usability of high-quality CVD-grown single-crystalline graphene. Nevertheless, the representative tensile fracture strength of ~50–60 GPa measured from our multiple samples is considerably reduced compared with the ideal strength of monolayer graphene (~100–130 GPa)7, which should be partly attributed to the edge defects implanted during the sample cutting process by FIB. To quantify this effect, we carried out atomistic simulations to measure the fracture strength of graphene ribbons with armchair, zigzag, and chiral edges under end-clamped tension  and the measured engineering tensile strengths of graphene ribbons with edge defects (σm) is close to our experimental measurements, as shown in the shadowed area in Fig. 5, exceeding one-half of the ideal strength of graphene , which implies that the defects created by FIB cutting are localized at the edge region, and the contribution of edge defects (rather than crack) toward the graphene membrane strength and toughness is less significant. It should be noted that the presence of large amount of internal and edge defects across the whole sample area would obviously modulate the crack behavior and fracture toughness, such as leading to multiple crack stages17, whereas in our high-quality single-crystalline monolayer samples with mostly edge defects, the effect appears less considerable. At last, we calculated the average atomistic principal stress in the region away from defect is marked as σa, and the peak principal stress identified near the defect is marked as σp. The simulation results show that although the value of σp characterize the local stress state of graphene, the ratio σp/σa displays a much wider span, indicating that the maximum principal tensile stress may not be a reliable indicator to determine the nucleation of fracture in single-crystalline graphene. This could be attributed to the significant anisotropy in fracture toughness of the hexagonal lattice that should be included in defining the criterion for failure40. Besides, in view of the likelihood of the presence of small in-plane misalignment of the tensile loaded sample between the two clamping ends, atomistic simulations are carried out to study the fracture strain and strength predicted for a free-standing graphene sample loaded with a non-uniform displacement field  of monolayer single-crystalline graphene was explored by the atomistic simulations coupled with TEM analysis on the fractured samples. The simulation results summarized in Supplementary Fig. 42 in practical engineering applications such as flexible electronics and ultra-strength composites7. Brittle fracture of the free-standing graphene initiated from the damaged lattice sites at edges near the clamping ends, whereas the ultimate stretchability and strength could be further optimized by controlling the edge states. This study reaffirms the high mechanical resilience of graphene, beyond the limit of local probe, and the measured sample-wide strain up to ~6% without inelastic relaxation brings the widely proposed strain engineering into practice27. At last, the in situ nanomechanical straining strategy we developed here for monolayer graphene, allowing for uniform, reversible and well-controllable strain modulation, provides opportunities for dynamically strain-tuned electronics and optoelectronics device applications of other 2D materials and their superlattices/heterostructures through elastic strain engineering43.In summary, by developing robust sample transfer, shaping and clamping techniques, in situ quantitative tensile test of free-standing single-crystalline graphene monolayers has been successfully achieved, showing near-ideal tensile stiffness and elastic modulus (~900–1000 GPa) with representative engineering tensile strength as large as ~50–60 GPa, despite that the reduction in strength arises mainly from edge defects. These results demonstrate the potential of large-area CVD-grown single-crystalline graphene44 by a 514-nm laser, the laser power was set as 1.7 mW to prevent thermal heating and damage of graphene. TEM characterization of the graphene samples before and after testing was carried out using a JEOL 2100 F TEM.CVD-synthesized monolayer graphene films on copper foil (from Sigma Aldrich) with PMMA coating (thickness < 100 nm) were transferred to the PTP device through an optimized chemical etching method. First, to release the graphene from the copper foil, the graphene with copper foil was immersed into Cu etchant (from Sigma Aldrich). Second, the graphene with PMMA was fished up with a piece of polyethylene terephthalate and released into distilled water for three times to remove Cu etchant thoroughly. Third, the graphene with PMMA coating was fished up with the PTP device and then left overnight in a dry cabinet so that the graphene attached firmly with the PTP device. Fourth, the PMMA was dissolved by acetone solution in a critical point dryer, which was used to reduce the surface tension of acetone during dissolving the PMMA layer and protect the suspended graphene on the gap of PTP device. Finally, the suspended graphene was cut into ribbon shape . Tensile strain (ε) can be well controlled and measured directly from the SEM image sequences/videos. The tensile stiffness of the sample can be calculated by subtracting the inherent stiffness of the PTP device from indenter load-displacement curves, and the 2D and 3D Young’s modulus can be calculated as k, t, l, and W are the tensile stiffness, the thickness (taken as 0.335 nm for monolayer graphene), the gauge length and the width of the sample, respectively20. It should be note that to interpret the measured stiffness of the whole sample in terms of the Young’s modulus E as a material parameter, assumptions have to be made for the stress/strain state in the sample. In our experimental setup, the end-clamped loading condition enforces a uniaxial strain state at the ends. However, the center suspended part is free of transverse load. As a result, for samples with a sufficiently large aspect ratio, the uniaxial stress assumption  applied on the free-standing graphene sample is then calculated through conventional E2D into E3D.Graphene on the PTP device was tested using a Hysitron pico-indenter (PI85) inside a FEI Quanta 450 SEM46. The adaptive intermolecular reactive empirical bond-order (AIREBO) potential function is used to describe the inter-atomic interaction in graphene47. For the parameters in AIREBO potential functions, the cutoff distance was adjusted to be 0.2 nm to avoid the spurious strengthening effect after structural failure48. The simulation started at 0 K and the temperature rise during tensile tests is on the order of 10 K. The effect of thermal fluctuation on mechanical responses is thus negligible. Uniform tensile strain was applied by deforming the simulation box and applying atomic displacements at a loading rate of 1 ns−1 before the atoms were allowed to move following their equations of motion. The atomic displacement was determined by assuming an affine deformation of the lattice, which is reasonable owing to the single-crystalline nature of our experimental samples. MD simulations at lower loading rates were also performed and compared with to verify that the deformation and fracture behavior are not affected by these dynamical effects. Tensile tests to study crack nucleation and propagation in clamped samples were carried out with two rows of aromatic rings at the clamping ends fixed or displaced (at a constant rate of 1 ns−1) in the simulations  along the loading direction, and two free boundary conditions in the width and out-of-plane dimensions. The length and width of ribbon are 20 nm and 10 nm, which were validated to exclude size effects on the conclusions we made. Defects were created symmetrically at the ribbon edges, as illustrated in Fig. Supplementary InformationDescription of Additional Supplementary FilesSupplementary Movie 1Supplementary Movie 2Supplementary Movie 3


In this study, a novel method based on the transfer of graphene patterns from a rigid or flexible substrate onto a polymeric film surface via solvent casting was developed. The method involves the creation of predetermined graphene patterns on the substrate, casting a polymer solution, and directly transferring the graphene patterns from the substrate to the surface of the target polymer film via a peeling-off method. The feature sizes of the graphene patterns on the final film can vary from a few micrometers (as low as 5 µm) to few millimeters range. This process, applied at room temperature, eliminates the need for harsh post-processing techniques and enables creation of conductive graphene circuits (sheet resistance: ~0.2 kΩ/sq) with high stability (stable after 100 bending and 24 h washing cycles) on various polymeric flexible substrates. Moreover, this approach allows precise control of the substrate properties such as composition, biodegradability, 3D microstructure, pore size, porosity and mechanical properties using different film formation techniques. This approach can also be used to fabricate flexible biointerfaces to control stem cell behavior, such as differentiation and alignment. Overall, this promising approach provides a facile and low-cost method for the fabrication of flexible and stretchable electronic circuits. Graphene has recently received a lot of interest as a promising conductive material for flexible electronics device fabrication due to its exceptional electrical, optical and mechanical properties4. Graphene has been used in combination with various flexible polymer substrates , polyimide (PI), and polydimethylsiloxane (PDMS))10 as well as other non-conventional substrates 14 through different fabrication processes for the development of next-generation flexible electronics21. The flexible nature of these substrates enables protection of graphene electrodes from various deformations, and facilitates their installation onto irregular, curvilinear surfaces22.The field of flexible and stretchable electronics has garnered increasing interest due to versatility for unique on-body applications including but not limited to portable energy-harvesting devices, electronic skin, wearable electronic devices, and sensors27. Polymer-based transfer methods, that serve as an intermediate step between the metal substrate and the final target substrate for the graphene, involve mostly poly(methyl methacrylate) (PMMA) and polydimethylsiloxane (PDMS) along with others such as PET, PI and polyvinyl alcohol (PVA), due to their surface energy and adhesion forces between polymer/graphene and polymer/target substrate interfaces28. A major challenge for polymer-assisted graphene transfer is protecting the graphene integrity and preventing the mechanical deformation and destruction during transfer. Moreover, polymer-assisted graphene transfer requires substantial processing, such as stamping, plasma etching, chemical etching, washing, and high temperature baking28. Despite the optimization and enhancements in transfer processes28, there are still concerns regarding the remaining residues after the transfer, deteriorating the electronic properties of graphene23. Other than polymers, thermal release tape (TRT)-based graphene transfer has also been used for the roll-to-roll (R2R) technique; however, this method requires applied pressure, etching and high temperatures (~100 °C) to separate the tape from the graphene28.Conventional methods such as chemical vapor deposition (CVD) can fabricate low-cost and large-scale graphene films on metal substrates at high growth temperatures (300–1000 °C), and the graphene is subsequently transferred to the substrate of interest30. However, these techniques typically require additional postprocessing, particularly high temperature sintering, which limits the substrate selection to thermally sensitive substrates, such as poly (PEN), PI or PET possessing relatively low glass transition temperatures30. Graphene transfer after printing may be required for some cases, where mostly polymer-assisted transfer processes described earlier are used29.Graphene-based flexible electronics can also be fabricated through various printing techniques such as inkjet printing, gravure printing, screen printing, and offset printing, providing high throughput large-scale production with low raw material consumption and reduced cost32, transfer printing33, or micro transfer molding34 are other potential fabrication methods. Graphene transfer via patterned sticky/adhesive tape peeling has difficulties in controlling the graphene feature size and resolution32. Although recent studies demonstrated certain improvements in patterned graphene transfer via “Stick-and-Transfer” process, this method still requires use of high graphene amounts to coat PDMS negative pattern features, additional tape peeling for surface cleaning, and is only valid for adhesive tapes, limiting the use of polymer based substrates22. Transfer printing provides high-resolution patterning of graphene through transferring of graphene patterns from a patterned template or mold to a receiving substrate36. However, it requires etching, PDMS stamping and removal39, or DMSO-based surface energy modification of PDMS molds40. In micro transfer molding, the patterned surface of a stamp is placed on the target substrate and the channels are filled with conductive graphene solution through capillary action. This process then requires vacuum drying and removing the stamp from the final substrate41.Transfer of graphene via sticky/adhesive tape peeling44. This situation not only limits the target substrate material selection but also the potential application areas, especially in the healthcare and biomedical fields45. Because of this, most polymer-based transfer methods are an intermediate step between the donor and receiver substrates and use sacrificial polymer carrier layers mostly limited by PMMA or PDMS stamping48, etching49, hot lamination/delamination50, or electrochemical bubbling51 to transfer the graphene patterns to the target substrate. Recent studies have demonstrated direct transfer of graphene on PET substrate based on selective dewetting; however, this method requires UV curable adhesive52. Similarly, CVD grown graphene was also transferred to PVA through drop casting and lamination53 and to PDMS surface via drop casting and peeling approach with low transfer efficiency54. There are some other studies reporting the direct transfer of conductive silver using polymer casting55; however, transfer of graphene via simple polymer casting directly to the target polymeric substrate could be an alternative, simple, fast, green and cost-effective approach allowing the use of variety of flexible substrate materials and eliminating the aforementioned processing steps.In summary, it is difficult to conduct the entire graphene-based flexible electronic device fabrication procedures such as CVD, printing and lithography, directly on the target flexible substrate due to the requirements of harsh chemical, physical or thermal treatments that can significantly deform the target polymeric substrates, particularly natural or synthetic biodegradable polymers possessing low thermal stability28. With this approach, we are able to fabricate high-resolution, small feature sized  conductive graphene patterns/circuits  on various flexible polymeric substrates. Briefly, our method consists of two main steps; (i) the formation of graphene patterns/films on substrates/molds via conventional methods such as CVD, channel filling or ink-jet printing and (ii) direct casting of target substrate polymer solution on the graphene patterned substrates/molds and direct transfer of graphene patterns to the target substrate via peeling off upon drying and film formation. This makes our method versatile allowing the use of different polymers including natural/synthetic, biodegradable/non-biodegradable polymers (such as Poly-L-Lactic Acid (PLLA), Cellulose Acetate (CA), Gelatin (GEL), Poly Lactic-co-Glycolic Acid (PLGA) or Whey Protein Isolate (WPI)) with well-defined characteristics and provides precise control of 3D microstructural and mechanical properties  of the target substrate material with high resolution graphene patterns (feature dimensions of ~5 µm width/depth). This process also enables graphene-based circuit design on biodegradable polymeric films which is not possible with chemically or thermally degrading, lithographic patterning techniques. Our method also requires the use of lithography only for the preparation of molds with high resolution and small feature sized patterns; however, it needs to be done only once. On a broader scale, the use of this new room-temperature facile method to fabricate biodegradable, biocompatible, flexible, and electrically-conductive graphene circuits could pave the way for various applications including tissue engineering, robotics, implantable heart sensors, brain-computer interfaces, or low-cost wearable sensors58.Here we propose a facile and versatile graphene transfer method at room temperature based on simple polymer casting that does not require thermal processing, etching, stamping or UV treatment. This process is simply based on the differences in the surface energies and adhesive forces between the graphene/mold and graphene/target polymer substrate created during the polymeric film formationThe developed method is focused on direct transfer of graphene-based patterns from rigid or flexible substrates to the polymeric flexible films via polymer casting. The method consists of three main steps; (i) Preparation of graphene-based patterns/films via channel filling, ink-jet printing or CVD on rigid or flexible substrates/molds; (ii) Casting of the target substrate polymer solution on the graphene-based patterns/films formed on substrates/molds; (iii) Drying of the solvent and formation of films followed by peeling off the films from the substrate/mold, transferring the graphene pattern from substrate/mold surface to the target polymeric film surface. The application steps of the graphene transfer via direct polymer casting on rigid Delrin and ink-jet printed flexible polyimide substrate was shown in Fig. 22. The difference in the surface energies between the polymer solution and graphene patterns upon the formation of polymeric films makes it possible to remove the graphene pattern precisely from the solid mold surface and transfer the patterned graphene onto the polymeric film surface. The work of adhesion at the graphene-polymer interface is higher than the work of adhesion at the graphene-mold interface, which makes the complete transfer of graphene possible as described in Equation 122d and γp correspond to the dispersion and polar components of surface energy (γ = γd + γp). The theoretically calculated work of adhesion between each material interface support the proposed hypothesis. For instance, WGraphene-Delrin (77 mJ/m2) < WGraphene-PLLA (90 mJ/m2); WGraphene-Delrin (77 mJ/m2) < WGraphene-PLGA (88 mJ/m2); and WGraphene-Delrin (77 mJ/m2) < WGraphene-GEL (88 mJ/m2).This method is also versatile and can be applied to many different polymeric materials including but not limited to PLLA, PLGA, CA, GEL and WPI films Fig. . It was Graphene-Polyimide (86 mJ/m2) < WGraphene-CA (94 mJ/m2). Therefore, this fabrication approach allows formation of high-resolution patterns on the surfaces of versatile polymeric films as long as their free surface energies are sufficiently different to enable strong adhesion to one another. Similarly, we also demonstrated the potential of the direct polymer casting and peeling approach for the transfer of CVD grown graphene as illustrated in Fig. Graphene-Cu = 46 mJ/m2)   using our direct polymer casting approach. After the transfer, we did not observe a significant change in the resistance of graphene films, which indicates the efficiency of the transfer.This graphene transfer approach can also be used to transfer the graphene patterns from one flexible substrate to another. For instance, we already demonstrated that the graphene patterns, ink-jet printed and post-processed , can easily be transferred to 3D microstructured and porous CA-based films using polymer casting approach Fig. . The suce 86 mJ/m < WGraphm2) Fig.  and quarm2) Fig.  as well m2) Fig.  substrat59.The conductivity of graphene was enhanced by applying thermal annealing  prior to the application of graphene solution to create the substrate patterns via the channel filling approach. The transferred graphene via polymer casting has the sheet resistance of ~0.2 kΩ/sq. Similarly, the ink-jet printed and laser annealed graphene was also transferred from flexible polyimide substrate to CA substrate, which has the sheet resistance of ~0.7 kΩ/sq61 along with additional peaks around 282.7 eV, which potentially stems from the existence of graphene layers on the sample  was observed66. Samples show a small D peak associated with lattice structure imperfections and edge plane defects in the graphene, as well as large G/2D peaks characteristic of sp2-hybridizated carbon (graphite/graphene structure)59. Graphene patterns displayed low (IG/ID) ratio (0.33 ± 0.01), which demonstrates that the transferred material is most adequately characterized as a multi-layer graphene structure69.The dense, compact and continuous filling of graphene patterns on the PLLA film surface are shown in Fig. The obtained graphene patterns showed significant stability after multiple bending and washing cycles Fig. . In addiAs an alternative to the flexible electronics, this novel method can also be used to fabricate stretchable electronics. For this purpose, we fabricated flexible and stretchable WPI films with conductive graphene patterns using polymer casting method. The initial length of WPI film with graphene patters was 5 cm Fig. . A cycliThese results clearly demonstrate that it is easy, fast, green and cost effective to fabricate graphene-based flexible and stretchable electronic devices on various biodegradable and biocompatible polymeric flexible substrates with proper 3D microstructural properties, small feature sizes and high-resolutions via the mentioned method. Therefore, the fabricated devices can be used in various biomedical and healthcare applications. Considering this, we conducted additional experiments to demonstrate the potential use of the developed devices with polymer casting method for controlling the stem cell differentiation.73. Very recently, we have demonstrated the successful transdifferentiation of mesenchymal stem cells (MSCs) into Schwann cells (SCs) using electrical stimuli through ink-jet printed and laser annealed graphene circuits on flexible polyimide substrates70. Although the transdifferentiation was successful, the non-biodegradable and non-porous nature of polyimide substrate limits the potential translation of this strategy for surgical implantation and clinical applications related to nerve regeneration. Therefore, the fabrication of such graphene-based circuit on biodegradable, 3D microstructured and porous substrates could make the surgical and clinical application of this strategy possible74. Therefore, the in situ-in vivo precise control on MSCs differentiation, migration and fate commitment upon the surgical implantation could be possible75. Considering this, we developed the same device  on biodegradable, 3D microstructured and porous PLLA films using our graphene transfer via polymer casting approach as described above.In our previous studies, we showed precise control of the transdifferentiation of stem cells using mechanical, physical, topographical, chemical and electrical cues for peripheral nerve regeneration purposes70. Our results showed that almost ~90% of the cells got immunolabelled with the selected SC markers suggesting the successful transdifferentiation . The flexibility in the polymeric material selection along with precise control on substrate properties, enables the potential use of the developed devices in biomedical applications or implantations. In addition, with this approach it is possible to obtain graphene patterns with high resolution, low feature size , high conductivity (sheet resistance of ~0.2 kΩ/sq) and high stability (maintains sheet resistance after 100 bending and 24 h washing cycles). Moreover, this is a novel, facile, versatile, scalable and cost-effective manufacturing approach that eliminates the need for expensive equipment , physical or chemical post-processing or complex transferring/stamping processes. Furthermore, this is an alternative approach to the conventionally used PDMS molding or Cu foil-based transfer methods. Overall, this promising method has the potential to pave the way for flexible and stretchable electronics fabrication, particularly for the biomedical applications including but not limited to brain-computer interfaces, robotics or lab/organ-on-a-chip.46. In another approach, a graphene pattern was ink-jet printed on a rigid or flexible polyimide substrate and thermally or laser annealed (to enhance the conductivity) to create graphene substrates76. In addition, graphene and silver nanowire films were grown on quartz and Cu foil substrates via conventional CVD method77.The Teflon or Delrin substrates with small feature size micropatterns were prepared using computer numerical control (CNC) machine while the silicon wafer molds with small pattern features were prepared using photolithography only once. The pattern dimensions were varied between 5 to 400 µm in width and dept. Graphene nanoplatelet (Sigma Aldrich) solution, sonicated and thermally annealed (at 75 °C) prior to application , was used to fill the micropatterns. The excess graphene on the substrate was removed and cleaned by sticking and peeling a commercially available scotch tape, which leaves the graphene in the patternsFollowing formation of graphene patterns on the substrates using the techniques mentioned above, the polymer casting solutions were prepared at desired formulations and concentrations. The polymer solution (could be Poly-L-Lactic Acid (PLLA), Cellulose Acetate (CA), Gelatin (GEL), Poly Lactic-co-Glycolic Acid (PLGA) or Whey Protein Isolate (WPI)) was then cast on the substrate with graphene patterns/films and left for drying. The 3D microstructure, mechanical properties, porosity and pore size of the films can be adjusted via well-established phase inversion techniques or use of pore forming or plasticizer agents. Upon the film formation, the film was peeled off and graphene patterns were transferred from the substrates to flexible film surface Fig. . This poThe stability of the graphene patterns on the polymer devices was tested through multiple washing, bending and peeling-off cycles. The conductivity of the devices was determined by building up a circuit and measuring resistance. The microstructure of graphene patterns and devices were characterized through scanning electron microscopy (SEM) (FEI Quanta 250 FE-SEM), x-ray photoelectron spectroscopy (XPS) (Amicus XPS) and Raman spectroscopy (Bruker FT-Raman Spectrometer) analysis. SEM samples were sputter coated with 2 nm iridium before the analysis and images were taken using secondary electron mode. Monochromatic Al Kα X-ray source (1486.6 eV) was used in XPS analysis with an electron take-off angle 45° from a normal sampling surface. Survey scans were collected from 10 eV to 1100 eV with a pass energy of 187.85 eV. Raman spectra were collected with a backscattering geometry, 1064 nm Nd:YAG laser and a spot size of about 1 mm.We investigated the potential use of the fabricated devices as biointerfaces to control the stem cell behavior. We used Brown Norway rat mesenchymal stem cell (MSCs), which were provided by Dr. Donald S. Sakaguchi in Genetics, Development and Cell Biology Department at Iowa State University.The MSCs isolation procedure from Brown Norway rats was conducted in accordance with the NIH Guide for the Care and Use of Laboratory Animals guidelines and the principles in the “Guidelines for the Use of Animals in Neuroscience Research” presented by the Society for Neuroscience. All animal procedures had the approval of the Iowa State University Institutional Animal Care and Use Committee and were performed in accordance with committee guidelines.2. MSCs were sub-cultured when they reached 80% confluency approximately every 2–3 days.MSCs were plated in T75 flasks in maintenance media (MM), consisting of α minimum essential medium , 20% fetal bovine serum , 4 mM l-glutamine (Gibco), and antibiotic–antimycotic (Invitrogen) and incubated at 37 °C and 5% CO5 cells were seeded on devices and electrical transdifferentiation into Schwan cell-like phenotypes (SCs) was conducted exactly as described in our previous work70. At the end of transdifferentiation, immunocytochemical analysis was performed on MSCs possessing SC-like phenotypes seeded on the device using selected SC markers, s100, s100β and p75 as described previously70.The ink-jet printed graphene patterns with finger dimensions of 400 µm finger width and 250 µm finger-to-finger spacing were transferred to the surface of PLLA film by applying polymer casting method and the obtained devices were used for electrical transdifferentiation of MSCs. 2 × 10