Crack-free transfer of CVD graphene

CVD derived single-layer graphene is well-suited for the fabrication of large-area membranes attributing to the scalability of the CVD process22,23. Post-CVD, graphene needs to be transferred from the catalytic metal foil to a porous support for the membrane fabrication. However, the conventional transfer techniques invariably introduce cracks and tears in the graphene film24, and therefore, so far the suspended, crack- and tear-free, single-layer graphene membranes have been limited to a few μm2 in area14,17,25. Among several transfer techniques developed so far, the wet-transfer technique has been investigated the most attributing to its versatility allowing transfer of graphene on a wide-range of substrates26,27,28,29. Briefly, the exposed surface of graphene lying on a metal foil is coated with a sacrificial mechanically reinforcing polymer film (typically 100–200 nm thick poly(methyl methacrylate) (PMMA) film). Subsequently, the metal foil is removed by etching the metal in an etchant bath leaving the polymer coated graphene floating on the bath. Finally, the floating film is scooped on top of the desired substrate, and the polymer film is dissolved away to expose the surface of graphene. This wet-transfer process has been proven to be quite successful in fabricating graphene-based devices on smooth non-porous substrates24. However, significant cracks and tears develop in the graphene film when a porous support is used, primarily because of a strong capillary force on the suspended graphene film during the solvent drying stage30. This issue can be mitigated if the mechanically reinforcing film is not removed, and yet somehow the graphene surface is exposed. Motivated by this, we developed a nanoporous carbon (NPC) film-assisted transfer method (Fig. 1), where at the end of the graphene transfer, the NPC film is left on top of the graphene film. Briefly, a solution of turanose and polystyrene-co-poly(4-vinyl pyridine) (PS-P4VP) was spin-coated on top of the as-synthesized CVD graphene. The block-copolymer film undergoes phase separation into hydrophobic and hydrophilic domains upon drying31. Subsequently, the film was pyrolyzed at 500 °C in the flow of H 2 /Ar, leading to the formation of the NPC film on top of graphene. Scanning electron microscope (SEM) images of the NPC/graphene film on Cu, and transmission electron microscopy (TEM) images of the transferred NPC/graphene film on a TEM grid revealed that the NPC film was 100 nm thick, and comprised of 20–30 nm sized nanopores (Fig. 2a–c), which should expose at least 50% of the graphene surface. Selected area electron diffraction (SAED) of the composite NPC/graphene film (Fig. 2d), displayed the typical diffraction peaks of a suspended single-layer graphene, representing periodicities of 0.213 and 0.123 nm32. NPC film (Supplementary Note 1, Supplementary Table 1) contributed to the SAED with broad rings, a characteristic of the amorphous structure (Supplementary Figure 1). We could not find any area representing only the NPC film, indicating that graphene and NPC film bonded strongly during the pyrolysis step and that the graphene did not peel off from the NPC film during the etching of the metal foil. This is highly important for crack-free transfer of graphene, otherwise poor interactions of graphene with a support film can lead to severe cracks and tears during the transfer step33. The NPC coated graphene was transferred from the Cu foil to a custom-made macroporous support (porous area of 1 mm2, Fig. 2e, f). The support was fabricated by laser drilling an array of 5 μm pores in a 50-μm-thick W foil25 (Supplementary Note 2). Inspection of the transferred film by the SEM confirmed that there were no visible tears or cracks in the transferred film (Fig. 2g). Interestingly, even a macroscopic fold as shown in Fig. 2g did not break the membrane, making this process highly promising for the scale-up of single-layer graphene membrane.

Fig. 1 Schematic of fabrication of large-area graphene membrane by the nanoporous carbon (NPC) film-assisted transfer method. A block-copolymer solution was spin-coated onto the CVD graphene supported on the Cu foil; pyrolysis was conducted to form NPC film on top of graphene. The copper was etched by 0.2 M sodium persulfate, after which the floating graphene/NPC film was rinsed with DI water. Finally, the NPC/graphene film was transferred onto the porous tungsten support Full size image

Fig. 2 Synthesis, transfer, and characterization of low-pressure chemical-vapor deposition (LPCVD) derived graphene. a Scanning electron microscopy (SEM) image of the nanoporous carbon (NPC) film coated on top of graphene. b Cross-sectional SEM image of the composite NPC film and graphene. c Transmission electron microscopy (TEM) image of the composite NPC film/graphene. d The electron diffraction pattern from the composite film shown in c. e SEM image of porous tungsten support. f SEM image of porous tungsten support. g SEM image of the transferred graphene on the tungsten support. h Histogram of I D /I G from LPCVD graphene. i–k High-resolution TEM (HRTEM) images of the intrinsic defects in graphene lattice. The unprocessed raw images are shown in Supplementary Figure 3a-c. Scale bars in a, b, and c are 200, 100, and 50 nm, respectively. Scale bar in d is 10 nm−1. Scale bars in e, f, and g are 500, 20, and 500 µm, respectively. Scale bars in i, j, and k are 1 nm Full size image

Gas transport through intrinsic defects of graphene

Using scanning tunneling microscope (STM) directly on the as-synthesized graphene supported on the Cu foil, we recently imaged the low-density of intrinsic defects in the CVD graphene25. These defects are essentially molecular-sized pores (missing 10–16 carbon atoms), formed by etching of the graphene lattice in the presence of residual oxygen in the CVD chamber, and are promising for the gas separation. In the current study, the density of defects, estimated by the carbon amorphization trajectory34 (I D /I G of 0.07 ± 0.02, Fig. 2h, Supplementary Figure 4), was 5.4 × 1010 cm−2 corresponding to a porosity of 0.025%. A survey of the graphene lattice by aberration-corrected high-resolution TEM (HRTEM) revealed several sub-1-nm nanopores with a pore-density of 2.8 × 1011 cm−2 (Fig. 2i–k). This small disagreement between the HRTEM survey and the estimate from the amorphization trajectory is expected especially at a low defect density. Nevertheless, the successful crack-free transfer of the CVD graphene allowed us to study the transport behavior of the intrinsic defects.

Graphene membranes were sealed in a homemade permeation cell using a metal face seal directly on top of the W support, ensuring a leak-proof measurement of the gas transport (details in the Methods). Typically, the feed side (a pure gas feed or a mixture feed) was pressurized to 1.5–7.0 bar, whereas the permeate side connected to a pre-calibrated mass spectrometer (MS) was maintained at 1 bar with an argon sweep (Supplementary Figure 2). Temperature of the membrane was varied between 25–250 °C. Single-component gas transport study from eight separate membranes revealed H 2 permeance in the range of 5.2 × 10−9–7.2 × 10−8 mol m−2 s−1 Pa−1 (15–215 gas permeation units, GPU) with H 2 /CH 4 , H 2 /CO 2 , and He/H 2 ideal selectivities ranging between 4.8–13.0, 3.1–7.2, and 0.7–2.0, respectively, at 25 °C (Fig. 3a–d, Supplementary Table 9–11). The H 2 permeance corresponded to a permeation coefficient of 1.0 × 10−23–1.3 × 10−22 mol s−1 Pa−1 based on the defect density of 5.4 × 1010 cm−2. This permeation coefficient is consistent with that of Bi-3.4 Å membrane reported by Koenig et al. (4.5 × 10−23 mol s−1 Pa−1)20. The H 2 /CH 4 selectivity was lower than that from Bi-3.4 membrane20, indicating a wider PSD of intrinsic defects in CVD graphene, compared to PSD from pores incorporated in micromechanically exfoliated graphene. Based on the achieved H 2 /CH 4 selectivities, the estimated percentage of larger nanopores yielding non-selective effusive gas transport is less than 25 ppm (refer to the Supplementary Note 4 and Supplementary Table 8 for more details). Interestingly, the H 2 /CO 2 selectivity was higher than that of the Bi-3.4 membrane where a selectivity of ca. 1.5 was reported. Membrane M8 displayed the best molecular sieving performance and was the only membrane displaying He/H 2 selectivity greater than 1, implying that the mean pore-size in M8 was less than the kinetic diameter of H 2 (0.289 nm).

Fig. 3 Gas separation performance of the intrinsic defects in graphene. a H 2 permeance across eight graphene membranes (M1–M8) as a function of temperature when using a single-component feed. b–d Ideal selectivities for various gas pairs from the eight membranes as a function of temperature; b H 2 /CH 4 , c H 2 /CO 2 , and (d) He/H 2 . e Extracted activation energies (average across all eight membranes) are plotted as a function of the kinetic diameters of various gases Full size image

The graphene membranes did not rupture during heating to up to 250 °C. The permeance of He, H 2 , CO 2 and CH 4 increased with temperature, indicating that its transport was in the activated transport regime. At 150 °C, the H 2 permeance increased to 3.3 × 10−8 – 4.1 × 10−7 mol m−2 s−1 Pa−1 (100 - 1220 GPU), with H 2 /CH 4 , and H 2 /CO 2 selectivities increasing to 7.1–23.5 and 3.6–12.2, respectively. We note that this H 2 /CH 4 separation performance from single-layer graphene with 0.025% porosity approaches the 2008 Robeson upper bound for polymers35 (assuming 1-µm-thick selective skin layer of the polymer membrane, Supplementary Figure 6). To understand the transport behavior, the activation energy for gas diffusion across the nanopores was extracted from the temperature-dependent gas flux using an adsorbed phase transport model developed using the concepts of adsorption and diffusion36,37,38,39 (Eq (1), refer to Supplementary Note 3 for details).

$${\mathrm{Flux}} = C_{\mathrm{o}}A_{\mathrm{act}}A_{\mathrm{sur}}{\mathrm{exp}}\left( { - \frac{{(E_{{\mathrm{act}}} + \Delta E_{\mathrm{sur}})}}{{RT}}} \right)\left( {f\left( {P_{\mathrm{A}}} \right) - f\left( {P_{\mathrm{R}}} \right)} \right)$$ (1)

where

$$f\left( {P_x} \right) = \frac{{P_x}}{{1 + A_{\mathrm{sur}}{\mathrm{exp}}\left( {\frac{{ - \Delta E_{\mathrm{sur}}}}{{RT}}} \right)P_x}}$$

Here, C O is the pore-density, E act is the activation energy for the gas diffusion across the nanopores, and ΔE sur is the adsorption energy of gas on to the graphene nanopore. A act is the pre-exponential factor for the gas diffusion across the nanopores. A sur is the pre-exponential factor for the adsorption event, representing changes in the overall entropy. T is the temperature, and P A and P R are the gas partial pressures on the feed and permeate sides, respectively. A comparison of E act for the four gases can indicate the ease with which the molecules diffuse across the nanopores, while a comparison of the pre-exponential factor, C o A act A sur , can indicate the relative number of pores participating in the molecular diffusion. Average E act across eight membranes for He, H 2 , CO 2 , and CH 4 were 14.7 ± 3.2, 20.2 ± 2.7, 31.3 ± 2.8, and 25.8 ± 4.8 kJ/mol, respectively, increasing as a function of kinetic diameter (Fig. 3e, Supplementary Table 2). A slightly smaller E act for CH 4 in comparison to CO 2 can be explained by the fact that diffusion of CH 4 takes place from a smaller number of pores (average C o A act A sur for He, H 2 , CO 2 , and CH 4 were 1.5 × 10−5, 2.6 × 10−5, 3.8 × 10−6, and 1.3 × 10−6, respectively, Supplementary Table 3), assuming A act A sur do not change significantly for CO 2 and CH 4 37. The activation energy for H 2 was similar to that from hydrogen-functionalized pore-10 reported by Jiang et al. (0.22 eV)3, providing an indication that the average pore in this study is close to that made from missing 10 carbon atoms, which is supported by the HRTEM images (Fig. 2i–k) and our previous STM findings25. We note that while the gas permeance varied across the eight membranes, the activation energies for diffusion across the nanopores were consistent across the membranes. This indicates that, while the PSD was uniform across the membranes, the density of the intrinsic defects varied across the membranes.

Separation of a gas mixture can elucidate the contribution of competitive adsorption on the overall separation performance from the nanoporous graphene membrane. To the best of our knowledge, gas mixture separation through a single-layer graphene membrane has not been reported. Here, the large-area of the graphene membrane enabled measurements of He, H 2 , CO 2 , and CH 4 permeances from an equimolar gas mixture. The competitive adsorption3,9,36,40 is expected to yield a reduced He/H 2 , H 2 /CO 2 , and H 2 /CH 4 separation factor (SF) compared to the corresponding ideal selectivities (IS). However, the separation factors were similar (He/H 2 and H 2 /CO 2 ) or higher (H 2 /CH 4 ) compared to the corresponding ideal selectivities (Fig. 4a–d). For example, for membrane M2, the H 2 /CH 4 SF was higher than IS (10.8 vs. 5.7 at 25 °C and 12.2 vs. 11.2 at 150 °C), while the H 2 permeance (3.3 × 10−8–2.2 × 10−7 mol m−2 s−1 Pa−1 between 25 and 150 °C) in the mixture case was similar to the single-component case (Fig. 4a). Similarly, the H 2 permeance did not reduce for the membrane M3 for the mixture case, while the H 2 /CH 4 SF increased to 18.0 from an IS of 14.2. For other membranes (M1, M4, M5, and M6), the H 2 permeance and the H 2 /CH 4 selectivity in the mixture case were similar to those in the single-component case. These results indicate that the competitive adsorption of gases on the basal plane of graphene does not play a significant role in overall transport especially when the transport is in the activated regime, and when the feed pressure is moderate (up to 8 bar in this study). We expect the competitive adsorption to play a role at a higher feed pressure (30–50 bar, refer to Supplementary Note 3 and Supplementary Equation 36), which will be investigated in future studies.

Fig. 4 The gas mixture separation performance of the intrinsic defects in graphene. a H 2 permeance from six membranes (M1–M6) as a function of temperature when using an equimolar mixture gas feed. H 2 /CH 4 , H 2 /CO 2 , and He/H 2 separation factors as a function of temperature are shown in b, c, and d, respectively Full size image

The graphene membranes were thermally stable (Fig. 5a). In general, all membranes were stable at least up to 150 °C. For instance, the performance of membrane M2, tested under three consecutive temperature cycles from 25 °C to 150 °C, did not change significantly. From cycle one to cycle three at 150 °C, the H 2 permeance decreased marginally (3.3 × 10−7 to 2.3 × 10−7 mol m−2 s−1 Pa−1), while the H 2 /CH 4 selectivity increased marginally (8.3 to 10.5). Moreover, the graphene membranes were also stable at least up to 8 bar of mixture feed at 100 °C (permeate pressure 1 bar, Fig. 5b, c), where the H 2 permeance and the H 2 /CH 4 separation factor did not change significantly.

Fig. 5 Stability test of the graphene membrane. a H 2 and CH 4 permeances of the membrane M2 with three consecutive temperature cycles. Gas permeance (b) and H 2 /CH 4 separation factor (c) from the membrane M2 as a function of transmembrane pressure difference at 100 °C Full size image

Ozone functionalization-based etching and pore-modification chemistry

The porosity of graphene lattice yielding the attractive H 2 permeance was only 0.025%. Theoretically, the H 2 permeance can be further increased beyond 105 GPU by increasing the defect density to 1012–1013 cm−2. On the other hand, the gas selectivity can be improved by constricting the nanopores. One way to achieve this is chemical functionalization of the pore-edge. Although, there are several potential chemical and physical routes to open pores in graphene, development of an in-situ etching method (inside membrane module) allowing a high degree of control is highly attractive. In this pursuit, we report a scalable ozone functionalization-based pore-etching and pore-edge-functionalization chemistry, improving the performance of the single-layer graphene membranes. We demonstrate that a controlled temperature-dependent oxidative functionalization of the graphene lattice with ozone-derived epoxy and carbonyl groups can be used to either etch molecular-sized pores in the CVD derived graphene or constrict the existing pores.

Oxidative treatment of graphene has been shown to incorporate sp3-hybridized sites (epoxy and carbonyl groups) on the basal plane of graphene41,42,43. When the functionalization density is high, such as that in graphene oxide, one can introduce nanopores in the lattice by thermal annealing44,45,46,47. Typically, the functional groups migrate and rearrange forming larger groups (such as lactone), and finally desorb as CO or CO 2 leading to a vacancy44. Ozone, in the gas phase, can be conveniently used to oxidize graphene lattice. To understand the evolution of functionalization, CVD graphene supported on a Cu foil was exposed to ozone at various temperatures (25 °C, 80 °C, and 100 °C) and time (1 min to 7 min). The evolution of oxidative groups on graphene was probed by micro-Raman and X-ray photoelectron spectroscopy (XPS). The relative intensity of the D peak with respect to the G peak (I D /I G ), which marks the extent of disorder in graphene48, increased from 0.07 to 4.0, while the intensity of the 2D peak decreased in intensity with the increasing reaction time and temperature, indicating that the sp3-hybridized sites in graphene increased after ozone treatment (Fig. 6a, b). XPS indicated that C–O and C=O were the major functional groups on the functionalized graphene (Fig. 6c, d and Supplementary Figure 5). The number density of functional groups increased with the reaction temperature and time, in agreement with the Raman spectroscopy. In general, the density of C=O groups was higher than that of the C–O groups, even when the functionalization was carried out at room temperature for a short exposure of 2 min. At 100 °C, the degree of oxidation approached that of graphene oxide (35, 56, and 65% of the oxidized carbon lattice with exposure times of 2, 5, and 7 min, respectively). Overall, the functionalization was reproducible and was simple to implement. HRTEM images of the ozone-functionalized graphene (80 °C for 2 min) indeed revealed a higher pore-density (4.2 × 1011 cm−2) compared to that in as-synthesized graphene (Fig. 6e–g). Moreover, the population of the sub-1-nm pores (87%) increased compared to that in the as-synthesized graphene (76%) (details in Supplementary Note 5).

Fig. 6 Characterization of the ozone-treated graphene. a Raman spectra of functionalized graphene under different functionalization conditions. b I D /I G for various ozone treatment. c, d C–O (c) and C=O (d) content of the graphene as a function of the functionalization time and temperature. e–g High-resolution transmission electron microscopy (HRTEM) images of the nanopores in the ozone functionalized graphene (2 min at 80 °C). The scale bar is 1 nm. The unprocessed raw images are shown in Supplementary Figure 3d-e Full size image

To understand the effect of functionalization on the performance of graphene membrane, the graphene membranes were exposed to ozone, in-situ, in the permeation setup after probing the gas transport from the intrinsic defects. With this strategy, the gas transport before and after the functionalization could be compared. Overall, the separation performance of the graphene membranes improved, reflected by an increase in the H 2 permeance or an increase in the H 2 /CH 4 selectivity or an increase in the permeance as well as the selectivity (Figs. 7 and 8). When functionalization was carried out at 25 °C for 2 min, the H 2 permeance decreased from 2.3 × 10−7 to 1.2 × 10−7 mol m−2 s−1 Pa−1, while the H 2 /CH 4 and the H 2 /CO 2 selectivities increased from 10.0 to 15.0 and 5.1 to 6.4, respectively, at 150 °C (M2, Fig. 7a, b), indicating pore-shrinkage. Interestingly, both E act-app (defined as E act +ΔE sur ) and C o A act A sur decreased after the functionalization (Supplementary Table 4 and Supplementary Note 3). While the changes in E act - app are complex to interpret because of relative changes in E act (higher activation energy due to the pore-shrinkage) and ΔE sur (increase in binding energy with functionalized pores), a 20-fold decrease in C o A act A sur for CH 4 (5.7 × 10−7 to 2.8 × 10−8, Supplementary Table 5) as a result of functionalization indicates reduced pore-density for the diffusion of CH 4 . We envision the functionalized pore-edges would shrink in size, providing higher resistance to CH 4 for the diffusion, and therefore resulting in a higher gas selectivity25. In contrast, functionalization at 100 °C led to an increase in gas permeance by 3-fold, while the gas selectivity change slightly (Fig. 7c, d). Here, E act-app did not change significantly after functionalization (Supplementary Table 6), while C o A act A sur for gases increased by an order magnitude (Supplementary Table 7) indicating an increase in the pore-density. Given that the high-temperature functionalization leads to a higher coverage of the C–O and C=O groups, it is highly likely that these functional groups formed new pores as indicated by the HRTEM analysis.

Fig. 7 Gas separation performance of ozone-treated graphene membranes. a, b Gas separation performance of M2 treated by 2 min O 3 at 25 °C, a gas permeance of H 2 and CH 4 , b gas selectivity of H 2 /CH 4 and H 2 /CO 2 . c, d Gas separation performance of M8 treated by 2 min O 3 at 100 °C, c gas permeance of H 2 and CH 4 , d gas selectivity of H 2 /CH 4 , and H 2 /CO 2 . e, f Gas separation performance of M6 treated by 1 min O 3 at 80 °C, e gas permeance of H 2 and CH 4 , f gas selectivity of H 2 /CH 4 and H 2 /CO 2 Full size image

We constructed a separation performance trajectory (Fig. 8), comparing the separation selectivity and hydrogen permeance before and after the ozone functionalization (Fig. 8). The overall trajectory trends clearly show that the gas separation performance of graphene membranes can be tuned by the ozone functionalization. A higher gas permeance (up to 300%) was achieved by generating new nanopore by ozone functionalization at 80–100 °C (membranes M7 and M8, Supplementary Figure 8 and Fig. 7c, d). A higher separation selectivity (up to 150%) was achieved by functionalization at 25 °C (M2, Fig. 7a, b). In one case, an increase in permeance, as well as separation selectivity, was obtained after ozone treatment at 80 °C for 1 min (M5, Supplementary Figure 7). Therefore, one can use ozone functionalization as a post-synthetic performance tuning method to enhance the separation performance of nanoporous single-layer graphene membranes.

Fig. 8 The evolution of gas separation performance after different ozone treatments. Permeance data for membrane M8 is at 200 °C, and all other (M2, M5, M6, and M7) is at 150 °C. Light and dark markers represent the gas performance from the intrinsic defects and the ozone-treated graphene, respectively Full size image