Scholarly article on topic 'Graphene-based materials: Synthesis and gas sorption, storage and separation'

Graphene-based materials: Synthesis and gas sorption, storage and separation Academic research paper on "Nano-technology"

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Progress in Materials Science
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{Graphene / "Graphene oxide" / Synthesis / "Hydrogen storage" / "Methane storage" / "Carbon capture" / "Gas storage" / "Gas separation"}

Abstract of research paper on Nano-technology, author of scientific article — Srinivas Gadipelli, Zheng Xiao Guo

Abstract Graphene-based materials have generated tremendous interest in a wide range of research activities. A wide variety of graphene related materials have been synthesised for potential applications in electronics, energy storage, catalysis, and gas sorption, storage, separation and sensing. Recently, gas sorption, storage and separation in porous nanocarbons and metal–organic frameworks have received increasing attention. In particular, the tuneable porosity, surface area and functionality of the lightweight and stable graphene-based materials open up great scope for those applications. Such structural features can be achieved by the design and control of the synthesis routes. Here, we highlight recent progresses and challenges in the syntheses of graphene-based materials with hierarchical pore structures, tuneable high surface area, chemical doping and surface functionalization for gas (H2, CH4, CO2, N2, NH3, NO2, H2S, SO2, etc.) sorption, storage and separation.

Academic research paper on topic "Graphene-based materials: Synthesis and gas sorption, storage and separation"

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Graphene-based materials: Synthesis and gas sorption, storage and separation

Srinivas Gadipelli *, Zheng Xiao Guo *

Department of Chemistry, University College London, 20 Gordon Street, London WC1H 0AJ, United Kingdom


Graphene-based materials have generated tremendous interest in a wide range of research activities. A wide variety of graphene related materials have been synthesised for potential applications in electronics, energy storage, catalysis, and gas sorption, storage, separation and sensing. Recently, gas sorption, storage and separation in porous nanocarbons and metal-organic frameworks have received increasing attention. In particular, the tuneable porosity, surface area and functionality of the lightweight and stable graph-ene-based materials open up great scope for those applications. Such structural features can be achieved by the design and control of the synthesis routes. Here, we highlight recent progresses and challenges in the syntheses of graphene-based materials with hierarchical pore structures, tuneable high surface area, chemical doping and surface functionalization for gas (H2, CH4, CO2, N2, NH3, NO2, H2S, SO2, etc.) sorption, storage and separation.

© 2014 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY license (



Article history: Received 8 September 2014 Accepted 9 October 2014 Available online 29 October 2014

Keywords: Graphene Graphene oxide Synthesis Hydrogen storage Methane storage Carbon capture Gas storage Gas separation


1. Introduction....................................................................................................................................................2

2. Gas sorption, storage and separation: principles and methodology............................................................5

3. Theoretical insights and predictions..............................................................................................................6

3.1. Hydrogen storage................................................................6

3.1.1. Expanded and pillared graphene layers..............................................................................6

* Corresponding authors. E-mail addresses: (S. Gadipelli), (Z.X. Guo). 0079-6425/© 2014 The Authors. Published by Elsevier Ltd.

This is an open access article under the CC BY license (

3.1.2. Pillared graphene oxide structures........................................... 7

3.1.3. Doped, functionalised and metal dispersed graphenes........................... 8

3.1.4. Graphane based materials................................................. 11

3.1.5. Hydrogen spillover....................................................... 11

3.2. CH4 and CO2: sorption, storage and separation........................................12

3.3. Gas separation membranes........................................................14

4. Experimental developments............................................................ 16

4.1. Chemically reduced/thermally exfoliated graphenes...................................16

4.2. Layered graphenes: simple expansion to pillaring with metals and organic ligands..........19

4.3. Spillover and dissociative atomic hydrogen sorption...................................25

4.4. Bottom-up solution-phase synthesised graphenes.....................................28

4.5. Highly porous graphene carbons from templating, CVD and KOH chemical activation........29

4.6. Electrochemical hydrogen storage..................................................31

4.7. Chemical hydrogen storage.......................................................31

4.8. Graphene membranes and protection barriers........................................33

4.9. NH3, NO2, H2S and SO2 sorption....................................................39

5. Summary and perspectives.............................................................43

6. Conclusion.......................................................................... 50

Acknowledgement.................................................................... 50

References.......................................................................... 50

1. Introduction

Graphene is a two-dimensional (2D) sp2 bonded carbon sheet, arranged in a hexagonal honeycomb lattice [1-4]. From a fundamental point of view, graphene is nothing but a single layer of graphite, which is an infinite three-dimensional (3D) material made up of stacked layers of graphene. The layers in graphite interact weakly through van der Waals (vdW) forces. In terms of properties graphene is unique; it is a soft membrane and at the same time possesses a high Young's modulus, and good thermal and electrical conductivities [4-6]. In addition, a single-layer graphene is a zero band gap material and highly transparent, exhibits optical transmittance of 97.7%. With its high theoretical specific surface area of ~2600 m2/g graphene provides a rich platform for surface chemistry [7-12]. The combined extraordinary physical and chemical properties of graphene, in turn, has ignited extensive research in nanoelectronics, supercapacitors, fuel-cells, batteries, photovoltaics, catalysis, gas sorption, separation and storage, and sensing [13-28]. A roadmap of graphene materials is described in a recent review [4]. It is important to note that most of the graphene properties are sensitive to structural defects and the number of layers [4,29-31]. Thus in order to exploit most of the proposed applications, the synthesis routes and conditions is important in tuning the structure and properties of graphene. There are a number of reviews on graphene related materials for possible applications in relation to optical, electronic, photocatalytic and electrochemical properties [1-50]. However, there is a lack of consideration on their important molecular interactions, e.g. for molecular adsorption and storage. This article attempts to address such issues.

Several large scale processing methods have been involved for different graphene based materials through either graphitic top-down or molecular carbon precursor bottom-up approaches, as summarised in Fig. 1. The flexibility in modification and functionalization of the graphene surface has opened up many possibilities for the development of tailored functional materials. For example, surface modification has been applied to tune the band-gap of single-layer graphene for microelectronic devices. Similarly, the intrinsic non-porous 2D graphene is tuned to highly porous 3D architectures for electrochemical devises (batteries, fuel-cells and supercapacitors) and gas sorption, storage, separation and sensing.

On a relatively large scale, graphene is mostly obtained from graphite precursors through oxidation-exfoliation-reduction, i.e. in the form of graphene oxide (GO) as schematically shown in Fig. 1a [4,33,49-52]. The GO structure contains abundant oxygen-rich functional groups; hydroxide and epoxide groups on the basal plane and carbonyl and carboxyl groups on edges of the graphene

Fig. 1. (a) Graphitic top-down approach and (b) molecular carbon precursor bottom-up approach for producing wide variety of graphene based materials in large quantities.

sheets. Thus the GO is hydrophilic in nature and soluble in water and several solvents. GO with its lamellar water, a largely expanded and tuneable layer structure provides a rich platform for engineering a wide range of functionalities and reaction sites for chemical modifications. Generally, the oxidation and reduction creates many defective sites on the graphene, which offers clear advantage in gas sorption, storage and separation and further functionalization.

Continued increase in energy demand and the urgency in reducing CO2 emission require rapid development of alternative and clean energy technologies [53-56]. Efficient ways of storing H2 [5663] and capturing CO2 [54,55,64-72] are key challenges in the development of hydrogen fuel-cell vehicles and carbon capture systems, respectively. Among the different methods, gas sorption and storage by physical adsorption in porous media is considered as a promising approach. Here the gas sorption and storage capacity is mainly governed by a high accessible surface area and pore structure. Many efforts have been made to synthesise a wide variety of tailor-made porous materials, such as zeolites, carbons, polymers, metal-organic frameworks (MOFs), and covalent organic frameworks (COFs) as adsorbents for H2 and CH4 storage, and carbon capture and separation from the flue gas [73-82]. The US Department of Energy (DOE) has set certain targets for H2 and CH4 storage materials for their practical applications [83-86]. Currently, the on-board H2 storage targets are: 5.5 wt% (gravimetric) and 40 g/L (volumetric) near ambient temperature. Similarly, a new CH4 storage program has set the following targets: 0.5 g(CH4) per g(sorbent) for gravimetric capacity and 11.741 mmol/cm3 (q = 0.188 g/cm3) for volumetric capacity, which corresponds to the density of compressed natural gas (CNG) at 250 bar and 298 K. The new volumetric target is equal to 263 cm3 (STP: 273.15 K,

Fig. 2. Graphene based materials for gas sorption, storage and separation.

1 bar) per cm3, which is significantly higher than the previous target of 180 cm3 (STP) per cm3 at 35 bar.

Carbon based materials have been considered for promising gas sorption, storage, and separation because of the abundance, robust pore structure, tuneable porosity and surface area, light-weight, high thermal and chemical stability, and easy synthesis in industrial scale. There is a considerable amount of interest in graphene related materials for gas sorption, storage and separation but still a lack of comprehensive review on such a topic [7,34,36,43,87-89]. Here, we present an up-to-date overview of the theoretical predictions and experimental results on the graphene-based materials towards gas sorption, storage, and separation with particular emphasis on H2, CO2, and CH4. In addition, a detailed account is also provided on the adsorptive removal of toxic gas pollutants, such as NH3, NO2, SO2 and H2S. Fig. 2 represents the overall overview of graphene-based materials and their possible applications in molecular adsorption and storage.

2. Gas sorption, storage and separation: principles and methodology

Gas sorption, storage and separation in carbon materials are mainly based on physisorption on the surfaces and particularly depend on the electrostatic and dispersion (i.e., vdW) interactions. The former can be tuned by introducing charge variations in the material, and the latter by chemical substitution. The strength of the interaction is determined by the surface characteristics of the adsorbent and the properties of targeted adsorbate molecule, including but not limited to the size and shape of the adsorbate molecule along with its polarizability, magnetic susceptibility, permanent dipole moment, and quadrupole moment. Li et al. summarise the adsorption-related physical parameters of many gas or vapour adsorbates, and herein Table 1 we show a few of those of interest, H2, N2, CO, CO2, CH4, NH3, SO2 and H2S [90]. For instance, an adsorbent with a high specific surface area is a good candidate for adsorption of a molecule with high polarizability but no polarity. Adsorbents with highly polarised surfaces are good for adsorbate molecules with a high dipole moment. The adsorbents with high electric field gradient surfaces are found to be ideal for the high quadrupole moment adsorbate molecules [91]. Normally, the binding or adsorption strength with a carbon nano-structure is relatively low for H2 and N2; intermediate for CO, CH4 and CO2; and relatively high for H2S, NH3 and H2O. Thus, surface modifications, such as doping, functionalization and improving the pore structure and specific surface area of nanocarbons, are important to enhance gas adsorption. For this purpose, graphene offers a great scope for tailor-made carbonaceous adsorbents.

Experimentally, at a given temperature, the quantity of adsorbed gas can be determined by an adsorption isotherm, generally carried out by one of the two methods: volumetric or gravimetric. The adsorption isotherm (namely equilibrium isotherm) characterises the adsorption equilibrium. The adsorptive isotherms of individual pure gases are also considered as a promising way to evaluate

Table 1

Some of the physical parameters of selected gaseous adsorbates.

Adsorbate Normal TC(K) VC Pc Kinetic Polarizability x 1025 Dipole Quadrupole

BP (K) (cm3/ (bar) diameter (cm3) moment x 1018 moment x 1026

mol) (nm) (esu cm) (esu cm2)

He 4.30 5.19 57.30 2.27 0.2551 2.04956 0 0

H2 20.27 32.98 64.20 12.93 0.2827-0.289 8.042 0 0.662

N2 77.35 126.20 90.10 33.98 0.364-0.380 17.403 0 1.52

O2 90.17 154.58 73.37 50.43 0.3467 15.812 0 0.39

CO 81.66 132.85 93.10 34.94 0.3690 19.5 0.1098 2.50

CO2 216.55 304.12 94.07 73.74 0.33 29.11 0 4.30

NO2 302.22 431.01 - 101.00 - 30.2 0.316 -

SO2 263.13 430.80 122.00 78.84 0.4112 37.2-42.8 1.633 -

H2S 212.84 373.40 98.00 89.63 0.3623 37.82-39.5 0.9783 -

NH3 239.82 405.40 72.47 113.53 0.2900 21.0-28.1 1.4718 -

H2O 373.15 647.14 55.95 220.64 0.2641 14.5 1.8546 -

CH4 111.66 190.56 98.60 45.99 0.3758 25.93 0 0

the adsorptive separation. In case of membrane separation, the kinetic diameter of test gas plays an important role, where gas molecules with a kinetic diameter smaller or larger than the membrane pore diameter are separated by molecular sieving. Molecular separation can also be facilitated by large differences in diffusion kinetics across a relatively thick porous membrane.

3. Theoretical insights and predictions

3.1. Hydrogen storage

3.1.1. Expanded and pillared graphene layers

The planar sheet of graphene with its inherent specific surface area of ~2630 m2/g [92] in the ideal case has further motivated theoretical calculations for gas adsorption from the well-established carbon nanotubes and graphitic structures. The adsorption of a monolayer of H2 on a single side of graphene sheet can lead to about 3 wt% of H2 (H/C = 0.18). However, it is thermodynamically impossible for H2 molecules to penetrate between the graphene layers of graphite. Deng et al. design a Li-doped pillared graphene sheets (Li-PGS) and single-wall carbon nanotubes (Li-P-SWNT) [93]. Grand canonical Monte Carlo (GCMC) simulations predict a H2 storage capacity of 6.5 wt% at 20 bar and room temperature in Li-PGS with a doping ratio of Li:C =1:3 and an interlayer distance of 1.0 nm (Fig. 3). The Li dopants act as positive (acidic) cores to attract H2 molecules with a binding energy of ~0.3 eV. An interlayer spacing of around 0.6 nm with Li:C = 1:6 leads to over 3 wt% of H2 adsorption. This interlayer spacing [94] or pore width [95] allows H2 molecules to interact both sides of graphene. Thus the spatial distribution of molecular hydrogen adsorption on the graphene plane is very delocalised (Fig. 4). The H2 interaction with graphene in bulk graphite is a localised phenomenon therefore adsorption generally restricted to the outermost graphene plane [96]. A maximum of 3.3 wt% of H2 is achieved with one monolayer adsorption between two graphene layers separated by 0.6 nm (Fig. 4). At most two H2 monolayers could fit between the graphene layers when the layers are separated by 0.9 nm, yielding a maximum capacity of 6.6 wt% of H2. Up to 16 kJ/mol binding energy is obtained for an interlayer distance of ~0.6 nm [94,97,98]. For larger pore widths or interlayer spacing, the binding due to the second graphene layer/surface weakens [95]. For smaller separations, the exchange repulsion reduce the free energy and becomes positive for interlayer separations of <0.5 nm. Thus energetically the H2-graphite system at shorter separations is purely repulsive. The gravimetric and volumetric H2 storage capacities in the C60 intercalated graphite (CIG) [99] approach the values obtained for Li-doped graphenes [93] at low temperatures. In another case, a capacity of 3 wt% of H2 is estimated at ambient in ~0.7 nm spaced carbon nanoscrolls followed by alkali doping [100]. First-principles calculations predict 3.4 wt% of H2 and binding energy of (10-22) kJ/mol in the Li and organic molecules (benzene and tetrahydrofuran) co-intercalated graphite with an interlayer

interlayer distance (A) Li: C =1 :x Temperature (K)

Fig. 3. Left: Dependence of H2 storage capacity on the interlayer distances (ILD) under various pressures for (a) PGS (white) and Li-GIC or Li-PGS (black); the doping concentration is Li:C = 1:6. The pressures are square, 50 bar; circle, 10 bar, and triangle, 1 bar. Middle: Effects of Li-doping concentration on the H2 storage capacity under various pressures. ILD = 1.0nm. Right: Temperature and pressure (unit: bar) effects on the H2 storage capacity in Li-PGS: Li:C = 1:3 and ILD = 1.0 nm. The horizontal line represents the US Department of Energy (DOE) target [93].

Interlayer separation, A interlayer spacing (À)

Fig. 4. Gravimetric (a) and volumetric (b) H2 storage capacities of layered graphite structures, calculated from the real gas equation of state, as a function of interlayer separation. The DOE targets (w = 6.5%, v =31.2 cm3/mol) are indicated by solid horizontal lines. (c) Storing H2 between rationally spaced graphene sheets. The figure shows probability densities for selected lowest eigenstates of the translational nuclear Hamiltonian. The lowest in-phase eigenstates for the double-layer structure are shown (interlayer distance of 0.8 nm) [94]. (d) H2 absorption as a function of graphite interlayer spacing at 298 K for 10 bar (circles) and 50 bar (squares) pressures [95].

graphene distance of ~0.77 nm [101]. Furthermore, the charged graphenes show enhanced H2 binding energy up to kJ/mol compared to ~5.8 kJ/mol in a pristine graphene. Similar binding energies are also achieved in Li-doped single-walled carbon nanotube (CNT) pillared graphene structures with intertube and interlayer distance of 1.5 nm and 1.2 nm, respectively [102]. The molecular dynamics simulations estimate H2 adsorption of up to 6 wt% at 77 K and 3 wt% at 300 K in the CNT pillared graphenes [103].

Furthermore, a detailed analysis and the theoretical limits of H2 adsorption and storage capacity of slit pores is also presented [104-106]. For example, practicable H2 adsorption capacity at room temperature is possible for graphene carbons with a slit-pore size between (0.8-1.1) nm, adsorption energy of p15 kJ/mol and a specific surface area of ~2600 m2/g. The porous materials with micropores smaller than 0.7 nm are not useful for storage application because of low delivery efficiency and large pores (>1.2 nm) have difficulty in fulfilling the volumetric goals at room temperature. The optimal volumetric storage is expected when the pore volume accommodate 1 or 2 layers of H2.

3.1.2. Pillared graphene oxide structures

The loosely stacked layers in GO similar to graphite but with a much wider interlayer spacing (0.60.9) nm, would be more convenient to store H2. However, it is very important to consider how accessible is the surface area between graphene layers when the GO layers are stacked together with the presence of surface functional groups and lamellar water molecules [107]. Under normal conditions, GO usually contains different oxy-functional groups, thus not all of the available sp2 carbons and interlayer spaces are accessible for gas sorption or storage. In recent investigations, some open GO structures are modelled with pillared GO layers, in which further increase in the interlayer distance to

about 1.1 nm is shown as optimal for practicable H2 storage [107,108]. Burress et al. design a series of idealised pillared-GO model systems with a phenyldiboronic acid linker, called GO frameworks (GOFs) (Fig. 5) [107]. The variation in linker concentration produces different interlayer distances, pore sizes, pore volumes and surface areas. The GOF-32 structure is characterised by one linker per 32 graphene carbons predicted to hold ca. 6.1 wt% of H2 at 77 K and 1 bar. In another case, Li-doped CNT pillared GO (Li-PGO) model structures with an interlayer distance of 1.1 nm and pore size of 2.3 nm show a gravimetric and volumetric H2 capacity >10 wt% and 55 g/L, respectively at 77 K and 100 bar [109]. The interaction binding energy of H2 in the system is ~15.5 kJ/mol. First-principles computations on the Ti-grafted GO show gravimetric and volumetric H2 storage of 4.9 wt% and 64 g/L, respectively [110]. Similar to the Li in Li-PGO, the Ti atoms bind strongly to the oxygen sites on the GO thus preventing them from clustering. It is also estimated that each Ti binds multiple H2 molecules with suitable binding energy of (14-41) kJ/mol. The interlayer distance between 1.2 and 1.4 nm shows an optimal volumetric density of (64-74) g/L. First-principles calculations on the covalently bonded graphenes (CBGs) show enhanced H2 binding energy of 20-150% or more compared to the isolated graphene [111]. The stable porous CBGs are formed through sp3 carbons at edges of the graphenes. CBGs with adsorbed Ti atoms show p4 wt% of H2 adsorption depending on the pore size and length of the metal chains.

3.1.3. Doped, functionalised and metal dispersed graphenes

Simple graphene-based nanostructures show weak binding energy thus low H2 adsorption capacity at ambient conditions. Thus chemical doping or structural modifications in the nanostructures has been carried out. A considerable amount of theoretical work has been reported on the modified graphene structures with doping and/or surface metal (alkali: Li, Na, K; alkaline earth: Mg, Ca; simple: B, Al; transition: Sc, Ti, V, Ni, etc.; and noble: Pd, Pt) dispersion. This strategy increases the binding energy and adsorption of H2 due to the polarisation of H2 molecules by an electric field and/or hybridization of H2 r or r* orbitals with transition metal d orbitals (the so-called the Kubas interaction [112]) [112114]. Lochan and Head-Gordon reported interesting theoretical results, which describes that the ionic

0.2 04 0.6 08 1


Fig. 5. Top: GCMC simulations for ideal GOF-n structures with n graphene carbons per linker. The structures of three examples with n = 64, 32, and 8. Bottom: The simulated H2 adsorption isotherms at 77 K for several representative GOF structures [107].

metal atoms, such as Li+, Na+, Mg2+ and Al3+, hold up to six H2 molecules with moderate to very strong binding energies between (12-340) kJ/mol, which further depends on the cluster size and complexes [115]. Furthermore, to avoid the issue of structural instability and poor reversibility in metal atomic dispersion (because of the large cohesive energy of bulk transition metals, the aggregation/clustering seems to occur instead of being atomistically dispersed) the structural or chemical defects/doping are introduced in the substrate [111,116,117]. It has been showed that the alkali- and alkaline earth-metal adsorption on graphene surface leads to a metallic state due to charge transfer (Fig. 6) [118,119]. For instance, as shown in Fig. 6, up to 16 wt% of H2 is predicted in the Li and B adsorbed graphene with a binding energy of ~24 kJ/mol [119-123]. H2 adsorption of up to 14.4 wt% per molecule with an average binding energy of ~40 kJ/mol in ethylene-Ti complex is estimated [124,125]. A tensile strain of 10% on graphene increases the adsorption energy of Li (Ti) atom by around 75% (71%) and the gravimetric H2 storage up to 15.4 wt% (9.5 wt%) with an optimal binding energy of about 19 kJ/mol [126]. By controlling the corrugation of individual layers of graphene, the density functional theory (DFT) calculations estimate up to 8 wt% of H2 [127]. The corrugation of the graphene sheet and the controlled inversion of its curvature show a fast storage (on convex sites) and release (on graphene concave sites) of H2 at ambient conditions.

n # of H2 per Li 1 2 3 4

ETRY HI d, - 1.77 A EL- OJ» cV • d, - 1.71 A •¡•d,- l.»A ^ E|-M»eV d, i d, - 1.80 A c d, -1.93 A y c Ki-oj«cV • El-0.22 cV d, - 1.72 A O • a d - 1.94 A « _ T6 E,-0.l7eV • Ej- 0.21 eV d, - 1.72 A m » « • d, - 2.12 A tf » c E.-«lI8IV • K4-0.20 eV


ION GEOM HI + H2 d, - 1.7« À El - 0.8» cV ^ Ei - 0.82 cV d, - 1.71 À **4,- 1.86 A E, - 0.05 cV d, - 1.70 À a d, - 1.92 A f • Ki- 0.41 r\ • Ej-0.2.1 cV - t, L t d, - 1.74 A g 4,- 1.98 A f \ Ej-0.18eV • Ej- 0.21 «V 4,- 1.82 A • 4,-UIA f I « Ei- 0.19 cV • EL,-0.21 cV

• • J, gd- 12.8 wt %

ADSORPT H1 + H3 • I------ • d, - 1.78 A It- 0.K1 CV E, - OJU cV • ^ d, - 1.71 A d, - ij«A 1 , o.o- , \ V- —*r i d,- 1.71 A C JP d ' ' * X C Ei - 0.41 cV Ei-0.24 *V i ^ d, - 1.71 A a W d, - 2.15 A • f> Ej-0.27*V o EI- 0.25 «V d, - 1.77 A JS * d, -2.13 A E|-0.05cV & E7- o.2o »v

Fig. 6. Top: (a) Various adsorption sites H1, H2 and H3 on the (2 x 2) cell, (b) charge accumulation, Aq+, calculated for one Li atom adsorbed to a single side specified as H1, (c) same as (b) for one Li atom adsorbed to H1 site, second Li adsorbed to H3 site of the (2 x 2) cell of graphene. Bottom: Adsorption sites and energetics of Li adsorbed to the (2 x 2) cell of graphene and absorption of H2 molecules by Li atoms. EL is the binding energy of Li atom adsorbed to H1 site, which is a minimum energy site. For H1 + H2 or H1 + H3 configuration corresponding to double sided adsorption, EL is the binding energy of second Li atom and El is the average binding energy. For H1, H1 + H2 and H1 + H3 configurations, E1 is the binding energy of the first H2 absorbed by each Li atom; En (n = 2-4) is the binding energy of the last nth H2 molecule absorbed by each Li atom, and En is the average binding energy of nH2 molecules absorbed by each Li. Shaded panel indicates the most favourable H2 absorption configuration [119].

The B-doping is found to form an electron-deficient structure that enhances the H2 interaction energy [128-131]. The energy surface seems to extend over several graphene carbon sites, making the surface more heterogeneous. B-doping in graphene also enhances the Li and Ca bonding strength on the graphene [130-133]. The DFT calculations estimate approximately (13, 9.9 and 7.9) wt% of H2 adsorption for the adsorbed Li, Al, and Ti atoms, respectively on the B-doped graphenes [133,134]. Lu et al. show a H2 adsorption of (6.4 and 6.8) wt% (at 298 K and 100 bar) in a Li- and Ca-decorated B doped (pore rim of polyphenylene) porous graphene (Fig. 7) [135]. Up to 15 wt% of H2 uptake is predicted in Be adsorbed on B-doped graphene with an average adsorption energy of ~29 kJ/mol [136]. The H2 adsorption capacity of 2.4 wt% is also proposed in Si-doped graphene [137].

H2 adsorption of (6-9) wt% is estimated in the Ca-adsorbed/decorated graphene surfaces [132,138141]. The Ca adsorption on zigzag edge, porous or covalent-bonded graphenes (CBGs) is also reported [142-144]. CBGs doped with B and decorated with Ca clusters show up to 6 wt% of H2 adsorption [144]. The Ca atoms prefer to bind more strongly on the sp3 vertex sites of the pores than to planar graphene sites. This leads to formation of Ca chain structures and exhibit mixed characteristics of multipole Coulomb and Kubas interactions with H2 molecules depending on the occupation of Ca s or d orbital-derived states. A similar Mg-doped GO shows 5.6 wt% of H2 adsorption at 200 K [145]. The hydroxyl groups are reduced and Mg atom is bound to the surface of GO by the ring-opening reaction of epoxy in the form of -(C-O)x-Mg (x = 1, 2). On the surface the positively charged Mg sites and negatively charged O sites produce local electric fields separately to adsorb H2 molecule with a binding

Fig. 7. Top: (a) Unit cell of 2D porous graphene sheet. Grey balls, C; white balls, H; green balls, substitutional sites for B or N doping; small green balls, adsorption sites of Li or Ca above the hexagonal centre. (b) Schematic diagram containing one pore. Bottom: Computed excess gravimetric H2 adsorption isotherms at 298 K. The dotted line corresponds to the 2017 DOE target. The insets show the snapshots of adsorbed hydrogen molecule density from GCMC simulations at 100 bar and 298 K. Grey balls, C; white balls, H; pink balls, B; purple balls, Li; green balls, Ca [135].

energy of (14.5 and 24) kJ/mol at O and Mg site, respectively. The calculations on Al-doped graphenes also revealed interesting results for H2 adsorption [146-148].

The H2 uptake by light transition-metal (LTM = Sc, Ti, V) atom decorated graphenes has also been studied. Each of the LTM atom seems to adsorb on both sides of graphene and hold up to four H2 molecules to give >6 wt% of H2 adsorption with the average binding energies between (29-48) kJ/mol [149]. The Sc decorated on either armchair or zigzag edges of graphene nanoribbons show >9 wt% of H2 adsorption [150]. Ni atomic dispersion on graphene yields about 3 wt% of H2 but is expected to reduce experimentally due to the clusterization and nanoparticle formation [151].

3.1.4. Graphane based materials

Graphane is nothing but the graphene with alternately covalent bonded hydrogen atoms on each carbon atom on both sides of graphene sheet, as shown in Fig. 8 [152]. Structurally graphane is crumpled, rather than planar, because each hydrogen atom bonded to carbon pulls it a small distance out of the plane. First-principles studies show that the chemisorption of hydrogen on both sides of graphene, i.e., with full coverage of hydrogen (CnHn, i.e., 8 wt%) is the lowest energy structure due to the strain associated with the sp2 to sp3 conversion [153]. Due to its stable structure and small size, graphane has been also considered as a potential candidate for H2 storage if doped with alkali- and alkalineearth metal (M) [154]. The H2 adsorption mechanism on these doped graphanes is explained by considering the fact that positively charged M+ ions polarise the H2 molecules, which are then held to the M+ ions by the vdW attractive forces. Ab initio calculations by Antipina et al. show that the graphane with chemically bonded alkali metals adsorbs up to 12 wt% of H2, with a binding energy of ~19 kJ/mol [154]. The calculations by Hussain and Ahuja's group also predict practicable H2 adsorption in the Li-and Ca-doped graphane with and without strain [152,155,156]. The optimised structure for the case of maximum coverage (5 H2 molecules, each on both sides of Ca atoms) is shown in Fig. 8.

3.1.5. Hydrogen spillover

The H2 adsorption by hydrogen spillover is a hydrogen dissociative atomic adsorption, and it has been found to be other alternative to enhance H2 adsorption in the carbonaceous materials. The spillover is expected in some of the graphene based materials with functionalized/doped structures [157162]. For example, the stability of hydrogenation states of graphene and conditions for hydrogen spillover are described by Han et al. [161]. The mechanism in metal-doped carbon materials and other solid porous media is explained by George et al. [162]. In brief, it can be explained in three elementary steps as shown in Fig. 9; (i) the activation or dissociation of H2 molecules on catalytic metal sites, (ii) a

Fig. 8. Left: (a and b) are the side and top view of the optimised structures of pure graphane, respectively, (c and d) show the side and top views of the optimised structures of Ca adsorbed graphane, respectively. Right: The optimised structure of the fully hydrogenated Ca adsorbed graphane with maximum number of hydrogen molecules adsorbed around each Ca atom. The extended structure of relaxed Ca adsorbed graphane with adsorbed H2 molecules is also shown [152].

transition of atomic hydrogen across the metal-receptor interface, and (iii) the migration of atomic hydrogen throughout the receptor surface. The whole process is denoted as 1, 2; 3, 4; and 5, 6, respectively. The receptor, 'R' is considered inert towards the unactivated H2 and thus only accepts the activated atomic hydrogen that migrates from the catalyst, 'C. The receptor surface is expected to be the key that simultaneously provides appropriate thermodynamic stability (binding energy) and kinetic mobility (migration). Fig. 9 also shows stability of the hydrogen chemisorption state of graphene in comparison with the chemical potential of H2 gas.

Effective dissociation of H2 molecules into H atoms is observed with the firmly bound catalytic Ptj metal cluster on B-doped graphene [160]. The H atoms then migrate from the bridge site of the H-sat-urated Pt4 metal to the supporting graphene sheet. By DFT calculations, Cabria et al. reveal that the adsorption and dissociation of H2 on Pd clusters on the graphene is due to either (a) the adsorption of the H2 in an activated state or (b) the dissociation of the H2 and the chemisorption of the two hydrogen atoms [163]. In an activated state, the H2 bond length is slightly stretched but not broken. The binding energies of the activated H2 are between (0.4-1) eV/molecule, close to the estimated energies required to obtain reversible H2 storage at room temperature and moderate pressures. Tight-binding model describes that the C-Pd-H2 interaction is much weaker than the C-Pd-H bonding [164]. The DFT studies also suggest that the graphene-like materials comprising of curved graphene fragments and many edge atoms, such as microporous carbon, strongly interact with the spillover hydrogen atoms [165]. At 300 K, Li anchored on the graphene vacancies show 6.2 wt% of H2 dissociative adsorption with a strong binding energy of ~84 kJ/mol [166]. Lueking et al. show that both the B substitu-tionally doped graphene and hydroxylated graphene satisfies the constraints for reversible room-temperature hydrogenation [167]. A significant enhancement in the binding of H2 on N-doped graphenes is expected through dissociative adsorption and diffusion on the surface with an externally applied field [168,169]. Under an applied electric field, the N-doped graphenes exhibit a relatively low dissociative adsorption energy barrier for H2 and low diffusion energy barrier of H atoms.

3.2. CH4 and CO2: sorption, storage and separation

Similar to H2 adsorption, the graphenes with ripples, defects, pores, functional groups, metal dispersion or doping and pillaring also show favourable CH4 and CO2 adsorptions [170-181]. The GCMC

V. Са5<"г> : fH.C.r) , «

Fig. 9. (a) Schematic description of the H2 spillover process. C and R indicate the catalytic metal island and the receptor surface, respectively. (b) Chemical potential (eV) of H2 gas with respect to the zero-temperature static energy of the H2 molecule. (c) Static chemisorption binding energy (eV) per H2 onto graphene. The inset shows the chemisorption geometry of 24 hydrogen atoms in the chair-like aggregated configuration. Small (pink) balls represent hydrogen atoms. The carbon atoms to which hydrogen atoms are attached are described with larger grey balls. Other parts of the graphene are denoted only with the wire frame [161].

simulations on fullerene intercalated graphene nano-containers show enhanced CO2 adsorption with increasing fullerene concentration. The pore with an effective diameter of ca. 0.5 nm is found to be optimum for the effective CO2 adsorption at low pressures due to the energetically favourable surface interactions [170]. In a relatively large slit-pore, (1.0-2.4) nm the high CO2 adsorption occurs at relatively high pressures due to multilayer molecular adsorption, similar to the H2 adsorption characteristics discussed above. In another study, molecular simulations demonstrate the CH4 adsorption in multilayer graphene nanostructures with and without Li-doping (MGNs and Li-MGNs) and function of interlayer distance [171]. The MGNs and Li-MGNs with an interlayer distance of (0.34-0.51) nm (0.34 + 0.17) show no CH4 adsorption due to steric effect of the adsorption space, whereas an interlayer distance between (0.68-2.04) nm yield a good CH4 adsorption (Fig. 10). A maximum uptake of 310 v(STP)/v of CH4 is obtained at 35 bar, with an adsorption energy of 27.2 kJ/mol. The structures also show a higher H2S-CH4 adsorption than CO2-CH4 in case of biogas separation. Several main influencing factors for CH4 storage and acid gas separation in MGNs and Li-MGNs are identified, including the binding energy between gas molecules and substrates, the optimal layer spacing for CH4 adsorption,

Fig. 10. Comparison of adsorption isotherms of pure CH4 in (a) MGNs, and (b) Li-MGNs with different graphene layer distances in the range of 0.34-2.04 nm at 298.15 K [171].

and the Li doping strategy. Li cation is able to enhance the binding of CH4 and H2S molecules with the graphene substrate surface via nonbonding interactions. The CO2 molecule interacts with Li cation through chemical bond attributed to the high electronegativity of O atoms.

DFT calculations on the graphene layer modified with Ti at high metal coverage (C2Ti) show a dissociative and associative adsorptive behaviour for both CO2 and CH4, respectively at 300 K and 1 bar [172]. First-principles calculations on the Ti-decorated GO (Ti-GO) show a selective adsorption of CO over CO2, N2, CH4 either in single component or mixture with computed binding energy of ~70 kJ/mol for CO, ~20 kJ/mol for N2 and <5 kJ/mol for CO2 and CH4 [173]. The strong interaction between CO molecule and Ti is a result of dative bond, i.e., hybridization between an empty d orbital of Ti and an occupied p orbital of CO. The plane-wave electronic DFT calculations suggest that CO2 binds stronger to the defective site on graphene surface (with one carbon atom missing, i.e., monovacancy) than a perfect graphene with a physisorption energy of ~20 kJ/mol, approximately 4 times stronger than a perfect graphene surface [174]. The adsorption and reactivity of CO2 on a defective graphene sheet shows strong physisorption energy of ~136 meV [175]. Chemisorption of CO2 is expected by forming a lactone group with the C atoms surrounding the vacancy after overcoming an energy barrier of ~1 eV. Further dissociation of the CO2 occurs through the formation of stable epoxy groups with ~614meV. Ab initio and molecular mechanics calculations show that the periodic nature of the vdW potential energy stored between CH4 and perfect sheet is altered due to the insertion of vacancies (missing lattice atom) and sinusoidal ripples. The CH4 molecule avoids to be around vacant cites and on top of the peaks [176]. The produced vdW energy surface indicates that around the vacant cites the energy increases. On the perfect graphene, DFT calculations reveal that the adsorption of CH4 is essentially driven by dispersion interaction with a reasonable binding energy of 0.17 eV [177].

First-principles simulations show the large CO2 uptake capacities (~0.4-0.6 g CO2/g sorbent) by decomposition-mediated adsorption in Ca-decorated graphenes [178]. In the CO2 uptake process, decomposition into CO and formation of CaO is found to be thermodynamically favourable. Thus enhancement in CO2 adsorption is attributed to a secondary process, where further absorption of CO2 in a CaO leads to formation of CaCO3. The Ca-decorated graphene also shows an improved CO2/ N2 selectivity, because of stronger binding of CO2 than N2, e.g. 2.731 eV/CO2 against 0.645 eV/N2 in case of 12.5% Ca-graphene. Ohba and Kanoh show the edge effects of graphene on Ar, CH4, N2, CO2, and H2O gas adsorption [179]. The edge sites provide strong Coulombic interactions thus molecules with a strong dipole or quadrupole moment (H2O or CO2) are selectively adsorbed on the edge sites. The dispersion interaction-dominated molecules (Ar, CH4, and N2) are selectively adsorbed on the basal planes. Furthermore edge functionalized graphenes with the polar groups, including -COOH, -NH2, -NO2 and -H2PO3, show enhanced CO2 and CH4 adsorption due to strong binding of activating exposed edges and terraces [180,181].

3.3. Gas separation membranes

Pristine and porous 2D graphene membranes with and without functionalization have been investigated for gas purification [182]. The pristine graphene is impermeable to He and H2 [183-185]. Calculations by Miao et al. show that protons can readily pass through a perfect graphene sheet with a low tunnelling barrier compared to a substantially higher barrier for H2 [184]. Fig. 11 shows initial and final energy states for the tunnelling of an H atom through the centre of a hexagonal carbon ring (hollow site). H2 is physisorbed at a distance of 0.286 nm with a high energetic activation barrier of 2.86 eV. The tunnelling of an H atom through the benzene ring causes the ring to expand and then contract between a diameter of 0.285 nm and 0.291 nm in the initial/final and transition states, respectively. The strong orbital overlap between the electrons from the H atom and the graphene at -21.62 eV and 2.08 eV for C-2s and H-1s, and C-2p and H-2s respectively, revealing a strong interaction between the H atom and graphene sheet at the transition state (Fig. 11b). The penetration barrier decreases exponentially with the size of the defects, expressed by the number of carbon atoms involved in formation of defect. The potential energy as a function of the distance of an H atom from graphene indicates that a larger defect size result in a smaller penetration barrier. For example, the barrier is reduced largely to 0.66 eV for DV (divacancies) and 1.17 eV for 555-777 (one reconstruction

-■—i—1—i—1—i---1-■—i—'—i--— -1-1-1-1-1-1-1-1-1-1-1-1-1—

-0.2 0.0 0.2 0.4 0.6 0.8 1.0 1.2 -25 -20 -15 -10 -5 0 5 10

Reaction coordinate Energy (eV)

Fig. 11. (a) Reaction energy profile for a hydrogen atom directly penetrating a graphene sheet. (b) PDOS (partial density of states) of initial and transition states for hydrogen penetration (zero energy corresponds to the Fermi level) [184].

from a DV) defects compared to the 2.64 eV of a perfect graphene. For this defect sites the stable states of H atom found at different distances between 0.12 nm and 0.15 nm.

First-principles DFT calculations estimate that the H2/CH4 diffusion selectivities are on the order of 108 and 1023 for the N- and all H-functionalized pores, respectively [186]. The pores are created by the removal of two neighbouring benzene rings on a single-layer graphene. A CO2/N2 selectivity of around 300 with a free energy barrier for permeation of 24.7 kJ/mol and 9.6 kJ/mol for CO2 and N2, respectively, is reported for the same porous graphene using classical molecular dynamics simulations [187]. A sandwich-like model of porous graphene with the dimensions of 10 nm x 10 nm and the nanopore density of 0.04 per nm2 shows a linear dependence of flux on CO2 pressure. The smaller pores of graphene with periodically ''one missing'' phenyl rings (the polyphenylene unit cell) and line-defect-graphene with missing octagons show a high selectivity for H2 (e.g., 1017 for H2/CO2 and 1023 for H2/NH3) separation from O2, N2, NH3, CH4, CO2 and CO, since CO2, CO and CH4 has larger kinetic diameter and H2 is about width of the pore [188-190]. A controllable selectivity and permeable separation of H2/N2 is demonstrated in a series of porous graphene membranes with different pore sizes and shapes created by removing 11-16 carbons [191]. For example, the pore size of ~0.65 nm (pore-16) show best selectivity of N2 from H2 because of the vdW interactions with graphene membrane that leads N2 molecules to distribute within a single molecule layer around the graphene surface. Gas permeation in various limits of gas diffusion, surface adsorption, or pore translocation as the rate-limiting step is described by analytical expressions [192].

First-principles DFT and MD simulations on 0.135 nm pore membrane show a high selectivity for H2 over other gases with the energy barrier of (0.12, 0.26, 0.25 and 0.82) eV for H2, N2, CO and CH4, respectively [193]. Clearly, it is the electron overlap between the gas molecules and the pore that causes the energy barrier and hinders the molecules from passing through the pores. Thus the energy barrier for gas molecules passing through the membrane generally increases with decreasing pore size or increasing molecule kinetic diameter. Lu et al. demonstrate a high permeable selectivity for H2 relative to CH4, CO, and CO2 in a B or N doped pore rim of (polyphenylene) of porous graphene (PG, one hexagon missing) [135]. The 4B-PG shows a highest selectivity (1.2 x 1078) for H2/CH4 much higher than the PG (6.0 x 1058) with respective energy barrier of 0.91 eV and 0.68 eV. Substitutional doping leads to a change in the original bond length of PG, where stretching and shortening of the bond is observed by B (BH) and N (NH), respectively. Thus the 3B-PG and 4N-PG with change in pore size possess high selectivity (5.0 x 1037 and 2.0 x 1036) for H2/CO (20 times) and H2/CO2 (5 orders) respectively, compared to the PG. In the N-doped pore rim graphene (replacement of the unsaturated two C atoms with two N) a Lewis-acid-base type interaction leads to a bound state of CO2 with a binding energy of about 21 kJ/mol [194], which show no barrier for N2. Due to similar kinetic diameters and no electric dipole moments, it is difficult to separate CO2 from N2; and pore functionalization is needed to enhance the electrostatic interaction. The modified graphene membrane with -CH3, -OH, -NH2,

-COOH shows a higher permeability for both N2 and CO2 compared to the unmodified one [195]. The results show that the porous graphene membrane with all-N modified pore-16 exhibits a higher CO2 selectivity over N2 (~11) due to the enhanced electrostatic interactions compared to the unmodified graphene membrane. Functional groups also increase the active surface area thereby enhancing the uptake and diffusion onto the membrane surface for more permeation. The CO2 adsorption ability of functionalized graphene is in the following order: -OH > -COOH > -NH2 > -CH3. The porous membrane with carbon atoms bonding to three carbons and to two carbons with hydrogen generates nonuniform charge surfaces and strongly interacts with polar adsorbate molecules. A high selectivity for SO2 adsorption (SO2/N2, SO2/CO2) is reported on fluorinated porous graphene due to reduced dispersive interactions, e.g., for CO2 [196]. Furthermore, in a porous graphene the enhanced diffusion rates of H2, O2 and CO2 up to 2, 5, 9 and 7,13, 20 orders of magnitude are obtained by application of 10% of the uniaxial tensile strain and symmetrical stretch strain, respectively [197].

Clearly theoretical simulations have offered further insights into the mechanisms of molecular interactions with graphene-based structures. Pristine graphene only show weak binding with most of the molecules. Enhancement of binding interactions can be achieved by careful design of micropo-rosities, surface doping, defect engineering, and high density of edge/boundary defects. However, the binding energy needs to be retained at a level to cause ''chemical activation" of the molecules, typically (10-40) kJ/mol, rather than their complete decomposition. Separation of such molecules is another challenge, given the little differences in the kinetic diameters of the molecules of interest. Modifications that can cause preferential binding of a certain molecules are desirable to promote selectivity of the sorbent materials. The following section will examine some of the experimental developments along those directions.

4. Experimental developments

4.1. Chemically reduced/thermally exfoliated graphenes

Graphenes in the bulk quantities have been synthesised by liquid phase oxidation-reduction/exfoliation of graphite (Fig. 1a). The degree of oxidation and exfoliation or reduction conditions yield individual single-layer graphenes to loosely packed multi-layer graphenes with a wide range of pore structures, different degree of residual functional groups and accessible specific surface areas (unless otherwise specified, hereafter the term 'specific surface area (SSA)', is generally determined using the Brunauer, Emmett and Teller (BET) [198] method from nitrogen gas adsorption-desorption isotherm measurement at 77 K). In many cases, the reduction of GO and agglomeration limits the microstructure and SSA. For example, in a report by Ma et al., the graphenes obtained by thermal exfoliation of GO followed by hydrogen reduction show a much limited SSA (156 m2/g) due to the severe agglomeration of crumpled graphene sheets [199]. The sample adsorbs about (0.4 and 0.2) wt% of H2 at 77 K and 290 K under 1 bar and 60 bar, respectively. The exfoliated graphenes show an enhanced heat of H2 adsorption of ~6 kJ/mol compared to carbon nanofibers and active carbons [200].

A considerably improved SSA and gas sorption in graphenes have been obtained from different post-synthesis treatments of GO. Srinivas et al. report an enhanced SSA (640 m2/g) in a hydrazine hydrate reduced colloidal suspension of exfoliated GO (Fig. 12) [201]. The sample is characterised with interconnected, wrinkled and disordered graphene sheets with regions of agglomerated multiple layers and a partial restoration of graphitic structure. The gravimetric H2 uptake measurements at (77298) K and up to 10 bar show a heat of adsorption of ~6 kJ/mol at low coverage, and the maximum adsorption of 1.2 wt% at 77 K and 0.1 wt% (and up to 0.72 wt% by extrapolating pressure to 100 bar) at 298 K (Fig. 12). A thermally reduced GO by simple out-gassing at 423 K overnight shows an improved H2 adsorption of about 1.2 wt% at 77 K and 1 bar with a SSA of 477 m2/g due to enhanced microporosity and pore volume of ~1.0cm3/g [202]. In another case, a mild chemical reduction of GO with glucose yields a highly enhanced SSA up to 1200 m2/g with a mesoporous structure [203]. Surprisingly, a high H2 adsorption of 2.7 wt% at 298 K under 25 bar is obtained, a much higher than for most of the carbon nanostructures. Lee and Park report H2 uptake of up to 1.8 wt% at 298 K under 100 bar (Fig. 13) in a series of exfoliated GO's by CO2 pressure swing [204]. GO exfoliation is carried


0 3000 6000 9000 12000 15000

Hydrogen pressure (mbar)

Fig. 12. Top: TEM image of graphene material. Bottom: Hydrogen adsorption-desorption isotherms of graphene material measured at different temperatures and pressures up to 10 bar using gravimetric analyser, for comparison, the hydrogen uptake isotherm for 1-2 im synthetic graphite powder is also measured. Inset shows the isosteric heat of adsorption of graphene powder [201].

out at 373 K under high pressure of CO2 up to 40 bar followed by rapid release to vacuum to produce graphenes with a maximum SSA of 547 m2/g. Magnesium combustion derived graphene sample shows a more or less linear H2 uptake with increasing H2 pressure up to 300 bar with the final capacity of 0.9 wt% at 293 K [205]. Bulk graphene sample is produced via high temperature combustion of magnesium in a CO2 atmosphere using dry ice black, which exhibits a more ordered carbon lattice than other bulk graphenes derived from GO.

Rao's group prepare various types of graphenes from thermal exfoliation of GO (EG), pyrolysis of camphor (CG), and conversion of nanodiamond (DG) [206,207], which shows a wide range of SSA: 46 m2/g, (280-1013) m2/g and (639-1550) m2/g for CG, DG and EG, respectively. H2 adsorption at 77 K and 1 bar reveals an average linear trend with the SSA, in good agreement with other porous carbons. From this trend, the single-layer graphene, i.e. for the SSA of 2630 m2/g, is estimated to hold about 3 wt% of H2, in other words, about 1.2 wt% of H2 for every 1000 m2/g of SSA, at 77 K and 1 bar. In addition, a considerably higher H2 adsorption up to 3 wt% at 298 K and 100 bar is reported.

Fig. 13. H2 uptake (top) at 298 K and N2 (bottom) at 77 K adsorption and desorption isotherms of the various exfoliated graphene oxide samples (filled and hollow symbols correspond to the adsorption and desorption profiles). The exfoliated GO samples obtained are designated by the applied pressure (unit: bar), i.e., GO (for the unmodified sample), E-P-10, E-P-20, E-P-30, and E-P-40 [204].

Guo et al. obtain a highly improved H2 adsorption of 4.0 wt% at 77 K and 1 bar in a hierarchical pore structure graphene with a SSA of 1305 m2/g [208]. Briefly, the sample is synthesised by exfoliation of GO at 423 K under vacuum followed by further heat treatment at 873 K. The pore structure analysis from the N2 (at 77 K) and CO2 (at 273 K) adsorption-desorption isotherm measurements reveal a hierarchical mesopore and ultramicropore structure. This ultra-microporosity leads to an enhanced H2 uptake at low-pressure as well as a high heat of H2 adsorption, 9.5 kJ/mol.

The graphene nanoribbons (GNRs) with in-plane and edge sp2- and sp3-hybridized carbons show a strong irreversible adsorption of CO2 (0.22 mmol/g) at 303 K and 1 bar [209]. GNRs are synthesised through CVD via aerosol pyrolysis using ferrocene/thiophene/ethanol precursors. Adsorption of 6.4 mmol/g of CO2 at 30 bar and 298 K is obtained in a CO2 pressure swing exfoliated GO with a SSA of 547 m2/g and a total pore volume of 2.47 cm3/g [204]. Furthermore, a greatly enhanced CO2 adsorption of 21.6 mmol/g at 11 bar and 298 K is reported in a hydrogen exfoliated graphene (HEG) with a wide distribution of mesopores and SSA of 443 m2/g [210]. Surprisingly even higher CO2

adsorption of up to 56 mmol/g is measured in a vacuum exfoliated GO compared to the maximum of (8, 15 and 39) mmol/g for zeolite-13X, activated carbon and conventional graphene nanosheets, respectively at 298 K and 30 bar (Fig. 14) [211]. The samples exhibit corresponding SSA and total pore volume of (480, 839, 1453 and 701) m2/g and (1.730, 0.396,1.382 and 0.727) cm3/g, respectively. The high CO2 uptake isotherms in graphene samples exhibit unconventional adsorption behaviour than other porous solids. At low or moderate pressures up to 10 bar, the CO2 adsorption in graphenes is not comparatively better than others. The improvement at high pressures is linked to the higher interior void volume.

The chemical functionalization and surface grafting of graphenes with basic groups is also demonstrated to obtain increased CO2 binding and adsorption capacities through enhanced Lewis base-acid interactions [212,213]. Amine-grafted GO synthesised from different amines, including ethylenedia-mine, diethylenetriamine and triethylene tetramine and aqueous GO by vigorous stirring and reflux at 353 K followed by washing shows about 1.2 mmol/g of CO2 adsorption in the breakthrough experiments [212]. An unprecedented CO2 uptake of over 70 mmol/g of CO2 (at 10 bar and 298 K) is obtained in the HEG-polyaniline composites [213]. However due to the very strong chemical binding of CO2, the heat treatment under vacuum is necessary to desorb the CO2.

4.2. Layered graphenes: simple expansion to pillaring with metals and organic ligands

As described in the theoretical predictions there is an optimum interlayer distance between graph-ene sheets to maximise the H2 uptake. Such a range of distance reflects the energetically favourable interaction of one H2 layer between graphene layers. In order to demonstrate the importance of an interlayer distance for effective H2 storage a variety of hybrid structures have been designed. Below we summarise the experimental results on graphene layer structures synthesised through intercalating various chemical species starting with simple metals/oxides, molecules to organic ligands and combination of these. The best example for this case is the well-known graphite intercalation compounds (GlCs) [214-217]. GlCs are highly ordered periodic layer structures, known for their very large graphene-graphene inter-layer spacing, made up of layers of foreign atoms/molecules sandwiched between single to n-graphene layers, depending on the stage-n compound. Stage-n represents the n-number of graphene layers present between two intercalated layers. For example, the stage-1 and stage-2 potassium GlCs, KC8 and KC24 contains one and two graphene layers between two adjacent potassium intercalated layers (Fig. 15) [218]. This results in increased interlayer distance over 0.540 nm, sufficient to reversibly adsorb H2 at low temperatures. In particular, the series of stage-2 GlCs of stoichiometry MC24 (M = K, Rb, Cs) are reversibly and rapidly adsorb H2 up to ~2H2/M, without dissociation and behaves as a molecular sieve (Fig. 16) [218,219]. The charge transfer to the graphene

Fig. 14. CO2 adsorption isotherms of GNs, zeolite-13X, activated carbons, and conventional graphene [211].

Fig. 15. (a) The \fl x \fl R19.11" structure, frequently used as an approximation of the low-temperature in-plane arrangement of KC24. This model has a stoichiometry of KC28. The sorption site (blue square) and neighbouring sites (red triangles) are depicted for a single trigonal subunit cell. (b) Schematic side view of a stage-2 alkali-metal graphite intercalation compound [218].

| ..............

o 4 ■

0.5 1.0 1.5

Surface Excess Adsorption (H2:K)

Fig. 16. (a) Volumetric H2 adsorption isotherms of KC24 measured at 77 and 87 K. The y axis gives the specific surface excess adsorption as a mole ratio of H2 to K. The inset enlarges the low pressure region of the isotherm. (b) Isosteric heat of adsorption (kJ/mol) as a function of surface excess adsorption [218].

planes lead to enhanced binding energy for H2 adsorption (8.4 kJ/mol) in KC24 that is about twice to bare graphite.

An enhanced H2 adsorption is observed in the herringbone graphite nanofibers (GNFs) when expanded the graphite lattice using intercalation of concentrated sulphuric/nitric acid mixture

followed by a thermal shock [220]. The SSAof 47 m2/g in GNFs is increased to (67 and 555) m2/g when treated at 973 K and 1273 K, respectively. Thus expanded GNFs show a multifold increase in H2 adsorption (1.2 wt% and 0.29 wt% at 77 K and 300 K, respectively at 20 bar) compared to unexpanded GNFs. In order to achieve an optimum interlayer distance, Kim et al. obtain the H2 isotherms in a series of pillared-GO at 77 K and 1 bar [221]. Interlayer distance is controlled using pillaring diaminoalkanes (NH2(CH2)nNH2, n = 2, 6, 10) and subsequent thermal annealing between (300-598) K. The structures show a maximum H2 uptake for an optimum interlayer distance of 0.63 nm, in good agreement with the first-principles calculations for graphitic materials. In their earlier work [222], the GO itself used to tune the interlayer distance without inserting other pillaring supports. A careful reduction in a controlled thermal annealing not only retains the layered structure and also shows improvement in SSA, up to 192 m2/g and pore volume, 0.689 cm3/g, compared to 57 m2/g and 0.076 cm3/g for the initial GO. Nuclear magnetic resonance (NMR) measurements confirm a considerable reduction of O- and OH- functional groups in the annealed GO. As high as 4.8 wt% of H2 adsorption at 77 K and 90 bar is measured in the annealed GO with an interlayer distance of 0.64 nm. A considerable reduction in the H2 uptake is observed with a little increase in layer spacing over 0.64 nm.

In another approach, the porous networks of graphenes engineered by bridging the graphene layers and metal-oxide particles through surface functional groups. For example, the GO-transition metal-oxide composites; GO/V2O5 and GO/TiO2 are developed through binding of C-OH species of GO and the oxygen on the transition metal-oxide via a dehydration reaction [223]. The hydrothermally designed porous graphene-Mn3O4 nanocomposites from the aqueous GO dispersion and MnO(OH)2 colloids yield an enhanced SSA up to 680 m2/g [224]. The composite shows ~2.5 mmol/g of CO2 adsorption at ambient pressure and temperature compared to ~1.75 mmol/g in a controlled graphene. The solvothermally synthesised graphene-iron-oxide hybrid structures reveal an enhanced micropo-rosity with SSA ranging between (418 and 901) m2/g [225]. The nanoporous graphene-polyoxometa-late (GPOM) hybrid structures are synthesised from in-situ hydrazine hydrate reduction of GO aqueous solution and cross-linker phosphomolybdic acid [226]. POM is a polyatomic ion, usually an anion that consists of three or more transition metal oxyanions linked together by shared oxygen atoms to form a large, closed 3D framework. The polynuclear metal-oxo structured POM is a versatile building block cluster for construction of functional hybrid materials. The hierarchical pore GPOM shows a pore volume of 0.57 cm3/g and SSA of 680 m2/g, considerably larger than SSA of 23 m2/g for GO and 8 m2/g for POM. A highly increased interlayer spacing of 3.2 nm and 3.6 nm is achieved in GPOM and o-GPOM (oxidation of GPOM by hydrogen peroxide), respectively from 0.4 nm in a reduced GO. H2 adsorption of 0.8 wt% and 1.3 wt% in GPOM and o-GPOM respectively, is obtained at 77 K and 1 bar. From the results, it is worth noting that the (micro)pore size distribution, adsorption enthalpy, exposed metal-sites and functional groups play a critical role in enhancing H2 adsorption.

Using a well-known reactivity between boronic acids and hydroxyl groups, Taner's group [107,227] develop a novel pillared GO structures, called GO frameworks (GOFs) by linking the various phenylbo-ronic acids between GO layers (Fig. 17). For the past few years, boronic acids have been considered as versatile building blocks for the construction of a variety of molecular architectures [78,228]. The interaction between boronic acids and diols has been used recently to construct a variety of crystalline porous frameworks, called covalent organic frameworks (COFs) [78]. Thus the rich surface oxy-func-tional groups on the graphene leads to the development of GOFs by forming strong boronate ester bonds as shown in Fig. 17. The structures are synthesised through a solvothermal method using methanol solution from various concentrations of GO and phenylboronic acid. GOFs show a systematic increase in the interlayer distance from 0.75 nm to 1.05 nm and SSA from 10 m2/g to 470 m2/g. About 1 wt% of H2 uptake at 77 K and 1 bar with relatively high (~9 kJ/mol) heat of H2 adsorption at low coverage is achieved. The H2 adsorption capacity for a given SSA and heat of adsorption appears twice as large as typical porous carbons ([82,199,201], typically about 0.5 wt% of H2 uptake at 77 K and 1 bar for every 500 m2/g of SSA and heat of adsorption of 4-6 kJ/mol). GOFs also show a good room temperature CO2 adsorption of ~2.7 mmol/g at 4 bar with again high heat of CO2 adsorption of 35 kJ/mol. The micropore structure with slit-like pores and strong surface interactions with functional groups are favourable factors enhancing their capacity for gas adsorption and binding. Similarly, cross-linked and functionalized graphene sheets from diazonium (product 1) and tert-butylaniline or 4-chloroan-iline (product 2) as reagents in chlorosulphonic acid or oleum are synthesised aiming at improved H2

Fig. 17. Representations of (a) boronic ester [107] and (b) GOF formation by benzenediboronic acid pillars in GO, and the H2 uptake in respective structures [227].

adsorption [229]. Products 1 and 2 show a calculated interlayer spacing of ~0.96 nm and ~0.76 nm and SSA of ~440 m2/g and 350 m2/g, respectively. Enhanced H2 adsorption of up to 1.6 wt% at 77 K and 2 bar is obtained compared to wt% in the starting exfoliated graphenes. The enhanced H2 adsorption again suggests importance of tuning the interlayer separation, microporosity, and func-tionalization. Zhou et al. design a porous graphene structures by cross-linking non-planar terpyridine complexes using azide-alkyne click reaction [230]. The structures show up to 440 m2/g of SSA and a CO2 adsorption of 2.6 mmol/g at 273 K and 1 atm.

The hybrid monolith aerogels of chitosan (CTS, an environmentally-benign biopolymer) with GO, prepared using freeze-drying, show dramatic increase in both the porosity (from 153m2/g to 415 m2/g) and cyclic CO2 uptake (from 1.92 mmol/g to 4.15 mmol/g at 25 °C and 1 bar) for a GO loading of 20 wt% [231]. The DAB poly(propyleneimine) dendrimer (DAB) cross-linked GO-based hybrids display good CO2 adsorption of mmol/g (P/P0 = 0.35) upon wetting [232]. Under humid conditions the sample exhibits ~3 times the adsorption capacity with fast kinetics, compared with the dry state, due to the presence of amino groups of DAB which can form carbamates through zwitterions (primary amines) and bicarbonates (tertiary amines). The formation of these groups is enhanced in the presence of a base (hydroxyl groups of water).

In a work by Matsuo et al., the pillared graphenes are synthesised by pyrolysis of various syl-sesquioxane bridged GOs [233]. The structures are highly microporous in nature (calculated pore widths of about 0.7 nm) and show an interlayer spacing between 1.34 nm and 1.6 nm and SSA up to 940 m2/g, which leads to an H2 excess adsorption of 0.6 wt% at ambient temperature and 200 bar

with a high heat of H2 adsorption of (8-11) kJ/mol. A further enhanced H2 adsorption of 1.4 wt% and 2.6 wt% at room temperature and 50 bar is obtained in a liquid crystal (LC) route prepared macroscop-ically ordered layer-by-layer 3D framework of GO and GO-MWCNTs composites, respectively (Fig. 18) [234]. MWCNTs act as 1D spacer between graphene sheets that result in large increased interlayer spacing up to ~7 nm. The high H2 adsorption is attributed to the synergistic effect of the well-spaced GO sheets and debundling of MWCNTs to enhance the accessible surface area. It is worth noting that the ~1.4 wt% of H2 in a LC-GO with its interlayer spacing of around 0.8 nm is also comparatively higher than an earlier reported value of 0.2 wt% at 77 K in GO [107].

The enhanced H2 uptake capacities are also realised in graphene incorporated clay materials [235,236]. The samples are prepared by impregnation of the natural clay (montmorillonite and sepi-olite) substrates by caramel in aqueous media, followed by a thermal treatment in the absence of oxygen, gives rise to graphene-like materials, which remain strongly bound to the silicate support. These carbon-clay nanocomposites shows ~0.4 wt% of H2 at 298 K and 200 bar related to the carbon mass. These uptakes are still far from the predicted value of (3-4) wt% from model simulations. A very high isosteric heat of H2 adsorption of 14.5 kJ/mol shows a strong stabilization of the H2 molecule within the structure. The controlled clays without graphene, thermally treated under the same conditions

Fig. 18. Top: Comparative H2 adsorption in GO, MWCNTs, GO-MWCNTs and rGO-MWCNTs at room temperature under different hydrogen pressures. Bottom: SEM micrographs of (a) GO paper, (b) layer-by-layer assembled GO platelets decorated by MWCNTs [234].

did not show any appreciable H2 adsorption. The samples also show about (0.6 and 1.7) wt% of H2 uptake related to the total mass of the system and carbon mass, respectively, at 77 K.

In another case the enhanced adsorptive gas storage is achieved by pore engineering of highly porous MOFs with GO [237-240]. The MOFs are a type of crystalline and highly porous solids made up of metal clusters linked with organic ligands. For example, Petit et al. synthesise the MOF-GO composites by solvothermal method, with the addition of GO up to 50% by weight to the MOF precursors [237]. The successful retention of porosity suggesting that the building process of the composites occur via the reaction/binding of the copper dimers from the HKUST-1 (Cu-MOF, also known as Cu-BTC) with/to the functional groups in GO (epoxy, carboxylic, hydroxylic, sulphonic). At the optimal GO content, the composites show about 15% and 18% increase in a micro- and total porosity, and SSA up to 1000 m2/g compared to a 900 m2/g in HKSUT-1. While the composites with the higher GO contents show smeared SSA (600-700) m2/g due to the formation of nonporous reduced graphene stacks from the excess, unreacted GO. H2 adsorption of (2.0-2.5) wt% at 77 K and 1 bar in these structures is consistent with the enhanced microporosity. Later, Liu et al. report a further enhancement in SSA and H2 adsorption as well as CO2, CH4 and N2 adsorptive separation in a similar Cu-MOF and GO composites, named as CG (Fig. 19) [238]. The authors also reveal more insights in the growth mechanism of the Cu-

0.0 0.2 0.4 0.6 0.8 1.0

P (atm)

Fig. 19. Top: HRTEM image of the composite CG-9. Bottom: CO2 adsorption isotherms of Cu-MOF, CG-3, CG-9 and CG-15 at 273 K. The SSA (m2/g), 1305,1470,1532, 500, and total pore volume (cm3/g), 0.518, 0.610, 0.645, 0.312 for Cu-MOF, CG-3, CG-9, and CG-15, respectively. In CG-n, n is a weight percent of GO added to the MOF precursors before synthesis [238].

MOF crystals on GO substrate in the composite through DFT calculations (in a VASP). The results show strong interaction energy of Cu to the GO defect of 15.12 eV, more favourable than those on the -OH (3.92 eV) and -O- groups (1.60 eV) of the GO surface. Furthermore, the epoxy groups and the defects of the GO layers act as seed and termination sites, respectively for formation of nanosized and well-dispersed Cu-BTC on the layers of GO (Fig. 19). Thus the composites show about a 30% increase in CO2 and H2 adsorption (from 6.39 mmol/g of Cu-BTC to 8.26 mmol/g of CG-9 at 273 K and 1 bar for CO2; from 2.81 wt% of Cu-BTC to 3.58 wt% of CG-9 at 77 K and 42 bar for H2). A similar CO2 adsorption is observed in a GO-Cu-BTC with stronger interaction energies (activation energies of (68.6 and 56.7) kJ/mol for CO2 on the composite and Cu-BTC, respectively) and enhanced CO2/CH4 adsorption selectivity, twice that of Cu-BTC for an equimolar CO2 and CH4 mixture [239]. Furthermore, the composites of Cu-BTC and GO modified by urea show enhanced CO2 adsorption of 4.23 mmol/g [240]. The incorporation of aminated GO into a MOF structure shows synergistic effect; the porosity modification with new micro-pores at the interface and basic groups on the surface, which are expected to increase the physical adsorption forces and acid-base interactions, respectively. This would not only increase the adsorption capacity but also the selectivity.

Sandwich-like graphene-silica sheets with incorporation of polyethyleneimine (PEI/G-silica, with about 12 wt% graphene and 60 wt% PEI, amine based solids for CO2 adsorption) show a high CO2 adsorption of 4.3 mmol/g at 348 K under 7.4 bar with good cyclic stability [241]. The G-silica sheets obtained through the hydrolysis of tetraethylorthosilicate on the surface of GO with the aid of a cat-ionic surfactant (cetyltrimethyl ammonium bromide) followed by pyrolysation at 973 K under Ar. PEI impregnation is made via a solution route using methanol solvent. PEI/G-silica sheets show a pore size of (0.7-2.0) nm with a SSA of 32 m2/g, much smaller than that of G-silica sheets (~2.0 nm and 930 m2/ g). Clear advantage of graphene sheets is observed in an enhanced CO2 adsorption at low equilibrium pressures (<1 bar), whereas the PEI/G-silica sheets show a significantly higher CO2 adsorption than PEI/silica sheets. Sui et al. synthesise 3D cross-linked GO hydrogels with a polyethylenimine (PEI) [242]. As shown in Fig. 20, the structures are developed through gel formation from aqueous GO dispersion with PEI solution at 25 °C followed by freeze-drying. The monoliths possess a low density in the range of (0.02-0.03) g/cm3 due to the open-pore structures interpenetrating the skeleton of GO sheets. This is also well observed in the XRD patterns. Compared with GO, an increase in the interlayer spacing up to 1.56 nm, indicating successful cross-linking of the GO sheets. Therefore, the open-structures show enhanced SSA of up to 476 m2/g and CO2 adsorption of ~2.4 mmol/g at 273 K and 1 bar compared to 1.8 mmol/g in a controlled GO with a high SSA (876 m2/g). Clearly, the functional basic sites enhance the CO2 adsorption capacity.

4.3. Spillover and dissociative atomic hydrogen sorption

As explained in Section 3.1.5, hydrogen spillover is a dissociation of H2 on the catalytic metal nano-particles, followed by subsequent atomic adsorption onto the substrate surface sites. It is considered as one of the effective ways to enhance H2 storage capacity in the carbonaceous structures [80,243]. In the study by Hu et al., the H2 uptake by low temperature physisorption and room temperature chemi-sorption is clearly demonstrated in a high Pd (about 24 wt%) dispersed graphene nanocomposite [244]. The chemisorptive H2 uptake in Pd nanoparticles shows a two-phase behaviour of isotherms. Initially formed solid solution region at low pressures (a-phase), is transformed to a hydride phase (p-phase) at higher pressure through a coexistence of two-phase region (a + p), plateau range, very similar to the typical pressure-concentration-temperature curves of H2 storage materials [245,246]. The isotherms also show a Pd particle size, (4-24) nm, dependence; a clear transition from a- to p-phase is observed with increase in a Pd particle size. A high H2 uptake of 3.2 wt% at 77 K and 1.1 bar for a given SSA of 230 m2/g and a pore volume of 0.3 cm3/g suggest a clear spillover effect. Furthermore, an exfoliated graphene (with a SSA of 755 m2/g) decorated by Pd and Pt nanoparticles, (25) nm, via an electroless deposition technique also shows enhanced room temperature spillover of H2 uptake [247]. The observed uptake of 4.3 H/Pd is much larger than that of actual chemisorbed hydrogen in the bulk Pd with a stoichiometry of 0.6 H/Pd. The room temperature H2 spillover capacity in a mixture of the activated carbon (AC, Maxsorb) as ''secondary spillover receptor'' and Pd-exfoliated graphene (in a 9:1 weight ratio) shows reversible H2 uptake of 0.82 wt%, 49% more (0.55 wt% of H2)

Fig. 20. (A) Illustration of the preparation process of the GO-PEI porous materials (GEPMs): digital pictures of aqueous GO dispersion (a), GO-PEI hydrogel (b), and GEPM (c) and schematic diagram of aqueous GO dispersion (d), GO-PEI hydrogel (e), and GEPM (f). (B) SEM image of the GEPM. (C) Wide-angle X-ray diffraction patterns of the GEPMs (solid square for shoulder peak position). (D) 77KN2 sorption isotherms of GEPM-1, GEPM-3, and GEPM-4 (solid symbols for adsorption and empty symbols for desorption). The isotherm of GEPM-3 is offset by 100 units. GEPM-1, GEPM-2, GEPM-3, GEPM-4, and GEPM-5 are for the different PEI to GO weight ratios, 3:1, 2:1,1:1,1:3, and 1:6, respectively 1242].

than Pd-free AC-exfoliated graphene under 80 bar [248]. The spillover is also revealed by a high isos-teric heat of H2 adsorption of ~14 kJ/mol, compared to ~8 kJ/mol for AC.

Ramaprabhu's group report up to 3 wt% and 4 wt% of H2 in a Pd (20 wt%) nanoparticle decorated graphenes before and after N-doping (ca. 7%), respectively at 298 K and 40 bar (Fig. 21) [249-252]. The values are comparatively higher than ~0.7 wt% and ~1.5 wt% in a Pd free graphenes before and after N-doping, respectively. A clear influence of N-doping can be observed. A combined physisorption and chemisorption process gives a high heat of H2 adsorption over 12.5 kJ/mol. The hydrogen exfoliated GO at 523 K is used to synthesise N-doping and Pd-decoration through nitrogen plasma treatment and reduction of PdCl2, respectively. The Pd/N-G sample accounts 48% of the spillover capacity enhancement in the total H2 storage capacity. Wang et al. develop a nonprecious and air stable Ni-B nanoalloy dispersed graphenes to enhance the H2 adsorption [253,254]. As much as 4.4 wt% of H2 is achieved at 77 K and 1 bar in the graphene with Ni (0.83 wt%) and B (1.09 wt%) (graphene-Ni0.83B109) compared to 1.35 wt% of H2 in the undoped graphene. The composites are obtained via NaBH4 chemical reduction of aqueous solutions of GO and Ni(CH3OOH)2. The H2 capacity seems excellent among above Pd-graphene and all other carbon-based materials. Up to 12 kJ/mol of isosteric heat of H2 adsorption is reported.

Further higher H2 storage capacity of over 5 wt% at 298 K and 20 bar and heat of adsorption, >30 kJ/ mol at low coverage and ~14 kJ/mol at high surface coverage is reported in the Pd/Hg nanoalloy and Pd nanoparticle loaded GO-like foams (GLFs) [255-257]. Porous GLFs (nominal composition C2OH) are synthesised by simple calcination of the molecular precursor, sodium chloranilate dihydrate in air at 573 K [255]. GLFs consist of interconnected bundles of turbostratically stacked few-layer graphenes

Pressure (MPa)

Fig. 21. H2 adsorption-desorption isotherms of (a) HEG and (b) Pd/f-HEG [249].

build-up of aromatic and aliphatic domains with oxygen rich functional groups on the surfaces. GLFs with SSA of 510 m2/g and pore volume of 0.89 cm3/g selectively absorb CO2 (about 2 mmol/g) over N2 (0.4 mmol/g), CH4 (0.7 mmol/g), H2 (negligible) and CO (0.7 mmol/g) at 10 bar and room temperature. The Pd4Hg particle loaded (12 wt%) GLFs (Fig. 22) with its inherent carbon-based radicals near the edge and defect sites yield an enormously enhanced H2 uptake of about 5 wt% at room temperature and 20 bar compared to a negligible H2 uptake in undoped GLFs [256]. The Pd-Hg alloy loading is obtained through NaBH4 reduction of aqueous GLFs, PdCl2 and HgCl2 mixture. In other study, Wang et al. show an important comparison of H2 uptake by spillover in Pd loaded (10 wt%) GLFs (Pd/GLFs) and super-activated carbon (Pd/AX-21), and oxygen modified AX-21 (Pd/AX-21-O) [257]. By comparing H2 absorption, the authors conclude that the existence of abundant oxygen functional groups on the surface and edges in Pd/GLFs lead to enhanced H2 uptake with a high heat of adsorption, >30 kJ/ mol at low coverage and ~14 kJ/mol at high surface coverage compared to those of Pd/AX-21 and Pd/

Fig. 22. Top: (a) HRTEM image and ED pattern (inset) of GLF. The image is inverted for clarity (carbon matrix is white). (b) HRTEM of GLF/Pd4Hg. Bottom left: H2 sorption-desorption isotherm (at 298 K) of the GLF, original GLF-Pd4Hg (3 cycles) and a 2nd GLF-Pd/Hg sample (prepared from trihydrate). Bottom right: kinetic uptake examples for equilibration pressures of 50 mbar and 4 bar (data from 2nd and 3rd cycle-original GLF-Pd4Hg) [256].

AX-21-O. The H2 dissociative adsorption is also observed in simple modified graphenic structures without metal dispersions. The volumetric H2 storage of 5 wt% is also reported in exfoliated turbost-ratic carbon nanofibers (CNFs) at 77 K and 100 bar compared to 1.5 wt% in the as-produced CNFs [258]. This enhancement is attributed to the simultaneous molecular and dissociative H2 adsorption in combination with more defective surface, increased SSA (from 103 m2/g to 227 m2/g) and pore volume (0.39-0.52 cm3/g). A graphite exfoliation technique using intercalation of acids and thermal shock is employed to expand the CNFs.

4.4. Bottom-up solution-phase synthesised graphenes

Using the bottom-up solution-phase method, graphenes of gram scales are produced from simple laboratory reagents of sodium and ethanol in a 1:1 M ratio by solvothermal method [259,260]. The

subsequent rapid pyrolysation leads to a porous graphene in relatively large quantities (~0.1 g for

1 ml ethanol). The sample with a high SSA of 2139 m2/g shows a high-pressure H2 uptake of 0.9 wt% at room temperature with a high heat of adsorption of 12 kJ/mol [261]. The adsorption is attributed to dissociation of H2 at the zigzag edges of the graphene, or combined effects of residual oxygen and sodium functionalities. In other study, the heteroatom (boron, phosphorous, or nitrogen) substituted carbon scaffolds (micrometer sized flake-like few-layered graphene structure assembled by (3-5) nm sized domains) show H2 capacities of (2.2-2.4) wt% per SSA of 1000 m2/g at 77 K and

2 bar relative to 1.8 wt% for the pristine carbon scaffold or wt% for common carbonaceous materials [262]. The heteroatom substituted scaffolds also show H2 binding energies between (5-9) kJ/mol.

4.5. Highly porous graphene carbons from templating, CVD and KOH chemical activation

As demonstrated above, all the methods; functionalization, exfoliation and pillaring hybrid designs generated a promising gas sorption and storage. However the structures have limited accessible surface areas and pore volumes. To obtain a relatively higher SSA with tuneable pore structure, self- or sacrificial-templating from the ordered structures, CVD and chemical activation methods are employed [263-267]. As shown in Fig. 23, the chemical activation of precursor GO with activator KOH yields a very high SSA of up to 3100 m2/g graphene-carbons. The exfoliated GO by thermal shock or microwave irradiation is thoroughly mixed with various amounts of KOH either by solution or dry milling followed by activation at temperatures between 873 K and 1173 K. For example, Srinivas et al. report a range of SSA (up to 1900 m2/g), pore volume (up to 1.65 cm3/g), pore size and size distribution in GODCs (GO derived carbons) by changing the GO/KOH concentration and activation temperature [263]. The large fraction of smaller pores is transformed to large pores when increasing KOH concentration and/or activation temperature. As shown in Fig. 24, the GODCs show favourable gas adsorption for CO2 and CH4, compared with other high surface area carbons for a given SSA.

(Graphene Oxide Derived Carbons)

Fig. 23. Top: Schematic showing the chemical activation of GO with KOH that creates pores [263]. Bottom: Surface morphology of nitrogen doped KOH activated GO: (a) SEM image, and (b) TEM image [264].


✓ -tv

^ . Porous carbons

500 1000 1500 2000 2500 3000 3500 BET SSA (m2/g)

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500 1000 1500 2000 2500 3000 3500 BET SSA (m2/g)

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- ¿y'A -

- 'O + □ GODCs O COFs + MOFs -

1000 2000 3000 4000 5000 6000 BET SSA (m2/g)

1000 2000 3000 4000 5000 6000 BET SSA (m2/g)

Fig. 24. The 300 K and high pressure CO2 at 20 bar (a and b) and CH4 at 35 bar (c and d) adsorption capacities of GODCs with respect to the BET surface area. (a and c) A comparative CO2 and CH4 adsorption capacities at 20 and 35 bar, respectively from a range of other porous carbons at around room temperature (298-300 K). (b and d) A comparative CO2 and CH4 adsorption capacities at 20 and 35 bar, respectively from a range of other porous organic frameworks; MOFs and COFs [263].

Furthermore, the N-and S-doping in GODCs with pyrrole and poly-thiophene, respectively show an enhancement in ambient CO2 uptake mmol/g) and the selectivity over CH4 and N2 [264-266], The selectivity is an important factor for the application in flue gas capture. For example, as shown in Fig. 25, the S-doped graphene carbons show highly enhanced and stable recycling CO2 adsorption capacity of >4.0 mmol/g over CH4, N2 and H2 (which show much lower adsorption of <1 mmol/g) at 298 K and 1 bar. The capacity is twice as high as for the undoped activated graphene carbons (~2 mmol/g) [263]. The CO2 capture capacity in the doped porous graphene seems comparable with or even higher than the various other porous media; N-doped porous carbons, amine functionalized silica and MOFs. The high microporosity and N-functionality show a high isosteric heat of CO2 adsorption between (30 and 56) kJ/mol. The CO2 adsorption of 1.82 mmol/g at 0.2 bar (is a partial pressure of CO2 in the flue gas stream) in the S-doped sample is also comparatively larger than other porous carbons. The increase in activation temperature yields enhanced microporous samples but same time the N- and S-doping content reduces; reduction from (7 to 2.7) at% is observed at 700 °C, whereas the lower activation temperatures lead to a decrease in SSA. Using porous, layer structure MgO as a template, nanomesh of one- to two-layer graphene in a gram-scale is produced by CVD method using methane as a carbon precursor [267]. After acid treatment to remove MgO template, the graphene sheets show a large corrugations with nanopores of ~1 nm, total pore volume of 2.35 cm3/g and SSA of 2038 m2/g. A high H2, CH4 and CO2 adsorption of ~1.5wt%, 14.5 mmol/g (at 90 bar) and 36.5 mmol/g (at 31 bar), respectively at 274 K is reported.

0.0 0.2 0.4 0.6 0.8 1.0 0.0 0.2 0.4 0.6 0.8 1.0

Pressure (atm) Pressure (atm)

Fig. 25. Left: CO2 adsorption isotherms for SG and a-SG samples at 298 K (filled symbols: adsorption, open symbols: desorption). Right: H2, N2, CH4 and CO2 adsorption-desorption isotherms of a-SG6 at 298 K [266].

4.6. Electrochemical hydrogen storage

The graphene materials with their high electrical conductivity and chemical stability also investigated for electrochemical H2 storage via electro-decomposition of water on a cathodically polarised electrode. The adsorption-desorption mechanism during charge/discharge phenomenon is simply written in terms of the following equations: [G] + nH2O + ne~ m [CHn] + nOH~, where [G] represents the graphene host. In steps; water reduces first (H2O + e~ ? H+ + OH~) and the nascent hydrogen then adsorbs into the graphene network due to the electric polarisation of the electrode under the external applied potential. In the reversible process, the hydrogen atoms discharge from the graphene and recombine with OH~ to make H2O. A reversible H2 storage of ~148 mA h/g is reported in a few layer graphene flakes that are synthesised through arc-discharging of pure graphite under hydrogen atmosphere [268]. The graphene working electrodes are prepared on Ni-foam using PVDF binder and press. The reversible H2 storage seems highly dependent on the structural properties of graphene flakes; crystallite size, wrinkles, defects, and inter layer spacing. Graphene-porous cobalt-oxide (Co3O4) nano-composites show a first discharge capacity of 241.9 mA h/g which is equal to 0.89 wt% of H2 [269]. The nanocomposite on a sheet of nickel foam, Ni(OH)2/NiOOH and Hg/HgO as a working, counter and reference electrode, respectively is employed in a 6 M KOH electrolyte at 298 K. As shown in Fig. 26 a further high discharge capacity of 1415 mA h/g, equivalent to 5.176 wt% of H2 is observed in the ful-lerene-like orthorhombic-structured Co3C nanoparticles, synthesised by ball-milling of Co and graphene powders [270]. Two obvious plateaus at 1.2 V and 0.4 V, respectively in the discharge curve of Co3C nanoparticles reveal an existence of two different H2 adsorption sites; first adsorbs into the interstitial sites/pores between Co3C nanoparticles and then diffuses into the interstitial sites between Co3C. The composition and ball-milling conditions show a profound effect on the cyclic stability of the composites. The continuous decay in a cyclic discharge capacity is attributed to the formation of p-Co(OH)2. Overall, the Co3C nanoparticles display a relatively high discharge capacity when compared with the traditional AB5 type H2 storage materials.

4.7. Chemical hydrogen storage

The hydrogenated graphene, also called as graphane, a covalently bonded hydrogen atoms to carbon atoms, (sp3 C-H bonds on basal plane) is essentially different from the physical H2 absorbed graphene structures. Number of synthesis methods; plasma, CVD, liquid and gas phase reactions are developed to obtain graphanes. Subrahmanyam et al. demonstrate that it can be used as a chemical H2 storage material [271]. Birch reduction of a few-layer graphene with excess Li in liquid ammonia produce the hydrogenated graphene that contains up to 5 wt% of H2 (Fig. 27). Spectroscopic studies

Fig. 26. Growth mechanism and hydrogenation behaviour of the fullerene-like Co3C (weight ratiosCo:GE = 4:1, 800 rpm): (a) starting nanoparticle of graphene and metal Co, (b) partly reacted Co3C, encapsulated with layers of the graphene, (c) the final quasi-spherical closed-cage nanoparticle of Co3C, (d and e) HRTEM images of the Co3C and (f) charge-discharge curves. (g) Capacities and cycle stability of the as-obtained G and Co3C electrodes (the composition CoxC, represents the weight ratios of Co:G = x:1; x = 4-7) at a 30mA/g discharge current density [270].

reveal the presence of sp3 C-H bonds and the hydrogen can be released up on heating or irradiation with UV or laser. Indeed, the hydrogenation of various carbon structures, fullerenes, nanotubes, and graphites via a dissolved metal reduction method with Li and methanol in liquid ammonia has been demonstrated early in 2001 [272]. These structures release about (4-5.4) wt% of H2 upon heating to 873 K. A simultaneous release of H2 and a small amount of methane is observed during the decompo-

Fig. 27. Change in the weight percentage of hydrogen of EGH and HGH with temperature. (Inset) The evolution of H2 is recorded by a gas chromatograph [271].

sition. Recently, a hydrogenated graphene with wt% of H2 (C13H)n, is synthesised directly from a Birch reduction of graphite powder [273].

Simplified techniques have also been developed to obtain chemical hydrogenated graphenes in a large scale compared to complexity involved in a hydrogenation of CVD/mechanically cleaved graphene in hydrogen plasma [274] or Birch reduction of graphene/graphite. For example, a gram scale of hydrogenated graphene directly from the GO is synthesised via in-situ generated atomic hydrogen through the H2 spillover using nickel as an active catalyst under ambient conditions [275]. The hydrogenated graphenes also produced directly by exfoliating GO under high pressures, (60-150) bar of hydrogen at high temperatures, (473-773) K [276]. The hydrogen, fluorine, and oxygen gas phase atomic covalent functionalization to graphene is also explained [277].

Earlier, Orimo and his co-works report a hydrogen concentration up to 7.4 wt% (CH0.95) in the ball-milled graphitic nanostructures (Fig. 28) [278-280]. Hydrogenation is attributed to the formation of dangling carbon bonds, the CHx covalent bonds. Interestingly the milled graphite also shows ~400 m2/g of SSA and H2 physisorption at cryogenic temperature because of random orientation of fragmented small coherent domains that consist of small stacks of 2-3 graphene layers. Very recently, a large scale (5 g for each batch), edge-hydrogenated graphene with 3 wt% of H2 (which corresponds to C/H = 2.3) is produced by ball-milling of graphite powder in a hydrogen atmosphere [281]. Electron microscopy studies reveal the mechanochemical cracking of graphite into small grain sizes of (0.11) im after milling of 48 h correspondingly, the SSA and pore volume increases respectively, (3437) m2/g and (0.002-0.391) cm3/g. The edge-selective hydrogenation during the ball-milling process is expected due to the reaction between reactive carbon species (radicals and ions) at the broken edges of graphenes.

4.8. Graphene membranes and protection barriers

A monolayer graphene membrane is impermeable to standard gases including helium [282]. The nanopore graphene membrane as molecular sieves is demonstrated by measuring the transport of a range of gases (H2, CO2, Ar, N2, CH4 and SF6) through the irradiated pores [283]. As shown in Fig. 29, a membrane fixed over predefined 5 mm diameter wells on SiO2 and pressurised with H2

400 300 200 100

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- CCH = 4.1 wt. % _CH „ . 6.0MPa ; CpH ^ 0.5wt. %

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400 600 800 1000 Temperature (K)

Fig. 28. Left-top: Surface area of the nanostructured graphite, as a function of milling time under initial H2 pressure of 10 bar (1 MPa) at room temperature. Left-bottom: Total hydrogen concentration in the nanostructured graphite, as a function of milling time under initial hydrogen pressure of 10 bar at room temperature [280]. Right: Thermal desorption spectroscopic plots of H2 for the nanostructured graphite prepared by milling for 80 h under various hydrogen pressures. Here, chemisorption hydrogen concentration (CCH) determined by oxygen combustion method and physisorption-like hydrogen concentration (CPH) at 60 bar (6 MPa) and 298 K determined by a volumetric method are listed as well [278].

gas is etched with ultraviolet light leads to a rapid leakage of H2 but preventing N2 from passing through it. A size selective permeation of gas molecules is demonstrated by creating different-size pores. The submicrometer thick, (0.1-10) im and area cm2 of GO membranes are found to be completely impermeable to liquids, vapours and gases, including helium, while allowing unimpeded permeation of water (H2O permeated through the membranes at least 1010 times faster than He) [284]. It is attributed to a low-friction flow of a monolayer of water through 2D capillaries formed by closely spaced graphene sheets. Diffusion of other molecules is blocked due to reversible narrowing of the capillaries in a low humidity and/or by their clogging with water. Mass-spectrometry measurements show that helium is less permeable in dry GO film than a 1-mm-thick glass. No permeation is detected for several polar and nonpolar organic liquids through GO membranes of 1 im in thickness. Reduced GO membranes with an interlayer spacing of 0.4 nm are much less permeable to water. In recent studies, the graphene membranes are also tested for different solvent filtration (desalination) and separation [285,286].

The single layer graphene grown by a CVD method on Ru, Cu and Cu/Ni is found to protect the underneath metal from air oxidation even after heating them in air [287-291]. The surface oxidation of bare copper and oxidation resistance with surface grown graphene is clearly demonstrated in Fig. 30 by controlled CVD growth of graphene directly on Cu foils with partial to full surface coverage [289]. The colour change of bare copper foil indicates a surface oxidation, whereas the fully covered graphene surface shows no sign of oxidation. In other case, the oxidation resistance of Fe and Cu foils are also demonstrated by coating them with reduced GO layers [290]. Nilsson et al. investigate the limitations of graphene as an effective corrosion-inhibiting coating on Pt(100) surface for O2 and CO using scanning tunnelling microscopy measurements and DFT calculations [291]. Graphene layer is found to protect Pt(100) surface against O2 and CO exposure up to a certain pressures of 10~4 mbar

0 1 2 3 4 5 0 25 50 75 100 125

y(pm) Time, t(min)

Fig. 29. (a) Schematic of a microscopic graphene membrane on a SiO2 substrate. The microchamber filled with 2 bar (200 kPa) of H2 (represented as red circles). Equilibrium was reached (pint = pext) by diffusion through the SiO2. (b) After pint > pext the membrane bulged upwards. Etching pore(s) in the graphene membrane bigger than H2 but smaller than the air molecules (mostly N2 and O2, denoted as green circles) (c) leaks out the H2 from microchamber, and (d) all the H2 molecules have leaked out of the microchamber, the membrane deflected downwards. (e) Deflection versus position, measured from 0 min (black) to 8 min (dashed blue) after etching. (f) Maximum deflection d versus t for one membrane separating H2 from air, measured by AFM. Black symbols represent the H2 leak rate before etching and red symbols the H2 leak rate after introducing selective pores into the graphene. Inset: optical image of the bilayer graphene flake used in this study, which covers many cavities in the SiO2 substrate (scale bar is 60 mm) [283].

and 10~6 mbar, respectively. At higher pressures, CO is observed to intercalate under the graphene coating layer.

Designing the polymer thin films or membranes to prevent gas or water molecules from permeating through is a major challenge in many applications for food and electronics. The graphene-polymer composites have been investigated as improved gas barriers due to the large aspect ratio of graphenes

Fig. 30. Photographs of bare Cu foil, partially and fully grown graphene on Cu-foils during oxidation at 443 K [289].

[292-296]. For example, a poly(ethylene-2,6-naphthalate)-graphene nanocomposite film with the addition of just 4 wt% of graphene shows 60% decrease of the H2 permeability [292]. Up to 90% decrease in N2 permeation is observed in thermoplastic polyurethane films reinforced with isocyanate treated GO of 3 wt% [293]. Whereas, the solution filtration synthesised PEI (polyethylenimine) func-tionalized reduced GO thin films with a brick and mortar structure show a significant decrease in the H2 permeation rate with increasing PEI content due to tight packing and filling up of gallery spac-ings of the film assembly [294]. A considerably reduced O2 and CO2 permeation is observed in a layer-by-layer assembly of GO and PEI composite (Fig. 31) [295]. A 10 bilayer film (~91 nm thick) with 0.1 wt% PEI and 0.2 wt% GO mixtures deposited on a 179 im PET shows O2 transmission rate of 0.12 cm3/m2/day. A high H2/CO2 gas separation with a selectivity higher than 383 is observed, could be due to strong binding of CO2 with the amine groups. A solvent blending prepared graphene-poly(ethylene vinyl alcohol) (RGO/EVOH) thin film (a thickness of about 0.2 mm) with only 0.5 wt% of RGO loading shows a significantly decreased permeable coefficient of O2 nearly 1671 times lower than that of neat EVOH film [296]. The permeability measurements are carried out by a differential volume-variable pressure method at room temperature with 37% relative humidity and O2 flow at 1 bar on 80 mm diameter film. The interesting gas barrier properties are attributed to a tightly packed nanobrick wall structure that creates super tortuosity and diffusion length for gas molecules.

Ultrathin (up to 20 nm) graphene and GO based membranes also show excellent gas permeation selectivities [297,298]. CVD grown large-area monolayer graphene films with increasing numbers of graphene sheets deposited on a Poly(1-methylsilyl-1-propyne) (PTMSP) substrate show improved O2/N2 selectivity suggesting that the gas diffusion through irregularly aligned defective graphene pores as well as slit-like interlayer spacings [297]. The GO thin-film membranes deposited on polye-

Fig. 31. Top: Schematic of the layer-by-layer deposition process. Steps 1-4 are repeated until the desired number of bilayers are achieved. Bottom: Images of (PEI/GO) 10 and (PEI/GO) 20 thin films on 179-im PET to demonstrate transparency [295].

thersulphone supports in humidified state show a high selectivity of CO2/N2, which is ideal for postcombustion CO2 capture. The gas permeance (in gas permeation units, GPUs) measured at a feed pressure of 1 bar, using the constant-pressure and variable-volume method shows very different gas transport behaviour (Fig. 32). Introduction of humidity in the feed stream enhances the CO2 permeance leading to enhanced selectivities of CO2/CH4, CO2/H2 and CO2/N2. The highly interlocked layer structure of GO membranes shows a gas permeation in the order of CO2 > H2 p He > CH4 > O2 > N2. As shown in Fig. 32d the separation performance of GO membrane in its dry and humidified state is higher than that reported for CO2 for polymeric membranes, including thermally rearranged (TR) polymers, polymers of intrinsic microporosity (PIM), and inorganic membranes, such as carbons, silicas, and zeolites. A more permeable, H2/CO2 selective is demonstrated by creating thermally generated irreversible and small defective pores on the basal plane due to partly released surface oxygen-containing functional groups by thermal reduction of GO membranes. Overall the enhanced CO2 selectivities are attributed to the CO2-philic permeation behaviour, which is further enhanced by the presence of water. It is predicted that the incorporation of a carboxylic acid group lead to the highest isosteric heat. Although CO2 is a nonpolar gas, the polarity of the individual C-O bonds

Fig. 32. Gas transport behaviour through ultrathin GO membranes. (A) Gas permeances of GO membranes as a function of molecular weight (method one; dashed line represents the ideal Knudsen selectivity) under dry and humidified conditions (amu, atomic mass unit). (B) H2 and CO2 permeances and H2/CO2 selectivity of method one GO membranes as a function of permeation time. (C) Gas permeances of GO membranes as a function of kinetic diameter (method two) under dry and humidified conditions. (D) Relation between CO2 permeability and CO2/N2 selectivity of method two GO membranes under dry and humidified conditions [TR; tetrazole PIM (TZPIM) and PIM; CMS]. Error bars indicate the SD of all raw data [297].

in the molecule allows for interaction with polar groups in GO. Thus, CO2 can act as a Lewis acid or a Lewis base and can participate in hydrogen bonding. Carboxylic acid groups on GO provide a preferential site for CO2 adsorption, consequently retarding CO2 transport in the nanopores by strongly trapping CO2 molecules. This phenomenon is counterintuitive, because strong affinity between penetrant and nanopore walls often leads to surface diffusion of the penetrant preferentially sorbed on the pore wall, resulting in flux enhancement of highly sorbed or condensable penetrants. As such, these membranes are promising materials for industrial CO2 separation processes related to petrochemical engineering (CO2 removal from natural gas), the environment (CO2 capture from flue gas), and biomass energy (CO2 recovery from landfill gas).

Ultrathin (~1.8-18 nm thickness) GO membranes prepared by vacuum filtration on anodic aluminium oxide (AAO) support from controlled dilution of GO dispersions show unusual H2/CO2 and H2/N2 separation selectivities as high as 3400 and 900 respectively, at 293 K [298]. Normally the micropo-

rous membranes show low H2/CO2 selectivity (<10) or are selective to CO2 over H2 at temperatures below 373 K due to a strong CO2 adsorption and blocking of H2 permeation. GO membranes with increasing thickness from 1.8 nm to 180 nm show exponential decrease in H2 and He permeances. Here the gas permeance is mainly attributed to the selective structural defects within GO flakes. The GO interlayer spacing has minimal effect on molecular transport. The H2/CO2 separation selectivity decreased with increasing temperature, resulting from the faster increase of CO2 permeance. A ~18 nm thick GO membrane on cellulous acetate support show H2/CO2 and H2/N2 separation selectiv-ities of 1110 and 300, respectively. As shown in Fig. 33, the performance of the ultrathin GO membranes is far above the upper bound for polymeric (black line) and inorganic membranes.

4.9. NH3, NO2, H2S and SO2 sorption

The emissions of NH3, NO2, SO2, and H2S from the industrial plants during number of combustion processes, wastewater treatment, and food and compositing process are toxic, corrosive and malodor-

Fig. 33. Top: Comparison of ultrathin GO membranes with polymeric membranes and inorganic microporous membranes for H2/CO2 mixture separation (50:50 H2/CO2): selectivity versus H2 permeance. The black line denotes the 2008 upper bound of the polymeric membrane for H2/CO2, assuming membrane thickness is 0.1 mm. Blue points (1-9) represent microporous inorganic membranes from the literatures; red points (10) indicate ultrathin GO membranes. The table at bottom explains points 1 through 10. ZIF, zeolitic imidazolate framework; MOF, metal-organic framework [298].

ous air pollutants. Due to growing demand for a more economical and improved process, there is always a need of search for new filtrate/adsorptive removal materials. Recently, the porous inor-ganic-graphene based materials for adsorptive removal of NH3, NO2, SO2 and H2S pollutants have been in focus of attention. Many theoretical and experimental studies are conducted for molecular adsorp-tive detection/capturing on graphene surfaces and GO based functionalised/inorganic hybrid porous structures. The DFT calculations suggest that adsorption of nitrogen oxides, NOx: NO, NO2, NO3 on GO is generally stronger than that on pristine graphene due to the presence of active defect sites, such as the hydroxyl and carbonyl functional groups [299]. The interaction of NOx with GO is expected to result in the formation of hydrogen bonds OH - O (N), weak covalent bonds C---N and C-- O, as well as the nitrous and nitric acid like moieties. First-principles calculations reveal that the oxygen groups on GO act as strong binding sites and induce dissociation of the NH3 into the NH2 or NH species by the H atom abstractions [300]. This further leads to the removal of surface oxygen species through the hydroxyl group hydrogenation and the ring opening of epoxy group. The reaction of NH3 with the hydroxyl and epoxy groups is generally exothermic. From the first-principles calculations adsorption of H2S onto pristine graphene is found to be very weak due to its small binding energy, large bonding distance and small net charge transfer [301]. Whereas, the Pt-decorated graphene bind up to seven H2S molecules to one side of a Pt-graphene system due to large binding energy and a short Pt H2S bonding length. The adsorption of several acidic gases (CO2, NO2 and SO2) on light metal (Li, Al) decorated graphene oxide (GO) is also investigated with the first-principles calculations [302]. It is found that Li and Al could be anchored stably by hydroxyl and epoxy groups on GO and acts as strong adsorption sites for CO2, NO2 and SO2 with improved binding energies up to (0.59, 2.29 and 1.08) eV, respectively, compared with Ti.

Experimentally, it is also identified that the acidic nature of GO with its carboxyl, hydroxyl, and epoxy groups act as energetic sites for chemical adsorption of gases, especially for polar molecules. A highly polar NH3 interaction with the GO through gravimetric uptake measurements is demonstrated back in 1962 [303]. About 180 mg/g of NH3 adsorption at room temperature and atmospheric pressure is observed. A physical adsorption at defective sites/or by hydrogen bonding with hydroxyl groups and a chemical interaction with the acidic sites and condensation in pores is proposed at low-and high-pressures, respectively. Therefore an increase in graphene inter layer distance proportional to the size of the ammonia molecule is observed.

Recently, there are renewed interests and increased investigations on NH3 adsorption in various types of GO based materials, due to the development of promising and efficient new synthesis methods for obtaining GO and its composites. Bandosz's group have been actively exploring the NH3, NO2, H2S and SO2 adsorption and interaction properties in various types of GO and its composites with polymers, metal oxides, and MOFs [304-327]. The NO2, NH3 and H2S molecules have a small kinetic diameter between (0.30-0.36) nm but exhibit different chemical properties (e.g., acid-base and redox properties). The amount of gas adsorption is demonstrated through the breakthrough measurements by passing a gas stream through a sample packed column with the option of either dry or up to 70% relative humidity at room temperature. The simple schematic of experimental setup is shown in Fig. 34 [304]. The effects of the degree of oxidation, reduction and exfoliation of the GO samples are examined for toxic gaseous adsorption. For example, a high NH3 breakthrough capacity of 61 mg/g is observed in GO, dried at room temperature compared to 28 mg/g in GO dried at 393 K and 17 mg/g in GO exfoliated at 573 K [305-307]. The acid-base reactions and intercalations between the GO layers are identified as the main mechanisms of NH3 retention. The NH3 adsorption is very sensitive to the additional functional groups present in GO when different synthesis methods are employed. For example, GO prepared by the Brodie method, NH3 is mainly retained via intercalation in the interlayer space of GO and reaction with the carboxylic groups at the edges of the graphene layers. On the contrary, the NH3 reaction with the epoxy, carboxylic and sulphonic groups is observed in the Hummers GO [305]. In addition, the role played by surface texture, degree of oxidation and interlayer space/or pore structure on the NH3 adsorptive retention is explained [306]. NH3 dissolving in lamellar water as ammonium ions or by forming hydrogen bonding with epoxy and phenolic surface groups is also proposed [307,308]. The interaction of NH3 with carboxylic groups is proposed as either Bransted:

Fig. 34. Breakthrough capacity experimental set up [304].

or Lewis acids:

The second reaction seams more likely to happen in the dry conditions.

A variety of GO-composites are synthesised with the introduction of surface acidic groups, metal salts or solid acids. The intercalation of aluminium and aluminium-zirconium polyoxycations (also called Keggin type cations; in case of aluminium oxycation, the cations have one tetrahedral aluminium surrounded by twelve aluminium in octahedral coordinations, [Al13O4(OH)24(H2O)12]7+) in GO layers show a highly increased interlayer spacing (up to 1.5 nm) and NH3 sorption due to additional Lewis and Bronsted acidic centres from the Keggin cations [310]. The GO-POM composites also show enhanced NH3 sorption and retention than GO alone due to the interaction with the both bridging and terminal oxygens of POM (hydrogen bonded to the terminal oxygens of the POM could enhance the

acidity) [311]. POM is a kind of Keggin polyanion (consists of a central tetrahedron PO4 surrounded by twelve octahedral MO6 (M = W or Mo)) and exhibits a high acidity. The enhancement of NH3 adsorption in clays-graphene composite is attributed to the synergistic effect of structure as well as graph-ene edge acidic groups and clay origin Bronsted or Lewis acidic centres [312]. The GO with simple metal-oxide (MnO2) nanocomposites are also tested for NH3 sorption [313].

Interestingly, GO in its dry condition show negligible adsorption for other pollutants; NO2, H2S and SO2. However, an enhanced H2S sorption (breakthrough capacity of 30 mg/g compared to zero in dry GO) is observed in a functionalised GO that is treated with NH3 at 1223 K [314]. The NH3 treatment results in decomposition and rearrangement of oxygen functional groups and incorporation of nitrogen mainly in pyridinic and quaternary entities. These nitrogen functionalities enhance the basicity and attract HS~ ions of dissociated H2S via electrostatic interactions. The oxidation to elemental S and a smaller amount of SO2 is also detected. The GO-metal oxide/hydroxide composites show enhanced sorption of NO2, SO2 and H2S by forming nitrates, sulphide, sulphite and sulphates through the reactive paths (1) and (2) shown below [315-320]:

An iron-oxide (a-Fe2O3 and Fe3O4)-pillared graphene composite with a hydrophobic nanospace shows high activities for O2, NO and NO2 due to the interaction of NO with active y-FeOOH and a-Fe2O3, and the formation of bridging nitrates [321,328]. Instant reduction of NO2 to NO is noticed. A favourable adsorption and conversion of NO2, and NO retention is also observed on GO-polyaniline (PANI) composites [322]. GO in its dry state show no NO2 adsorption, however the sorption and retention of NO2 is observed with the presence of N-methylformamide between GO layers due to polar and acid-base interactions.

The MOF-GO composites show an improved NH3, NO2, and H2S sorption compared to the constituents because of synergistic effect related to the new pore space in the interface between MOF units and GO and a reactive interaction of metal sites and ligands either by hydrogen bonding (MOF-5) or coordination and subsequent complexation (HKUST-1) [323-327,329]. For example, the Lewis interactions of NH3 and the Cu sites, (70-80) kJ/mol and the reaction of NH3 with the ligands, (70-100)kJ/mol exhibit stronger binding energy. The NH3 adsorption in Fe-MIL shows retention of the framework structure and thus could be useful for adsorbent regeneration. As shown in Figs. 35 and 36 initially dark blue HKUST-1 and its GO composites turned to light blue upon adsorption of NH3, NO2 and H2S [323,329]. With the progress of the adsorption test, the colour of the sample further changed to a light blue or blackish tint during exposure to NH3 and H2S, respectively. The first colour change is attributed to the coordination to the Cu sites due to the presence of a lone pair of electrons on the adsorbate molecules. The second colour change is likely caused by the formation of new complexes like CuS. The reactive decomposition in HKSUT-1 is attributed to the formation of the Cu(NH3)4+, CuS and Cu(NO3)2 for adsorption of NH3, H2S and NO2, respectively. In Zn-MOF-GO composites, NH3 interact through hydrogen bonding with the oxygen atoms of the ZnO4 units whereas for Fe-MIL-GO composites, the acid-base reaction thought to be the reactive pathway between the NH3 and water coordinated Fe sites.

Long et al. show an oxidative absorption of SO2 in the porous GO foams, prepared by freeze-drying technique [330]. As shown in Fig. 37 the GO foams exhibit a high porosity and resemble a sponge-like appearance. The oxidative adsorption of SO2 to SO3 by reduction of GO at room temperature is observed in a dry condition. Whereas, the reaction is intense with the GO water suspension; the GO is reduced quite fast and SO2 becomes SO4~.

Fig. 35. Schematic of the adsorption process of NH3, H2S and NO2 on the HKUST-1 with evidence of the colour changes and the identification of the reactions products [323].

Fig. 36. Mechanism of the de-sulphurization process of MOF-5/GO composites [329].

5. Summary and perspectives

As discussed above and shown in Figs. 1 and 2, there exist a wide range of graphene/GO based materials; starting with expanded graphite, intercalation, exfoliation, chemical reduction, pillared layers, self-assembly, thin-film membranes, functionalization, doping, metal dispersion, porous template and chemical activation. Those are investigated for gas sorption, storage and separation. In particular, the experimental H2 and CO2 uptake results are comparatively analysed in various graphene based materials, as indicated in Tables 2-5, along with the SSA and the measurement conditions (particu-

Oh 2 4 8 12 24 36 60 96 144 2l6h

Fig. 37. Top: SEM image of GO foam, and GO solution before (brown) and after (dark) coming in contact with SO2 gas, clearly showed a reduction of GO. Bottom: Colour change of GO foams exposed to SO2 gas in dry form. The SO2 gas is introduced into the column containing the GO foams, the colour of the foams gradually changed from brown to black. This suggests that the GO reacted with the SO2 gas and the extent of the reaction gradually increased with time [330].

Table 2

Experimentally reported H2 adsorption capacities in exfoliated and reduced GO based materials.

Sl. no. Type of graphene material Specific surface area (m2/g) H2 uptake (wt%) 77 K 298-300 K Ref.

1 Exfoliated GO 156 0.4 (1 bar) <0.2 (60 bar) [199]

2 Exfoliated GO 300 1.75 (54 bar) 0.1 (40 bar) [200]

3 Exfoliated GO 470 - 0.7 (40 bar) [249]

4 Exfoliated GO 477 1.2 (1 bar) - [202]

5 Exfoliated GO 477 1.4 (l bar) - [254]

6 Diamond derived graphene 520 0.68 (1 bar) 2.5 (100 bar) [207]

7 Expanded graphite nanofibers 67 & 555 0.19 & 0.39 (1 & 0.29 (20 bar) [220]

20 bar)

0.45 & 1.2 (1& 20 bar) 0.03 (20 bar)

8 Hydrazine reduced GO 640 0.68 & 1.2 (1& 10 bar) 0.72 (100 bar) [201]

9 Exfoliated GO 925 1.38 (1 bar) 3.1 (100 bar) [207]

10 Glucose reduced GO 1206 - 2.7 (25 bar) [203]

11 Outgassed GO and annealed GO 480 & 1305 1.4 & 4.0 (1 bar) - [208]

12 Magnesium combustion derived graphene 236 0.85 (65 bar) 0.9 (300 bar) [205]

13 Camphor, Diamond and GO derived 46-1550 0.05-1.7 (1 bar) - [206]


14 Exfoliated GO by CO2 pressure swing 141-547 - 0.7-1.8 [204]

(100 bar)

larly, pressure and temperature). As specific surface area and pore volume are the key porous parameters to determine the physisorption capacity Fig. 38 shows the H2 uptake in a wide variety of graph-enes against the SSA and pore volume. The structure modification effect on the binding, heat of H2

Table 3

Experimentally reported H2 adsorption capacities in various pillared graphene based materials.

Sl. no. Type of graphene material Specific surface area (m2/g) H2 uptake (wt%) 77 K 298-300 K Ref.

1 K-GIC, KC24 - 1.2 (1 bar) - [218]

2 GO-MWCNTs - - 2.6 (50 bar) [234]

LC-GO 1.4 (50 bar)

3 Self-pillared GO 57-192 1.1-4.8 (90 bar) 0.28-0.46 (90 bar) [222]

4 Diazonium linked graphene 260-440 1-1.6 (2 bar) - [229]

5 GOFs 26-470 0.35-0.92 (1 bar) - [107,227]

6 GPOM 580 & 680 1.3 & 0.8 (1 bar) - [226]

7 GMNPs 418-901 0.88-1.44 (1 bar) - [225]

8 Silylated GO 562-942 - 0.2-0.4 (100 bar) [233]

Table 4

Experimentally reported H2 adsorption capacities in doped, functionalised, activated, templated, bottom up graphene based materials.

Sl. no. Type of graphene material Specific surface area (m2/g) H2 uptake (wt%) 77 K 298-300K Ref.

1 20 wt% Pd-graphene - - 3.0 (40 bar) [249]

2 N-doped 146 - 1.4 (40 bar) [250-

Pd-N-graphene 1.5 & 4.5 (40 bar) 252]

3 Expanded nanofibers 227 5 (100 bar) - [258]

4 Exfoliated GO-24 wt% Pd 80-230 0.3-0.5 0.1-0.26 (1 bar) [244]

(1 bar)

5 Ni-B nanoalloy decorated graphene - 2.8 (1 bar) - [253]

272 4.4 (1 bar) 1.9 (1 bar, 273 K) [254]

6 Exfoliated GO & its Pt and Pd composites 755 - 0.07 (57 bar & 303 K) [247]

480-545 0.16 (57 bar & 303 k)

7 Pd-graphene/carbon composites 300/3328 - 0.82 (80 bar & 298 K) [248]

6 12 wt% Pd4Hg alloy decorated GLF 450 - 2.3-5.25 (20 bar) [256]

7 10 wt% Pd decorated GLF 687 - 0.98 (100 bar) [257]

8 Bottom up & functionalised graphenes 200-900 0.5-1.7 - [262]

(2 bar)

9 Template CVD graphene 2038 - 1.5 (90 bar & 274 K) [267]

10 Bottom up graphene 2139 - 0.9 (100 bar) [261]

adsorption is shown Fig. 39. Earlier, the H2 uptake capacity against the SSA of a wide range of other porous carbons is generalised and is normally accepted that for every 500 m2/g of SSA, the H2 uptake capacity at 77 K is increased by ~0.5 wt% at 1 bar and wt% at high pressures [82,331,332]. Similarly, at room temperature and pressures up to 100 bar the H2 storage in a wide variety of porous carbons is found to be well below 1 wt% [333]. From Fig. 38 one can clearly understand that the modified graphenes with carefully tuned pore structure or pillaring or chemical modifications (doping, func-tionalization and metal dispersion) performed much better than simply exfoliated graphenes. At 77 K, most of the exfoliated/reduced graphenes seems to exhibit more or less similar uptake capacities compared to the other porous carbons with respect to the porosity. Normally the exfoliation leaves very high mesoporosity due to the highly curved nature of graphenes. However, it is understood that the modified graphenes enhance the microporosity and functional active sites. Thus comparatively a high H2 uptake is seen in optimised GO structures, with doped, defective, and metal/alloy nanoparticle dispersed graphenes due to the synergistic effect of enhanced microporosity, binding sites or combined molecular and spillover/dissociative atomic hydrogen adsorption. Most importantly these graphene based materials exhibit very promising H2 uptake capacities between (1-4) wt% at 300 K compared to the maximum of up to 1 wt% in most porous solid adsorbents, i.e. MOFs, zeolites and nanostructured carbons. This is further evidenced by a high heat of H2 adsorption mostly between

Table 5

Experimentally reported CO2 adsorption capacities in a variety of graphene based materials. The respective available surface area and references also included for readers further interests.

Sl. no. Type of graphene material Specific surface area (m2/g) CO2 uptake (mmol/g) Ref.

1 GO 31 1.7 (1 bar & 273 K) [242]

GO-PEI hydrogel 253 2.5 (1 bar & 273 k)

Reduced GO 876 1.84(1 bar & 273 K)

2 Outgassed GO 477 0.7 (1 bar & 273 K) [202]

2 Exfoliated GO 925 8.6(1 bar & 195 K) [207]

3 Hydrogen exfoliated GO 443 21.6 (11 bar & 298 K) [210]

4 Exfoliated nanoplatelets 480 56 (30 bar & 298 K) [211]

5 Exfoliated GO 701 39 (30 bar & 298 K) [211]

6 Graphene-polyanilene - 70 (10 bar & 298 k) [213]

7 Amine grafted GO Nonporous 1.2 (15% CO2, 40 ml/min, 303 K) [212]

8 GO-Mn3O4 541 2.5 (1 bar & 298 K) [224]

9 GOFs 470 2.7 (4 bar & 300 K) [107]

10 GLF 510 2.0 (10 bar & 298 K) [255]

11 Exf-GO by CO2 pressure swing 140-547 2.3-6.4 (30 bar & 298 K) [204]

12 KOH activated GO 1894 2.0 & 19 (1 bar & 26 bar & 298 K) [263]

13 KOH activated GO & N-doping 1360 4.0 (1 bar & 300 K) [264]

14 KOH activated polythiophene-reduced GO 1396 4.5 (1 bar & 298 K) [266]

15 Template CVD graphenes 2038 36.5 (31 bar & 274 K) [267]

• 1 77 К & 1 bar

▼ ■

m AA A A

►* " A

A . i 1 1 1

500 1000

BET SSA (m2/g)

j 3 01 ra

a. 2 l"

1 1 1 * ► 1 1 1 1 77 К & >1 bar

► -

> / 1 » 1 : 1 A 1 ■ к ■ 1 . . 1

300 600

BET SSA (m2/g)

500 1000

BET SSA (m2/g)

6.5 7.0

Total pore volume (cm /g)

Fig. 38. Experimentally obtained H2 uptake capacity versus surface area in various types of graphene based materials; (a) 77 K and 1 bar, (b) 77 K and high pressure, >1 bar, (c) ~300 K and all pressures. (d) H2 uptake capacity at 77 K and 1 bar against pore volume of various graphenes. The materials can be identified with the corresponding values given in Tables 2-4.

0 5 10 30 35

Heat of H2 adsorption (kj/mol)

Fig. 39. Experimentally derived heat of H2 adsorption in various types of graphene based materials.

(8-15) kJ/mol in the modified graphenes, compared to a typical (4-6) kJ/mol in other porous carbon adsorbents. These facts are clearly encouraging and help to minimise the gap between DOE targets and experimental results. Similarly, the highest CO2 capacity is reported in functionalised/doped and highly porous activated graphene materials (Table 5). The graphene based materials are also found to exhibit enhanced CH4 storage that surpasses the DOE target and other traditional highly porous carbons and most of the MOFs. The graphene-based composites also show enhanced NH3, NO2, SO2 and H2S sorption and retention. The theoretical and experimental results on the graphene related membranes for gas separations are very encouraging.

Because of multifunctionality of the graphene-based materials many approaches have been adapted to develop various robust interconnected porous network architectures with various functional groups. Clearly, by controlling the synthesis parameters and tuning the pore parameters; increasing SSA and microporosity, the chemical modification of graphenes could be a very promising in obtaining practicable gas sorption, storage and separation. There are many possible routes still available to engineer graphene structures. For example, a number of robust 3D hierarchical porous graphene structures have been developed through self-assembly, sol-gel and cross-linking chemistry [43,334-338]. A vacuum centrifugal evaporation has been introduced as a strategy for a large scale synthesis of GO sponges with 3D interconnected network hierarchical structure [334]. Subsequent thermal annealing of the GO sponges result in recovery of sp2 graphene structure, forming graphene sponges with a large SSA and typical porous structure. An autoclaved leavening and steaming of GO layer films lead to reduced GO foams with open porous and continuously cross-linked structures [335]. Compared to reduced GO films, the reduced GO foams show greatly improved performance as selective organic absorbents. Furthermore, the doping of the graphene foams is performed to enhance its functionalizability. These graphene foams and its thin film forms could be ideal candidates for an efficient gaseous separation.

Optimised pillared graphene structures with amine-based functional groups could produce superior CO2 adsorbents (pillared GOs with NH2 groups like silicas with NH2). Much stable structures may be developed due to strong nucleophilicity of amines and epoxy groups on GO through a ring opening reaction without any activation. Using sol-gel method low-density graphene aerogels with tuneable SSA, (584-1200) m2/g and very high pore volume, (2.96-6.4) cm3/g are synthesised [336]. In another method by gelation of a GO suspension under basic conditions, a robust 3D macroassembly of graphene sheets is obtained with SSA> 1300 m2/g [337]. Furthermore, the chemical activation of graphene aerogels with KOH results in highly increased microporosity (up to 0.69 cm3/g) and SSA (up to 1715 m2/g) [338]. Similarly, CO2 activation of the mechanically robust, centimetre sized self-assembled interconnected networks of single-layer graphene aerogel monolith increases the SSA up to 3000 m2/g compared to the typically 400 m2/g for the non-activated foam [339]. Thus, adjusting synthetic parameters allow a wide range of control over surface area, pore volume and pore size, as

well as the nature of the chemical cross-links (sp2 versus sp3) making these materials promising candidates for gas sorption and storage applications.

A method for preparing functionalized 3D carbon-based architectures, consisting of high SSA mes-oporous carbon spheres (MCS) with a particle size of 60 nm intercalated between graphene sheets (GS) is demonstrated by Lei et al. [340]. The MCS itself consists of graphene sheets in small domains and exhibits a SSA and pore volume of 2396 m2/g and 2.9 cm3/g, respectively, obtained from silica template and ferrocene by CVD method. MCS-GS composite shows hierarchical porous architecture with SSA of 1496 m2/g and pore volume of 3.36 cm3/g with enhanced microporosity. Further thermal treatment of GMC-GS in NH3 yields N-doped product. The N-containing basic groups could be very useful in obtaining improved CO2 adsorption at low-pressures due to enhanced acid-base interactions. In addition, a wide variety of high and tuneable surface area and pore structure assemblies of graph-enes with chemical functionalization are demonstrated through template, pyrolysis and chemical activation methods [341,342]. For example, the silica and zeolite based templates are conveniently used to produce a very high SSA and pore volume graphenes and carbons. In particular, by controlling the pore size of silica one could obtain much more enhanced microporosity than the reported methods of producing mesoporous graphene structures. Other possible methods, such as soft co-polymer or sacrificial salt or nanoparticle templating, could provide desirable structures as well. The chemical activation is another promising method in which the tailor made porous graphene structures are highly possible by simply choosing the right activator; KOH, NaOH, Na2CO3, CO2, steam, etc. together with concentration and temperature [342]. The chemical modification, functionalization and hybrid structures of graphene based highly porous materials [343-368] could have attractive gas sorption and storage properties. The boron- and nitrogen co-doped and porous graphene like materials should help in obtaining higher gas sorption and storage. Zhang et al. show that a very efficient and industrially scalable approach of synthesising defect/wrinkle mesopore structure graphene sheets with a very high SSA up to 3523 m2/g and total pore volume up to 2.4 cm3/g [366]. The material obtained through the following two standard industry steps: (1) in-situ hydrothermal polymerisation/carbonisation of the mixture of cheap biomass or industry carbon sources with GO to get the 3D hybrid precursors; and then (2) a KOH chemical activation step. An enhanced surface area and pore volume is also obtained with the addition of GO to mesoporous carbon. Tuneable surface area and porosity parameters are also demonstrated in mesoporous carbons with the controllable addition of GO.

Recently, graphenic type fragments are synthesised from zeolites/MOFs as templates, with either self- or addition of small molecular carbon precursors [369-375]. Upon carbonisation, the organic linker ligand in MOF transforms into single layers functionalised graphene fragments with new hierarchical high slit-like micropores and mesopores while retaining the MOF framework porosity. These carbons with new pore types show ultrahigh porosities i.e., simultaneous high surface area (up to 3300 m2/g) and micro- and total-pore volumes (up to 5.53 cm3/g) [369,372]. These functionalised carbons yield exceptionally high H2, CH4 and CO2 uptakes compared with the capacities of those benchmark microporous materials under identical conditions. In particular, a gravimetric (volumetric) H2 uptake of 3.25 wt% and 7.3 wt% (50 g/l) are achieved at 77 K under 1 bar and 20 bar, respectively [369,370]. At 100 bar and 298 K, these carbons show ~0.94wt% of H2 storage. The functionalised MOF carbons also show record high CH4 and CO2 uptakes [372-375]. For example, the room temperature uptake of up to 4.6 mmol/g at 1 bar and over 27 mmol/g (119 wt%) at 30 bar are again some of the highest values reported in the literature for porous carbons. The overall synthesis and gas uptake mechanism of such graphenic type MOF carbons is schematically shown in Fig. 40.

The novel ultrathin, robust nanocomposite membranes by incorporating GO sheets into silk fibroin matrix through heterogeneous surface interactions in organised layer-by-layer manner show outstanding mechanical properties, with a tensile modulus of 1.45 x 106 bar (145 GPa), an ultimate stress of more than 3000 bar (300 MPa), and a toughness of above 2.2 MJ/m3 [376]. These nanocomposite membranes in their support and free-standing state could be valuable for potential applications in protective molecular coatings, permeable membranes for separation and delivery. A new avenue for graphene-based ultra-films anchored on solid supports with a high stability and controlled thickness via a layer-by-layer assembly is demonstrated using amino-substituted p-conjugated compounds, including 1,4-diaminobenzene, benzidine, etc., as cross-linkages [377]. A thermal annealing leads to a reduction of the films thus could yield suitable pore space for molecular sieving. Furthermore, the

(25-400) "C (600-800) °C >900 °C

(1) (2) (3) (4)

300 600 900 1200 Temperature (°C)

Fig. 40. Top: Schematic illustration for steps in fabricating hierarchical porous graphene carbons (HPCs) from a precursor MOF-5; (1) MOF framework/crystal structure (up to 400 "C) with well-defined pore structure in the micropore region, (2) decomposition, and metal-oxide formation, ZnO and carbon between 600 and 800 "C yielding poor pore development, (3) at >900 "C, ZnO reduction and evaporation of Zn and CO yields highly porous carbon with a hierarchical pore structure and (4) surface morphology of HPCs by TEM and SEM. Middle: Mechanism involved in the carbonisation process of MOF-5 through simultaneous thermogravimetric analysis (TGA), mass spectroscopy (MS) and powder X-ray diffraction (PXRD). The first mass-loss between 400 "C and 600 "C, is attributed to the framework decomposition leading to a major release of carbon containing gaseous products (mostly CO2, CO, C6H6 and a small amount of H2 and CxHy hydrocarbon mixtures), where a formation of crystalline hexagonal ZnO occurs (sharp XRD peaks). The second mass-loss starting at ~900 "C is due to further release of CO2 and CO with Zn through the reduction of ZnO by carbon via ZnO + C/Zn(g) + CO. The carbonisation yields materials with either a featureless XRD pattern or very weak and broad XRD peaks between 20" and 25", and at ~45" of two-theta indicating the disorderly oriented tiny graphenic type fragments in the structures. It is worth noting that the reduction of ZnO by carbon during carbonisation yields a pure porous carbon. Bottom: Porosity and CO2 uptake characteristics of HPCs. The 77KN2 adsorption isotherms on the millimetre-sized MOF-5 derived HPCs show simultaneously high surface area, up to 2734 m2/g, and exceptionally high total pore volume, up to 5.53 cm3/g. In the HPCs, micropores are mostly retained and meso- and macropores are generated from defects in the individual crystals, which is made possible by structural inheritance from the MOF precursor. The resulting HPCs show a significant amount of CO2 adsorption, over 27 mmol/g (119 wt%) at 30 bar and 27 "C. Inset also show a direct relationship between the CO2 adsorption capacity and the surface area: a CO2 uptake of 10 mmol/g for every 1000 m2/g increase of surface area under 30 bar and 27 "C [372].

amine groups could show selective gas permeations. A simple and effective method of preparing functional nanoporous graphene/carbon composite membranes from MOFs is also demonstrated and reviewed [378,379]. The resulting symmetric dense or asymmetric composite membranes exhibit good performance in gas separation and liquid separation via pervaporation.

6. Conclusion

In summary, the rich chemistry of graphene nanosheets have offered great potential for the design and development of many combinations of interconnected porous graphene-based materials as promising molecular adsorbents, with high thermal/chemical stability, tuneable specific surface area and pore structure. The other advantages are the relatively low cost and easy scale-up associated with the manufacturing of the porous graphene carbons. Graphene and its composites have been used as adsorbents for low temperature physisorption and high temperature chemisorption of H2, CO2, CH4, etc., showing promising results, especially when compared with conventional high surface area activated carbons. As highlighted in this report, the interest in gas sorption and storage by physical adsorption has grown considerably over the last few years. While many computational studies prove strong potential of graphene and its derivatives, practical developments have advanced markedly in molecular pillaring, metal dispersion, doping, functionalization and 3D interconnected networks to enhance gas sorption and storage. Clearly, there still exist several challenges in practice, such as the effective control of the desirable binding sites, porosities and functionalities; the balance of functional properties with structural integrity; and further tuning of molecular selectivities in gas mixtures for purification and sensing.

More specifically, there have been many attempts to use carbon adsorbents as reversible H2 storage materials. Most of the measurements so far have been conducted at 77 K, which limits the practicality. Similarly, carbon capture at ambient conditions is not satisfactory. As summarised in Tables 2-5, and Figs. 38 and 39, the remarkable synergistic properties from combining GO hybrid structures showed promising room temperature H2 storage and high pressure CH4 and CO2 adsorption capacities. For the given rich chemistry and functionalizability of graphene, the number of new compounds are continuously being tailor-made. Thus if carefully tuned the microporosity and the high surface area of the chemically modified graphene materials could be very promising for H2 and CH4 storage and CO2 capture. Similarly, the gas permeation separation on graphene and GO based membranes is highly dependent on GO sheet size and thickness, layer structure and assembly, residual lamellar water content, and deformation of thin-layered structures during dewetting can all effect. There is much scope in tuning and controlling the parameters to develop practically useful membranes.


This work was supported by EPSRC Grant Nos. (EP/L018330/1 and EP/G063176/1).


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