It has been found that the addition of only 1.2 wt% modified GO  increased the tensile strength, Young's modulus and storage modulus of EVA composites by 80%, 50% and 24%, respectively. These increments surpass many recent claims on relevant composites. Excellent molecular level dispersion was also observed from the fracture surface SEM images. Being amine-rich with high electron-donating capability and mechanically robust, the nanocomposite served as an outstanding tribopositive material, thereby generating 7.49 V and 4.06 μA output voltage and current, respectively, when employed in a triboelectric nanogenerator (TENG). The high electrical outputs led the device to light up 43 blue LEDs instantaneously upon hand pressing, demonstrating that the nanocomposite is indeed a promising candidate for harvesting green energy. Moreover, the nanogenerator displayed outstanding cyclic performance stability (even after 8000 cycles) and environmental durability.While the demand for lightweight high-strength nanocomposites is immense, their progress has been severely limited due to inferior filler dispersion and filler–matrix interface adhesion. This article reports a novel modification of graphene oxide (GO) encapsulated by the copolymer of polydopamine (PD) and polyethylenimine (PEI)  The fabrication of flexible high-performance EVA–graphene oxide nanocomposite triboelectric nanogenerators and the lighting of 43 blue LEDs. Graphene is regarded as the “magic bullet” for the composite world, owing to its remarkably high mechanical, electrical, thermal conductivity and gas and chemical barrier properties. Despite its huge merits and the large amount of published literature on graphene nanocomposites, scientists are still exploring the hidden potentials of this material in order to expand the scope of its application, because the improved properties achieved by graphene are still far lower than the theoretical predictions. Critical challenges that severely limit the continuous advancement of graphene for structural applications is its poor compatibility, non-uniform dispersion and weak interface adhesion with the polymer matrices due to the lack of adequate surface functionalities.1,8–10 Moreover, the strong interlayer van der Waals forces cause the graphene nanosheets to form irreversible agglomerates (or restack), making it extremely difficult to handle. Thus, the prevention of restacking should, therefore, be attained prior to incorporation in any system. Moreover, the level of enhancements has been found to depend significantly on the degree of dispersion, alignment and filler–matrix interface adhesion. Hence, to overcome the above drawbacks and enhance the chemical affinity towards polymer matrices and to successfully harness maximum reinforcement, the surface modification of graphene has become imperative.Among a variety of polymer nanocomposites, polymer–carbon nanomaterial composites, especially carbon nanotubes (CNTs) and graphene-based ones, are of the highest priority of current research interests as they deliver high electrical, thermal, mechanical and barrier properties to the composites at significantly low amounts.e.g., hydroxyl, epoxide, carbonyl and carboxylic acid groups)1,8–10 that can be used as an anchorage for further surface treatments depending on the needs. Here, ethylene vinyl acetate copolymer (EVA) was chosen as the matrix, since it has widespread applications in numerous areas including paints, adhesives, coatings, textiles, cables, automotive plastic fuel tanks, etc.11–14 The GO was modified with polyethylenimine (PEI) through the assistance of polydopamine (PD) coating via the Michael addition reaction, as demonstrated in Therefore, this article aims herein to design a novel surface modification process for graphene oxide (GO) in order to develop mechanically stronger polymer nanocomposites even at significantly low amounts of loading. GO was selected, as it is the most widely preferred derivative of graphene material, containing various oxygenated functional groups .14–17 The working mechanism of this device is based on the contact electrification effect, where energy conversion occurs via surface charge transfer resulting from periodic contact between two dissimilar tribopolar materials, a triboelectric positive and a triboelectric negative material.15–17 Interestingly, the nanocomposite-based TENG showed great performance to the triboelectric response, with an output voltage and current 4 and 4.2 times higher, respectively, than that of the pristine EVA device. Moreover, the EVA nanocomposite TENG showed highly stable performance against severe constant cyclic load and when exposed to ambient conditions for several weeks. Finally, the novelty of this research lies in the concept of the modification of GO, whereby the large reinforcement, excellent molecular-level dispersion and high triboelectric performance of the composite can be achieved concurrently at low filler loading.Besides the structural importance, we explored a new potential application of this composite material as a sustainable energy harvesting device, called a triboelectric nanogenerator (TENG). Recently, energy harvesting and conversion devices using lightweight materials have become an emerging trend in current science and technology, as they play an important role in building and driving self-powered nanodevices. To the best of our knowledge, this is the first article that introduced EVA/GO nanocomposites as a potential futuristic material for triboelectric energy harvesting application. Though graphene–polymer nanocomposites are a new attempt in TENG, Harnchana via several dehydrogenation steps, forming o-quinone units that can instantly cyclize to indole moieties and/or form C–C bonds. The Michael reaction takes place at the o-quinone moieties either at the carbonyl carbon atom or at the β-position of the inherent Michael system in an alkaline medium.18 This leads to the self-polymerization of dopamine and the simultaneous copolymerization with PEI. At the same time, the oxygenated groups on GO can form strong covalent and hydrogen bonds with the PD–PEI layer, making it very suitable in dispersing various solvents and polymer matrices when added.−1. Other characteristic peaks for the C PBM data was replaced with SVG by xgml2pxml:<glyph-data id="z.dbd" format="PBM" resolution="300" x-size="8" y-size="10" xml:space="preserve">00000000000000000000000000000000111111110000000011111111000000000000000000000000</glyph-data><svg xmlns="http://www.w3.org/2000/svg" version="1.0" width="13.200000pt" height="16.000000pt" viewBox="0 0 13.200000 16.000000" preserveAspectRatio="xMidYMid meet"><metadata>Created by potrace 1.16, written by Peter Selinger 2001-2019</metadata><g transform="translate scale" fill="currentColor" stroke="none"><path d="M0 440 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z M0 280 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z"/></g></svg>O and CC bonds, C–OH stretching vibrations, epoxy (1224 cm−1) and C–O vibrations from alkoxy groups were seen at 1725, 1624, 1351, 1224 and 1039 cm−1, respectively.1,10 After modification with the mixture of PD–PEI, some noticeable spectral changes appeared in the measurement. The new absorption peaks at 2930 and 2831 cm−1 corresponded to the nonsymmetric and symmetric stretching vibrations of methylene, respectively, while the new peaks at 1649 and 1569 cm−1 represent the stretching vibrations of NHCO, 1453 cm−1 for the aromatic CC bond, and 1220 cm−1 for the C–N bands.2,19 It is important to note here that the amine groups in PEI (owing to highly branched structure) will also react with the carboxyl and epoxy groups of GO at the same time, resulting in the appearance of amide and C–N bonds in the respective spectrum. The peak at 3320 cm−1 can be ascribed to the N–H group of PEI. To further ascertain the attachment of PEI onto GO and to learn about the surface chemical compositions, XPS analysis was carried out. 20 This result clearly endorsed the successful attachment of PDA and PEI onto the GO surface. While the pristine GO sample had C and O contents of 69.8 and 30.2%, respectively, the PEI–rGO sample showed 80.5% C, 14% O and 5.5% N contents. However, for curiosity when we checked the XPS spectra for the PD–rGO sample, only 3.7% N content was found .FTIR analysis was performed first to verify the chemical changes in GO as a result of the modification process. −1, corresponding to the first-order scattering of the E2g phonon of sp2 C atoms and structural defects induced by the surface functional groups, respectively. It is well-known that the height of the D band implies the size of the inplane sp2 domains.1,3,8 Therefore, the increase of the D peak intensity indicates the formation of more defects in sp2 domains in the structure. As can be seen in ID/IG ratio of PD–PEI–rGO is higher (1.12) than that of pristine GO (0.84), indicating that surface modification had led to an increase in the number of disorders or defects in the graphene sheets.Next, Raman spectra were measured to understand the crystal structure, disorder and defects in the above GO specimens. As can be seen from 1 Following that, the decomposition, which brought about 20–22% more mass loss, was due to the thermal decomposition of the carbon skeleton. On the other hand, the modified GO showed much higher thermal stability, with only 5% weight loss at 240 °C and 28% total loss at 800 °C. The losses at above 240 °C can be attributed to the degradation of the polydopamine and PEI molecules. The TGA results also revealed that there is ∼15–16 wt% of PD–PEI in the modified graphene oxide (PD–PEI–rGO). The higher thermal stability clearly indicates that there are strong interactions between the GO and PD–PEI moieties, also endorsing the above FTIR spectrum that showed the formation of amide bonds between GO and PD–PEI due to the modification. The presence of such large amounts of PD–PEI in GO also suggests that this can be a good filler for polymer nanocomposites.In order to understand the effect of modification on the thermal behavior of the GO, TGA analysis of pristine GO and PD–PEI–rGO was conducted, and the corresponding results were compared in i.e., 0.4, 0.8 and 1.2 wt%) were added into the EVA matrix to fabricate various nanocomposites. E) was obtained after the addition of PD–PEI–rGO into EVA, and the values increased with increasing PD–PEI–rGO content. It was interesting to see that the addition of the same amount of modified GO (0.4 wt%) led to an increase in the T.S of the nanocomposite by 4.4-fold that of the pristine GO/EVA nanocomposite.To explore the potential of PD–PEI–rGO as a filler material in the area of polymer composites, they were blended with ethylene vinyl acetate copolymer (EVA) matrix to create the nanocomposites. The composites were fabricated by solution casting followed by a melt-blending process. The tensile and DMA test results of the composites will be discussed next, as they indicate the capability of the material for structural performance. Low amounts of PD–PEI–rGO  strong interfacial adhesion between the matrix and the modified graphene nanosheets (owing to strong hydrogen bonding between the carbonyl of EVA and the amine of the PEI), (ii) effective load transfer from the polymer matrix to the graphene nanosheets, and (iii) homogeneous dispersion of the graphene nanosheets throughout the polymer matrix. The higher mechanical strength can also be explained by the greater mechanical interlocking between the EVA polymer chains and modified graphene oxide nanosheets, which possess higher surface roughness. Interestingly, these values were found to be much higher than that of many earlier reports on EVA nanocomposites  is considerably lower than that of the nanocomposites. After the incorporation of GO, regardless of the type (unmodified or modified), the storage modulus of the nanocomposites increased significantly. While the storage modulus of neat EVA was found to be 1981 MPa, nanocomposites containing 0.4 wt% pristine GO displayed a storage modulus value of 2160 MPa, which was ∼9% higher than that of the neat EVA. Interestingly, when the same amount of modified graphene oxide was added , the storage modulus reached to up to 2320 MPa (17% higher), nearly double that of the pristine GO-containing nanocomposite. The gradual addition of PD–PEI–rGO  further increased the storage modulus to 2377 and 2457 MPa, which was ∼20% and 24% higher than the storage modulus of the neat EVA, respectively. Such large increments in the mechanical strengths of the EVA nanocomposites with the use of such a small amount of graphene oxide (1.2 wt%) clearly suggest the great potential of our functionalization process and that of modified GO. The tan δ maximum value, which corresponds to the Tg of the polymer, was also noted to increase with the addition of PD–PEI–rGO  are insulating in nature, the generated charges at each surface would, therefore, remain there , leading to no electric potential difference between the two electrodes. As soon as the compression force is released, a strong potential difference between the top and bottom electrodes was created due to the electrostatic induction effect or strong dipolar difference. As a consequence, the electrons began to flow from the negative to the positive potential through the external circuit, resulting in the experimentally observed positive electric output signal  on the screen. The electron transfer occurs until an electrostatic equilibrium is established. Upon further pressing the TENG, the electrostatic equilibrium was disrupted, and the backflow of electrons took place from the bottom electrode to the top electrode until a new equilibrium was reached. This resulted in the generation of the opposite current, thereby causing reversed electrical signals to be observed. This is considered as the second half of the triboelectric charge cycle.To understand the multifunctional role of the EVA nanocomposite, especially for energy harvesting application, triboelectric measurements were recorded. The construction of TENG is shown in Voc) and the short-circuit current (Isc) between the pristine GO/EVA nanocomposite and the PD–PEI–rGO/EVA nanocomposite-based TENGs, respectively. While the peak output voltage and current for pristine GO/EVA nanocomposites were about 1.89 V and 0.97 μA, respectively, the PD–PEI–rGO/EVA composite TENG showed electrical outputs of about 7.49 V and 4.06 μA. These values were ∼4 and ∼4.2 times higher than those of the output voltage (1.89 V) and current (0.97 μA) of pristine EVA, respectively. The higher electrical output for the PD–PEI–rGO/EVA nanocomposite TENG can be attributed to the higher surface polarity and electron-donating ability owing to the presence of the amine (from PEI) and OH groups (from DA) when rubbed with a fluoropolymer. Moreover, the embedded GO nanoplatelets in the composite can act as charge trapping sites, thereby promoting the relative dielectric property of the film, leading to an increase in the triboelectric output performance of the composite.14 To verify the feasibility of the above nanogenerator for practical uses, we connected 43 blue LEDs directly  to the output of the nanogenerator without any circuit. i.e., dancing or jumping.Next, the cyclic performance stability and time-dependent durability of the above TENG was tested, as it determines the quality and shelf-life of the device. From We reported a novel and effective approach for GO modification to produce mechanically strong EVA nanocomposites with multifunctional capability. The surface modification was carried out by a unique mussel-inspired polydopamine-assisted Michael addition reaction. This approach showed greater promise in terms of molecular level dispersion and the interface adhesion of graphene nanosheets within the EVA matrix owing to strong hydrogen bonding interactions. As a result, the achieved polymer nanocomposites showed radical improvements in the tensile strength (80%) and storage modulus (24%) even at only 1.2 wt% of graphene content, clearly demonstrated the robustness of the functionalization process. Besides the high reinforcing capability, the addition of modified GO (PD–PEI–rGO) turned the nanocomposite into an excellent triboelectric material owing to the amine-rich surfaces. As a result, the composite triboelectric nanogenerator showed outstanding electrical responses with an output voltage of 7.49 V and current 4.06 μA, thereby lighting up the 43 blue LEDs instantaneously. Such a high output voltage also implied that our device would be very effective for large-scale industrial applications. This also suggests that the performance of the device can be amplified further by increasing the graphene content as well as the degree of amine functionalization. Considering the above fascinating achievements, it can be concluded that our functionalization strategy was indeed a versatile platform for GO modification and can hopefully attract great attention from various sectors to develop other lightweight and strong composites for wider applications.Flaky graphite powders (300 mesh) were purchased from Huadong Graphite Factory . Dopamine, Trizma base and polyethylenimine  with an average molecular weight of 25 000 were purchased from Sigma Aldrich, South Korea. Ethylene vinyl acetate copolymer  was obtained from India. Polyvinylidene fluoride (PVDF) film was purchased from Starlab Scientific Co., Ltd, Korea.1 The as-synthesized GO was then thoroughly dispersed in 10 mM Tris (pH 8.5) buffer solution using a probe sonicator at 20 Hz frequency for 30 min. Next, an equal amount of dopamine (DA) and PEI (1 : 1) was added into the GO solution, and the mixture was magnetically stirred for 6 h at 50 °C. After that, the mixture was filtered by vacuum filtration and washed with large amounts of water and methanol repeatedly in order to remove any unattached DA and PEI. Then, the product was freeze-dried overnight and labelled as PD–PEI–rGO. The ‘r’ represents the partial reduction of graphene oxide.Firstly, GO was synthesized by a modified Hummers method from graphite flakes following our earlier report.Various PDA–PEI–rGO/EVA nanocomposites were prepared by solution casting followed by the melt blending process. Briefly, PD–PEI–rGO of 100, 200 and 300 mg were dispersed separately in 25 ml of toluene by ultrasonication at room temperature for 1 h. Concurrently, EVA (25 g in 100 ml) was dissolved in toluene separately. Next, the above PD–PEI–rGO dispersions were added separately into the EVA solution and mixed thoroughly with constant stirring at 100 °C for 8 h and then cast on a glass Petri dish. After complete drying in air followed by vacuum, the above three different nanocomposite films  were cut into small pieces and melt blended in a Haake mixture (MiniLab) at 145 °C to obtain the desired composites. Dicumyl peroxide  was used as a curing agent during the blending process. The extruded strips were then made into sheets (1–1.2 mm) by compression molding in a hot press at a temperature of 150 °C with a holding pressure of ∼10 MPa for 5 min. For comparative study, a GO/EVA nanocomposite was also fabricated under identical conditions.i.e., EVA nanocomposite and PVDF films, used to fabricate the device. The process of fabrication is as follows: a 1.7 cm × 1.7 cm × 1 mm piece of PD–PEI–rGO/EVA nanocomposite film was prepared. Then, the film was attached to the center of a strip of copper tape, which acted as the top electrode. The Cu tape was further attached to a PET film in order to protect the electrode from damage and as a hard supporting layer for easy handling of the devices. On the other hand, a PVDF film of the same size (but the thickness was 0.25 mm) was taken, and one side was coated with aluminum using a metal ion evaporator. This became another electrode for the device. The coated PVDF was similarly attached to a PET support. Then, the device was assembled by placing the two active surfaces face-to-face, separated by a 2 mm soft spacer and connected to the external circuit through copper wires. For comparison, a pristine EVA/PVDF TENG had also been fabricated keeping all parameters the same as earlier. It is important to note here that for the triboelectric experiment a nanocomposite of higher PD–PEI–rGO content  was used.All TENGs used in this study were vertical contact types. There were two active materials, −1 resolution. The surface elemental compositions of the above samples were determined by X-ray photoelectron spectroscopy (XPS) on a K-Alpha Thermo Scientific spectrometer with monochromatic Al Kα X-ray radiation. All XPS spectra were corrected according to the C 1s line at 284.8 eV. The Raman spectra of GO and PD–PEI–rGO were measured by an Alpha 300R WiTec confocal Raman microscope. The thermal stability of all specimens was tested using a TA Instruments TGA 2950 from 30 to 800 °C at a heating rate of 10 °C min−1 under nitrogen gas medium. The tensile tests of the EVA nanocomposites were performed at room temperature with a crosshead speed of 50 mm min−1. The storage modulus of the composites was measured by a dynamic mechanical analysis instrument (DMA Q800 TA) in the tension mode at a heating rate of 5 °C min−1. The surface morphology of the tensile-fractured samples was observed by field emission scanning electron microscopy  after coating with platinum. Triboelectric performance of the TENGs was measured as follows: the device was fixed onto a metal zig, and a 20 Hz square wave input was applied to the device from a shaker , which was controlled by a function generator (Keysight 33220A). The bottom electrode was pressed up by a shaker with 14 N force, while the top electrode was fixed onto a lab bench. The input force was measured by a force transducer (Brüel & Kjær 8230-001). The output voltage and current data were collected using a pulse analyzer (Brüel & Kjær 3560-B-030) and PC.The presence of surface functional groups of the pristine GO and modified GO powder were analyzed by Fourier transform infrared spectroscopy  spectra at a scan rate of 16 with 4 cmWe have no conflict of interest.NA-001-C9NA00142E-s001


However, conventional exfoliation fabrication methods and layer‐by‐layer transfer techniques have various limitations. The CVD synthesis of high‐quality large‐area graphene and hBN multilayer heterostructures is essential for the advancement of new physics. Herein, the authors propose an in situ CVD growth strategy for synthesizing wafer‐scale AAB‐stacked single‐crystal graphene/hBN/graphene trilayer van der Waals heterostructures. Single‐crystal CuNi(111) alloys are prepared on sapphire, followed by the pre‐dissolution of carbon atoms. Single‐crystal monolayer hBN is synthesized on a plasma‐cleaned CuNi(111) surface. Then, a single‐crystal monolayer graphene is epitaxially grown onto the hBN surface to form graphene/hBN bilayer heterostructures. A controlled decrease in the growth temperature allows the carbon atoms to precipitate out of the CuNi(111) alloy to form single‐crystal graphene at the interface between hBN and CuNi(111), thereby producing graphene/hBN/graphene trilayer van der Waals heterostructures. The stacking modes between as‐grown 2D layers are investigated through Raman spectroscopy and transmission electron microscopy. This study provides an in situ CVD approach to directly synthesize large‐scale single‐crystal low‐dimensional van der Waals heterostructures and facilitates their application in future 2D‐material‐based integrated circuits. An in situ CVD synthesis of high‐quality single‐crystal graphene/hBN/graphene trilayer van der Waals heterostructures is presented here. A single‐crystal CuNi(111) film is produced on sapphire and then subjected to carbon dissolution. Single‐crystal monolayer hBN and graphene are then synthesized. Carbon atoms diffuse to the hBN–CuNi(111) interface and form graphene. This study paves the way for future 2D‐material‐based large‐area integrated circuits. The van der Waals heterostructures consisting of graphene and hBN used in most studies were fabricated using mechanical exfoliation and layer‐by‐layer transfer techniques. The exfoliation of single or multilayer graphene and hBN from bulk materials preserve the high crystal quality and intrinsic physical properties of thin materials. However, disadvantages such as small sample size and interlayer chemical contamination limit their applications in scientific research and industrial applications. The chemical vapor deposition (CVD) synthetic technology has facilitated the production of high‐quality, large‐area, single‐crystal monolayer graphene and hBN. Many researchers have synthesized graphene and hBN heterostructures using the CVD method. Currently, graphene/hBN, hBN/graphene, and hBN/graphene/hBN have been grown on Cu substrates in situ via CVD. Although the graphene/hBN/graphene heterostructure is considered an ideal system to investigate thermoelectric transport in the cross‐plane direction and has demonstrated significant application in tunnel transistors, the in situ CVD growth of high‐quality graphene/hBN/graphene heterostructures has been rarely reported. Hence, there is an urgent need for a CVD growth method for fabricating large‐area single‐crystal graphene/hBN/graphene heterostructures to be used in both fundamental research and industrial applications.2D van der Waals heterostructures are currently regarded as promising materials owing to their special electronic structures.4, 5FigurexNi1‐x film. Raman spectroscopy and transmission electron microscopy (TEM) with selected area electron diffraction (SAED) were used to verify the crystal quality of the as‐grown graphene and hBN layers and to identify the stacking mode of the heterostructures. The proposed in situ CVD growth method can be extended to the synthesis of various multilayer graphene‐based van der Waals heterostructures.In this study, we synthesized wafer‐scale single‐crystal graphene/hBN/graphene multilayer van der Waals heterostructures with an AAB‐stacking mode via in situ CVD growth. First, to create an ideal substrate, we fabricated single‐crystal CuNi(111) films on a two‐inch sapphire wafer Figure1. Carb22.12O3(0001) wafer was treated with an acid mixture to remove scratches and residual stress by etching the surface, producing an atomically smooth surface consisting of a terrace‐and‐step structure. Then, 450‐nm‐thick Cu followed by 50‐nm‐thick Ni were deposited on the Al2O3(0001) substrate through e‐beam evaporation to form a Cu90Ni10 alloy on Al2O3. After annealing at 1050 °C for 2 h in the CVD system with flowing hydrogen (50 sccm) and argon (350 sccm)  substrate  film. Hence, the hydrogen‐argon plasma was used to remove the graphene islands, which created a clean surface of the CuNi(111) film. Subsequently, ammonia borane was heated to 80 °C and then introduced into the CVD growth chamber under Ar carrier gas flow, leading to the formation of single‐crystal hBN islands on the as‐fabricated CuNi(111)/sapphire substrate at 1050 °C  of 15.2 cm−1, indicating that as‐grown hBN is a high‐quality monolayer.By utilizing the specific solubility of carbon in the CuNi alloy,33 metC Figure . Some wh2.2Figure To decrease the influence on the pre‐grown hBN layer, we conducted the fast growth of top‐layered graphene to let the graphene film quickly cove the hBN surface as a protective layer. The SEM image showed a uniform surface with no apparent adlayers on the as‐grown graphene/hBN heterostructure  belonging to graphene, confirming the co‐existence of hBN and graphene  and the E2g (1361 cm−1) peak of hBN . The FWHM of the 2D peak of as‐grown graphene was ≈39 cm−1, indicating that the mismatch angle between graphene and hBN is <1°. Thus, from the Raman spectroscopy results, the stacking mode between the hBN layer and top‐layered graphene is regarded as AA stacking. The analysis of the electron diffraction patterns of the bilayer heterostructure allows to gain a deeper understanding of the relative orientation of the two films. Therefore, we separated the as‐fabricated graphene/hBN film from the CuNi(111)/Al2O3 substrate and transferred it onto a TEM grid to collect the SAED patterns. The SAED pattern in Figure  The SAED pattern corroborates with the Raman spectrum, indicating that as‐prepared graphene/hBN represented an AA stacking mode between layers.After growing the high‐quality hBN monolayer, we synthesized single‐crystal monolayer graphene directly onto the hBN/CuNi(111) surface to form the graphene/hBN heterostructures by re‐introducing methane into the CVD system at 1030 °C Figure3a. Thise Figure . A flat e Figure . Raman mr Figure . A typice Figure . The 2D N Figure , which i2.3 ≈1.3% between graphene and Ni(111); and ≈1.6% between graphene and hBN), under the interplay of the under‐layered single‐crystal CuNi(111) surface and upper‐layered single‐crystal hBN monolayer, the precipitated carbon atoms nucleated and epitaxially grew to form a single‐crystal monolayer graphene film at the interface, resulting in the formation of graphene/hBN/graphene trilayer van der Waals heterostructures  and 2D (2706 cm−1) signals assigned to graphene and the E2g signal associated with hBN  film gradually precipitated to the interface between CuNi(111) and hBN. To prevent the formation of BCN during the entire process, we optimized the hBN growth parameters to ensure that the pre‐grown hBN monolayer had fully covered the CuNi(111) surface. Considering the small lattice mismatches among graphene, CuNi(111), and hBN (3%–4% between graphene and Cu (111);45 ≈1. Figures . A smoot Figures . The typN Figure . The blu−1 was fitted by two Lorentzian peaks centered at ≈1352 cm−1 (D peak of graphene) and 1362 cm−1 (E2g peak of hBN) , and a higher component of carbons in the CuNi film can lead to the growth of more graphene layers at the interface, resulting in the formation of van der Waals heterostructures with multilayered graphene under the hBN film (Table The peak at ≈1359 cm) Figure . The inc  film on a sapphire wafer and subjected it to carbon dissolution to incorporate a specific concentration of carbon atoms into the CuNi(111) film. A high‐quality single‐crystal hBN monolayer was prepared on the surface of the CuNi(111) alloy after plasma cleaning. Following this, a monolayer graphene film was grown on top of hBN to form an AA‐stacked graphene/hBN heterostructure. Subsequently, a high‐quality single‐crystal graphene monolayer was formed at the interface between the hBN layer and CuNi(111) surface through a controlled decrease in the temperature of CVD growth, resulting in the formation of an AAB‐stacked graphene/hBN/graphene trilayer van der Waals heterostructure on the CuNi(111)/sapphire substrate. This study introduces a strategy for the contamination‐free in situ CVD synthesis of wafer‐scale single‐crystal graphene/hBN/graphene van der Waals heterostructures. The proposed method will provide a foundation for the development of more promising functionalized 2D materials for industrial applications.42O3(0001), 2 in., c‐plane, and double‐side polished from 11‐D Tech) was cleaned using hot acetone, isopropanol, and deionized water for 5 min. A mixed acid solution (H2SO4:H3PO4 = 3:1) was used to etch the surface at 320 °C for 30 min. The sapphire was then rinsed with deionized water and cleaned using oxygen plasma. Subsequently, Cu and Ni films were deposited on the Al2O3 (0001) wafer. The resulting CuNi/Al2O3 was loaded into the CVD system for annealing at 1050 °C under H2  and Ar  atmospheres at a pressure of 750 Torr for 2 h. After annealing, a single‐crystal CuNi (111) film was formed on the Al2O3(0001) wafer.The sapphire substrate  as a precursor, with H2 (15 sccm) and Ar (7 sccm) as the carrier gases. During the CVD growth process, the precursor was heated and maintained at 80 °C. After the growth of hBN, upper‐layered graphene was synthesized on the hBN surface in a CVD system at 1030 °C for 20 min, using a mixture of CH4  and H2 (20 sccm). After the growth of upper‐layered graphene, the chamber temperature was slowly decreased to 1000 °C over 30 min under CH4 (5 sccm) and H2 (10 sccm) gas flow. After the completion of the CVD growth process, the sample was removed from the furnace zone and quickly cooled under the same gas flow and pressure.hBN was synthesized on the CuNi(111) surface in a CVD system for 40 min at 1050 °C using ammonia borane  as a protective layer; an aqueous NaOH solution was used as the electrolyte. Upon applying a constant current with CuNi(111) as the cathode, H2 bubbles were formed at the interface between the CuNi(111) substrate and the PMMA/graphene/hBN/graphene film, leading to the detachment of the PMMA/graphene/hBN/graphene layers from the CuNi(111)/Al2O3 substrate. After rinsing with deionized water, the stacked layers were transferred to the target substrate. Eventually, PMMA was removed using hot acetone, rinsed with isopropanol, and dried in air.The as‐grown heterostructures were transferred onto a SiO2O3(0001) substrates. AFM  was used to determine the surface roughness and layer thickness. TEM images and SAED patterns were acquired using a transmission electron microscope  at an acceleration voltage of 200 kV.Raman spectra and maps were obtained by the confocal Raman spectroscopy  using 488 and 532 nm lasers. SEM  equipped with EBSD was used to characterize the surface morphologies of the CuNi alloy and the as‐grown 2D materials. XRD  patterns were obtained for the Cu(111) and CuNi(111) films on the AlThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


SSZ‐43 crystals are intergrowths of two polytypes: ≈90 % orthorhombic polytype A with ABAB stacking of the 12‐rings, and ≈10 % monoclinic polytype B with ABCABC stacking. Molecular modeling performed on the idealized Si‐SSZ‐43 structure along with empirical relationships for zeolite selectivity in boron‐ and aluminum‐containing synthesis gels were used in a combined approach to design new di‐quaternary ammonium organic structure‐directing agents (OSDAs). Experimental trials demonstrated that the new OSDAs produced SSZ‐43 over a broader range of compositions than previous mono‐quaternary OSDAs.The structure of zeolite SSZ‐43 was determined by 3D electron diffraction, synchrotron X‐ray powder diffraction, and high‐resolution transmission electron microscopy. The SSZ‐43 framework forms one‐dimensional, sinusoidal 12‐ring channels from 5 2 compositions.After two decades since its discovery, the structure of zeolite SSZ‐43 has been determined as a novel framework comprised of a primitive orthorhombic crystal system with monoclinic faulting that forms unique, 1D 12‐ring sinusoidal channels. A combination of molecular modeling and empirical relationships for zeolite selectivity was used to design two new di‐quaternary OSDAs and produce SSZ‐43 across a wide range of high SiO Standard Oil Synthetic Zeolite‐52), a small‐pore, cage‐containing, aluminosilicate zeolite with 8‐ring windows was discovered by Lee and Zones at Chevron in 2001,SFW by the Structure Commission of the International Zeolite Association) was determined in 2013.CHA topology, and signaled its potential use as an emissions abatement catalyst. Initially, commercial scale‐up of SSZ‐52 was hindered by the complexity of the original OSDA, but computational modeling efforts later identified alternative organo‐cation guest molecules to permit a more economical synthesis route.SFE) and SSZ‐31 (*STO). However, since its discovery over two decades ago, the structure of SSZ‐43 remained unsolved.Zeolites are widely used in catalytic, adsorption, and separation processes.6, 7 Tc‐axis in a sinusoidal manner, making this structure most unusual among other known 1D, large‐pore zeolites. Additionally, we report a new computationally guided synthesis route for SSZ‐43 that was developed by combining molecular modeling of the idealized SSZ‐43 framework structure with experimentally derived relationships in zeolite phase selectivity for boron‐ and aluminum‐containing synthesis gels. This approach was used to design two different di‐quaternary ammonium OSDAs, which were identified according to their stabilization energies in the idealized framework structure and further validated through laboratory experiments. Using the newly discovered OSDAs, both boron‐ and aluminum‐containing, high‐silica  SSZ‐43 zeolites could be synthesized in hydroxide (OH−) media.Developing new or modified medium‐ (10‐ring) and large‐pore (12‐ring) zeolites is often key to enabling significant advances in refining and petrochemistry because the confining void spaces and pore architectures of these frameworks are often well suited for interactions with small hydrocarbon molecules in fuels, lubricants, and petrochemicals. Thus, there is considerable interest in understanding the structural details of novel zeolites, such as SSZ‐43. However, as is often the case with synthetic zeolites, samples of SSZ‐43 are polycrystalline and prohibit structural solution by established single‐crystal X‐ray diffraction techniques. In addition, SSZ‐43 is a partially disordered material, making structural derivation challenging from powder X‐ray diffraction (PXRD) patterns alone. Fortunately, recent developments in three‐dimensional electron diffraction (3D ED) in combination with high‐resolution transmission electron microscopy (HRTEM)15, 166 cathode operating at 200 kV using Instamatic software.2 adsorption data for the calcined B‐SSZ‐43 sample (OSDA Molecule 3), are provided in Section 2 and Figures S1 and S2 of the Supporting Information. From 3D reciprocal space reconstructed from the cRED data using REDp software,a=12.51, b=29.03, c=17.30 Å. From the 2D slices, the reflection conditions can be deduced as the following: h0l: l=2n; 0kl: k=2n, which gives two possible space groups Pbc21 (No. 29) and Pbcm (No. 57). The highest centrosymmetric space group Pbcm was chosen for structure determination. Datasets from seven different crystals were merged to increase the data completeness and redundancy  by using SHELXTEdtool.Pbcm space group. The structure was further refined by SHELXLInitially, the average structure of SSZ‐43 was solved in the orthorhombic space group 461butterfly units connected via 4‐rings similar to *STO . The butterfly unit can be constructed using two mel [415262] composite building units (CBUs), one jbw [64], and two incomplete cages. As shown in Figure butterfly units along the a‐axis forms the SSZ‐43 5461 slab. The neighboring 5461 slabs are rotated by 180° around the a‐axis in the orthorhombic polytype A and shifted by 1/3a in the monoclinic polytype B, and further connected to form layers containing 12‐rings  or down (D) in a UDUUDUDUDDUD configuration. The layers are further connected to the next sheet with the same configuration via 6‐rings. Thus, one unit cell of SSZ‐43 contains two 6‐ring “honeycomb” blocks that are either shifted by ±1/2b relative to one another in the orthorhombic structure, or by ±c in the monoclinic structure as shown in Figure The framework structures of both polytype A and B of SSZ‐43 contain 14 unique T‐atoms  and can be assembled by using the 5c‐axis, which are arranged in an ABAB stacking sequence along the b‐axis . This disorder results in the appearance of sharp spots for reflections with h=3n and streaks for reflections with h≠3n.The orthorhombic SSZ‐43 framework hosts sinusoidal channels with 6.5×6.5 Å apertures along the n Figure b indicatλ=0.99995 Å) was used for further analysis. The orthorhombic unit cell parameters of a=12.324(2), b=28.435(1), c=16.898(1) Å as derived from profile fitting, were used to refine the structural model obtained from cRED data, with restraints applied to all Si−O distances (1.61 Å) and O−Si−O angles (109°). The refined atomic coordinates were then used to generate the base layer for DIFFaX simulationb‐axis for the original orthorhombic unit cell, the layer in parallel to the original ac plane was set as the base layer and the stacking vector was set as 1/2b. Then, 1/3a and 2/3a shifts were applied to generate the second and the third layers, respectively. To ensure that the desired layer stacking sequence was produced, the DIFFaX‐simulated PXRD patterns for the structurally ordered end‐members, i.e., orthorhombic polytype A (ABAB‐stacking) and monoclinic polytype B (ABCABC‐stacking), were verified by comparison to the calculated PXRD patterns from the refined ED structural models as shown in Figure θ=4.65° (d=12.32 Å) shifts to higher angles and merges with that of the (110) reflection at 2 θ=5.07° (d=11.31 Å) to form a broad peak at ≈5° when α≥0.2. Comparison of the synchrotron PXRD pattern for SSZ‐43 with the simulated data  within SSZ‐43 and the other zeolite structures considered in this study are not known. Additionally, we limited the modeling to van del Waals (vdW) interaction energies between OSDA molecules and the target zeolite framework without including other inorganic cations  or H2O molecules in the calculation. First, we determined the optimized vdW interaction energies between the six previously known OSDAs for SSZ‐43n=4 OSDA molecules per unit cell. The optimal OSDA loadings agreed with experimental results obtained by thermogravimetric analysis (TGA). As shown in Table −1 with the decahydroquinolinium‐derived organo‐cations (Molecules 4–6) being most favorable.After the structural models of the SSZ‐43 polytype A and B were derived, a molecular modeling study on OSDA‐zeolite interaction was performed to design new, energetically favorable OSDAs for this framework, and to computationally guide the experimental synthesis of borosilicate and aluminosilicate SSZ‐43 by using these identified molecules (Supporting Information Section 3). We note that pure SiOSTO, polymorph I), a one‐dimensional large‐pore zeolite, and SSZ‐35 (STF), a one‐dimensional medium‐pore zeolite with large internal cavities, are two zeolites that also crystallize from gel compositions and synthesis conditions similar to those used for SSZ‐43. From a structural standpoint, both SSZ‐43 and *STO contain the same mel [415262] composite building units, these units form the 5461butterfly units that are further connected to themselves to assemble the same building layers. The vdW interaction energies between OSDA Molecules 1–6  and SiO2/B2O3 (SBR) in Table inter‐zeolite comparisons of interaction energies for a given OSDA are not straightforward since the zeolite host framework density (FD) and the percentage of framework atoms available to participate in vdW interactions could vary between different zeolite frameworks.[STF) has the lowest FD (16.9 T×10−3 Å−3) among the three zeolites considered . Furthermore, SSZ‐35 (STF) is also the only example here where all framework atoms are exposed to the one‐dimensional channel to interact with the occluded OSDA molecules. Thus, as shown in Table 2 SSZ‐43 structure or *STO, then the modeling results in Table STO. However, neither of these structures could be produced experimentally with Molecule 8, because this OSDA is much more energetically favorable in STF so that SSZ‐35 could be formed under a broader range of synthesis conditions and gel compositions.As indicated by the modeling predictions and experimental results in Table −1 and prevents the crystallization of SSZ‐43 across a broad range of synthesis gel compositions. However, in the case of SSZ‐31, the ≈0.5 kJ (mole Si)−1 interaction energy increase between Molecules 5 and 7 is not significant enough to affect the product phase selectivity. These results demonstrate that both OSDA‐zeolite stabilization energies and synthesis gel chemistry must be considered simultaneously when designing new molecules or identifying conditions to steer selectivity toward a zeolite of interest while preventing the formation of more thermodynamically stable, undesired structures.Often a minor modification to the OSDA can significantly impact the vdW interaction energy between the guest molecule and framework host, resulting in phase selectivity changes for the crystalline products. For example, OSDA Molecules 5 and 7 , SSZ‐31 prefers to crystallize from high Si/heteroatom gels, while SSZ‐35 is generally favored at lower Si/heteroatom ratios. Thus, if OSDAs with more favorable vdW interaction energies in SSZ‐43 can be developed, then the experimental conditions to synthesize SSZ‐43 may be expanded to improve its selectivity across broader ranges of gel compositions and avoid formation of the competing *T>120 °C, autogenous pressures) during hydrothermal synthesis. Organo‐cations with two charge centers  were preferentially screened in an attempt to steer the crystallization products toward SSZ‐43 and away from *STO and STF. Since the periodicity of SSZ‐43 along the 12‐ring pore (c‐axis) is ≈16.9 Å for the orthorhombic polytype A, or approximately twice the length as monoclinic *STO  and STF , we hypothesized that longer chain di‐quaternary OSDAs could more effectively stabilize SSZ‐43 by providing a better match of size and shape between the guest molecule and host zeolite. For each simulation, the minimum vdW interaction energy between the di‐quaternary OSDA and the idealized all‐silica SSZ‐43 structure was calculated by using the molecular modeling procedure described in Section 3 of the Supporting Information, and then used to rank the OSDA candidates from most‐to‐least energetically favorable. Based on the combined modeling predictions and experimental results for Molecules 1–6 in Table −1  were excluded from the ranking. Next, the vdW interaction energies of these remaining di‐quaternary OSDA candidates in the all‐silica STF and *STO frameworks were computed. As previously discussed, the interaction energies for Molecules 1–12  and SSZ‐43 structures. So here we elected, as a starting point for the screening process, to exclude di‐quaternary OSDAs with ESTF<2×ESSZ‐43 and E*STO<ESSZ‐43 from consideration.To demonstrate this predictive capability for SSZ‐43, we initiated a search for new OSDAs by screening an internal library of existing quaternary ammonium compounds that have been used successfully to synthesize zeolites and remain stable under the harsh conditions −1, respectively, with optimal loadings of n=2 OSDAs per unit cell for both OSDA‐zeolite systems. The lowest energy configurations of Molecules A and B inside the one‐dimensional sinusoidal channels of SSZ‐43 as predicted by molecular modeling are shown in Figure Of the original ≈300 candidate OSDAs from the internal library, five di‐quaternary molecules were selected for experimental trials to synthesize B‐SSZ‐43 and Al‐SSZ‐43 according to the method described in Section 1.2 of the Supporting Information. Out of these five molecules, two OSDAs (denoted Molecule A and Molecule B) successfully produced pure SSZ‐4318H32N42+).SEM images along with the PXRD pattern of the as‐made Al‐SSZ‐43 sample synthesized from the new di‐quaternary OSDA Molecule A  indicate that the crystallites exhibit morphological and structural features similar to those in the B‐SSZ‐43 sample synthesized from the original mono‐quaternary OSDA Molecule 3 Figure a. TGA reSTO) and SSZ‐35 (STF) phases could not be avoided in either very high SAR (>300) or low SBR (<40) gels. For example, SSZ‐31 (*STO) crystallized in the presence of Molecule A from pure silica (SAR=∞) gel, but SSZ‐43 formed as Al(OH)3 was added to the synthesis gel (SAR<500). By contrast, SSZ‐43 was the dominant product in the borosilicate system with gel compositions of SBR>40, but at higher B concentrations (SBR ≈10–40) product selectivity shifted toward SSZ‐35 (STF). Similarly, Molecule B also produced B‐SSZ‐43, but required the use of seed crystals in the aluminosilicate gel to crystallize Al‐SSZ‐43 and eliminate the formation of ZSM‐12 (MTW) as a frequently observed by‐product. Interestingly, the interaction energy of Molecule 3, the original mono‐quaternary OSDA, was predicted to be ≈0.2 kJ (mole Si)−1 more favorable in SSZ‐43 than Molecule B  and physical  factors affect the kinetics of zeolite nucleation and growth. In such cases, empirical relationships for zeolite phase selectivity are useful supplements to computationally guided OSDA design.While Molecule A successfully produced SSZ‐43, formation of competing SSZ‐31  constructed from connecting 5*STO) and SSZ‐35 (STF). By using the idealized structural model of SSZ‐43, molecular modeling studies were performed to screen and identify thermodynamically favorable OSDA candidates based on the minimum vdW interaction energy between the guest molecule and host framework. Experimental trials showed that two predicted new di‐quaternary OSDAs could selectively produce SSZ‐43 across a broader range of compositions  than the previously known mono‐quaternary compounds. The strategy reported here demonstrates that consideration of both OSDA‐zeolite stabilization energies and gel chemistry are needed for rational design of new structure‐directing molecules to steer zeolite phase selectivity in systems that are known to produce multiple competing zeolite structures under similar synthesis conditions.An ongoing chemical challenge for synthesizing SSZ‐43 as an aluminosilicate in thermally stable high SAR compositions that are relevant for catalysis, adsorption, and separations is preventing the formation of competing phases, such as SSZ‐31  should be addressed to the authors.Supporting InformationClick here for additional data file.


Mass production and commercial adoption of graphene-based devices are held back by a few crucial technical challenges related to quality control. In the case of graphene produced by chemical vapor deposition, the transfer process represents a delicate step that can compromise device performance and reliability, thus hindering industrial production. In this context, the impact of poly(methyl methacrylate) (PMMA), the most common support material for transferring graphene from the Cu substrate to any target surface, can be decisive in obtaining reproducible sample batches. Although effective in mechanically supporting graphene during the transfer, PMMA solutions needs to be efficiently designed, deposited, and post-treated to serve their purpose while minimizing potential contaminations. Here, we prepared and tested PMMA solutions with different average molecular weight (AMW) and weight concentration in anisole, to be deposited by spin coating. Optical microscopy and Raman spectroscopy showed that the amount of PMMA residues on transferred graphene is proportional to the AMW and concentration in the solvent. At the same time, the mechanical strength of the PMMA layer is proportional to the AMW. These tests served to design an optimized PMMA solution made of a mixture of 550,000 (550k) and 15,000 (15k) AMW PMMA in anisole at 3% concentration. In this design, PMMA-550k provided suitable mechanical strength against breakage during the transfer cycles, while PMMA-15k promoted depolymerization, which allowed for a complete removal of PMMA residues without the need for any post-treatment. An XPS analysis confirmed the cleanness of the optimized process. We validated the impact of the optimized PMMA solution on the mass fabrication of arrays of electrolyte-gated graphene field-effect transistors operating as biosensors. On average, the transistor channel resistance decreased from 1860 to 690 Ω when using the optimized PMMA. Even more importantly, the vast majority of these resistance values are distributed within a narrow range (only ca. 300 Ω wide), in evident contrast with the scattered values obtained in non-optimized devices . These results prove that the optimized PMMA solution unlock the production of reproducible electronic devices at the batch scale, which is the key to industrial production. Among devices . Various devices –17. Amon devices . The che devices –22. The 1–2% HCl . Concern1–2% HCl . Anneali1–2% HCl , enablin1–2% HCl ,22–23. S1–2% HCl . PMMA wi1–2% HCl . It mustTherefore, the way toward clean graphene processing appears to lie in the optimization of the PMMA-assisted transfer. In this context, we propose an optimized approach for a clean mass transfer of graphene samples over wafer-scale areas. A PMMA mixture was developed by balancing the AMW and weight percentage in anisole to guarantee a reliable transfer at a negligible contamination level, even without any post-treatment at high temperature. The supporting layer formed by spin coating presents high mechanical flexibility and strength for the transfer process and appears easy to dissolve afterward. We validated the impact of the optimized process in the mass fabrication of arrays of receded-gate graphene field-effect transistors for biosensing applications.2 areas were compared (We transferred graphene by using PMMA with two AMWs (15k and 550k), which were dissolved in anisole at two weight ratios (2 and 4 wt %). PMMA with 950,000 (950k) AMW (at 4 wt %), commonly used for microfabrication processes as an e-beam resist, was used for further comparison . Optical microscopy analysis was carried out to visually evaluate the presence of PMMA residues after the transfer process of graphene single crystals using PMMA with different weight percentages and AMWs –g. As decompared –f. The a2) revealed less than ten residues, indicating an extremely clean transfer process. B2 PMMA allowed for up to six transfer cycles, representing an intermediate, yet acceptable, mechanical support. This proves that the PMMA mixture features good mechanical strength and cleanness . The transferred graphene samples were investigated via Raman spectroscopy to evaluate crystallinity, layer number, and structural defect level [−1) and 2D (ca. 2700 cm−1) bands are typical of monolayer graphene [−1 (I(D)/I(G) ratio is very low (down to 0.03), meaning that no or few defects could be detected. I(2D)/I(G) > 1.6 (average of 2.1 ± 0.3) and a FWHM(2D) of 34.2 ± 3.0 cm−1  re3.0 cm−1 ,i, implyE2g symmetry) at the Brillouin zone center, which is an in-plane vibration of sp2 carbon atoms [EF|, which sets the carrier concentration. Due to the method and materials employed for the graphene transfer being the same except for the PMMA mixture, we consider that the differences in Raman spectra between the different samples can be attributed to adsorbed PMMA residues. Such residues could absorb water and oxygen molecules, conferring p-type doping to graphene [−1), indicating that the charge carrier concentration did not rise as much when using B2 PMMA. This shows that B2 leaves behind a much lower (if any) density of residues. This redshift observed upon reducing the PMMA residue concentration is consistent with studies on advanced methods for cleaning PMMA from graphene [−1, respectively (F|), further indicating that the p-type doping caused by adsorbed PMMA is less intense for the B2 samples. This result again supports that employing B2 PMMA yields fewer residues and may help in avoiding post-transfer treatments for advanced PMMA residue cleaning of graphene, such as annealing and ion beam irradiation [Ra) of ca. 0.2 nm , and O–C=O  bonds, respectively [3-hybridized C–C/C–H bonds in graphene. In 3, and O–C=O bonds are, respectively, 17.6%, 14.6%, and 15.0% of the main peak intensity (sp2 C–C). In 3 bonds, implying the crucial role of PMMA-15k in the mixture.XPS was employed to analyze the graphene samples transferred using C4 and B2 PMMA . Figure ectively –36. PMMAThe optimized graphene transfer process was statistically validated in a batch fabrication context by comparing the channel resistance of electrolyte-gated graphene field-effect transistors GFETs,  designed2 area), while the optimized mixture  yield a minimum value of ten residues in the same area. More abundant PMMA contamination on graphene translated into a more intense p-type doping, as evidenced by (i) the position of the Raman G peak, which was blueshifted by ca. 6 cm−1 (between C4 and B2 samples), and (ii) the FWHM of the G peak, which appeared broader in B2 samples (23 cm−1 vs 14 cm−1 of the C4 samples). The XPS analysis showed a markedly increased presence of C–H, C–C, O–CH3, and O–C=O bonds in C4 graphene compared to B2 samples, corroborating the previous findings. We validated the impact of the optimized B2 process in the mass fabrication of arrays of electrolyte-gated GFETs. The channel resistances of thousands of GFETs prepared using B2 and C4 PMMA were measured with a probe station in air. The resistance distributions were analyzed and fitted with Gaussian curves. The resistance distributions were centered at 690 Ω (FWHM = 303 Ω) and 1860 Ω (FWHM = 567 Ω), respectively, proving that the optimized PMMA mixture enables the production of reproducible arrays of electronic devices with consistent properties.Monolayer graphene films and single crystals were transferred using PMMA with different AMWs and weight percentages in anisole. Repeated transfer cycles among water baths revealed, as expected, that PMMA with higher AMW and weight percentage allowed for a better mechanical support to graphene. Optical microscopy, Raman spectroscopy, and XPS carried out to evaluate the amount of PMMA residues on graphene after the transfer processes showed that PMMA with higher AMW resulted in a more significant number of residues: PMMA-950k AMW (C4) yielded a maximum of 180 residues (in a 650 × 500 µm2) was 40 and 80 min, respectively. To finalize the process, an argon flush of 500 sccm was conducted to cool down samples until the furnace reached room temperature.Single-crystal and large-area graphene were obtained on Cu foil via catalyst-assisted growth in a low-pressure CVD system . A 25 µm thick annealed Cu foil , serving as a metal catalyst, was placed in a graphite enclosed cavity during the whole process. The temperature for annealing and growth was kept stable at 1040 °C by PID thermal controllers. The Cu foil was first annealed in argon atmosphere  for 30 min in a quartz tube furnace. In the growth process, the gas mixture of argon (250 sccm), hydrogen (100 sccm), and methane (1.2 sccm) was subsequently introduced into the quartz chamber, where a reaction pressure of 4.0 Torr was kept constant through the variable frequency-driven pumping system. The growth time for single graphene crystals (250–350 µm) and large-area graphene films .For the preparation of the various PMMA solutions, PMMA-15k  and PMMA-550k  powders were dissolved in anisole  in different proportions, as indicated in 3 solution surface for 3 h to etch away the Cu foil. After the etching process, PMMA/graphene floating on the surface of the FeCl3 solution was rinsed with DI water three times. A rinse in 2% HCl solution was done to remove metal precipitates. At last, the PMMA/graphene was washed in DI water three times and scooped up with a target substrate (SiO2/Si wafer). The sample was dried in a vacuum chamber (ca. 10−4 Torr) at room temperature for 2 h. For PMMA removal, the entire sample was vertically dipped into an acetone bath for 4 h. After that, the exposed graphene on a SiO2/Si substrate was again vertically dipped into IPA and then DI water bath for 1 h. Finally, the graphene on the receiving substrate was blow-dried with N2.PMMA/graphene/Cu foil was placed over the 0.5 M FeClA selective oxidation method was adopted to rapidly identify the as-grown graphene, enabling the direct optical inspection of the graphene domains without the laborious transfer process. Following this method, the Cu substrate with graphene was first oxidized in ambient air on a hot plate at 200 °C for 2 min. The graphene film on the Cu substrate serves as a protection layer, preventing the underlying Cu surface from oxidation because of its high chemical/thermal stability and impermeability to gases and liquids. In contrast, the surrounding exposed areas of the Cu foil surface exhibited high reactivity and were readily oxidized to copper oxides with a noticeable color change. The apparent color contrast between the oxidized and non-oxidized Cu surfaces made the synthesized graphene domains easy to be observed in an optical microscope equipped with a CCD camera  with an excitation laser wavelength of 532 nm. A laser power of ca. 2 mW was used for all measurements. The backscattered laser light containing the Raman bands from 1200 to 3000 cm−1 was collected by a CCD camera  integrated with the WITec system. The characteristic Raman signature collected from a p-doped Si wafer at 520.7 cm−1 was employed as standard calibration. Raman mapping was conducted by raster scan, where the step size of the laser spot moving over a selected area is 1 µm, and the exposure time of 0.4 s was taken at each point of the mapping. In the maps, the intensity of D and 2D bands were normalized to the G band intensity. Corresponding statistics were extracted from the maps.Large-area graphene films and single graphene crystals transferred onto SiO2. The detection system contains a double-focusing 180° spherical sector analyzer with a mean radius of 150 mm and an energy range of 0 to 5 keV. The pressure in the analysis chamber was ca. 5 × 10−10 Torr, and the analyzer had a pass energy of 20 eV.The chemical components of PMMA residues were analyzed by X-ray photoelectron spectroscopy  using a non-monochromatic Mg Kα source with an analysis spot smaller than 2 mmThe surface topographies of graphene were investigated by a Bruker Dimension Icon atomic force microscope (AFM), using PPP-NCH (NanosensorsTM) cantilevers with a tip radius smaller than 20 nm, a force constant of 42 N/m, and 250 kHz resonance frequency. The AFM measurement was carried out in tapping mode. A 633 nm laser light aimed at the back side of the cantilever tip was reflected toward a position-sensitive photodetector, which provides feedback signals to piezoelectric scanners that maintain the cantilever tip at constant height (force) above the surface, thus, reproducing its topography.2 (200 nm thick) wafers. Two arrays of devices were fabricated with different process steps . Both wafers started with the patterning of Cr/Au contacts (deposited by magnetron sputtering) using direct-write laser lithography and ion milling. The fabrication of the two wafers followed slightly different steps, as described below.Receded-gate graphene field-effect transistors were fabricated on an 8″ Si/SiOWafer 1: A stopping layer (Al2O3/TiWN/AlSiCu/TiWN) was patterned by lift-off, followed by the CVD growth of a multi-stack layer of SiO2 and Si3N4 to passivate the current lines. After this, a thin Al2O3 layer was deposited by sputtering and patterned by wet etching to protect the gate during the graphene etch. The C4 PMMA/graphene films were then transferred onto the patterned wafer until all device areas were covered. After removing the PMMA, graphene was patterned using optical lithography and oxygen plasma etching. Finally, the sacrificial layer was removed by wet etching.Wafer 2: An additional layer of Al2O3 was deposited on the Au layer as protection. After that, the two wafers followed different fabrication processes. A residue-free transfer process was used, using a sacrificial metallic mask  patterned by lift-off to protect the entire wafer except for the areas around the channel, source, and drain, on which the graphene film would make electrical contact. The B2 PMMA/graphene films were then transferred onto the wafer and patterned by O2 plasma, followed by the sacrificial layer removal. Previous to the passivation, Al2O3 was selectively removed to improve the adhesion of the oxide passivation to the surface of the chips. A stopping layer (Cu/AlSiCu/TiW) for the reactive ion etching (RIE) process was sputtered, and the SiO2/SiNx multistack passivation layer was deposited by CVD. The passivation layer was patterned by lithography and etched by RIE until revealing the stopping layer on the contact pads and graphene transistor channels. Finally, the stopping layer was removed by wet etching, exposing the graphene channel.W = 75 µm and L = 25 µm. The total number of GFETs condidered in this study was 1755 on Wafer 1 and 4200 on Wafer 2. The source–drain resistance of the GFETs was measured with a semi-automatic DC measurement probe station for 8″ wafers. The setup uses a 40-tip probe head with 250 µm spacing between the tips; each probe allowed us to measure 20 devices simultaneously. A current was injected into each transistor to reach an output voltage of 1 mV. Histogram distributions for the graphene channel resistances were plotted (using MATLAB scripts) to compare the effects of B2 and C4 PMMA-assisted graphene transfer on device performance. Based on our observations of the device performance, a threshold resistance of 15000 Ω was selected as the cut-off resistance value identifying working graphene channels . A maximum resistance value of 1 MΩ was set to identify the non-operating channels (due to imperfect transfer leading to heavily damaged or lacking graphene). Multiple Gaussian fits were performed to isolate the trends in the resistance distributions.In both wafers, the GFET channels had nominal dimensions of File 1Additional experimental data.


Person-centered palliative care (PC) fosters patient and care partner participation during medical appointments. The current study used Bandura’s Social Cognitive Theory to examine whether the manner in which healthcare providers respond to patient and care partner questions during initial PC appointments is associated with their self-efficacy in question asking. We analyzed provider responses to direct and embedded questions during initial outpatient PC appointments at a large academic medical center. Provider responses were coded in terms of timeliness  and completeness (how fully providers answered questions). Provider responses to direct questions were more timely  compared to responses to embedded questions , t(127) = -3.55, p < 0.001, whereas completeness was similar for both direct  and embedded questions , t(127) = -1.59, p = 0.114. On average, receiving more timely responses was not significantly correlated with self-efficacy, r(49) = -0.038, p = 0.791, nor was receiving more complete responses ρ(49) = -0.178, p = 0.211. These findings suggest that timeliness and completeness of responses may not be factors that affect self-efficacy, thus future research should continue to explore other factors, such as empathy present in provider responses. These findings also suggest that future research should investigate the role and effectiveness of embedded questions as an information-seeking behavior in patients and care partners given that providers took longer to respond to them.